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0 DEVELOPMENT OF MICROSTRUCTURE DURING SINTERING AND ALUMINIUM EXPOSURE OF TITANIUM DIBORIDE CERAMICS
Thesis submitted in partial fulfilment of the requirements for the degree Doktor Ingenipr
Norwegian University of Science and Technology Department of Physics
April 1997 DISCLAIMER
Portions of this document may be illegible in electronic image products. Images are produced from the best available original document. ACKNOWLEDGEMENTS
The work presented in this thesis was carried out at the Department of Physics, Norwegian University of Science and Technology (NTNU) in the period March 1993 - April 1997. Financial support was provided by the Norwegian Research Council who paid my salary and Norsk Hydro who helped paying for the experiments. Their support is gratefully acknowledged. My supervisor has been Professor Ragnvald Hpier. Thanks a lot for the encouragement and guidance during these years Ragnvald!
Important parts of the work were carried out during a four months stay at the Max- Planck-Institut fur Metallforschung in Stuttgart in 1995. I am indebted to Professor Manfred Ruble for the invitation to visit Stuttgart and to Dr. Joachim Mayer for his hospitality and valuable discussions and suggestions to my work. Jurgen Plitzko and Jorg Thomas are acknowledged for their assistance with the Zeiss and VG microscopes. Thanks also to the rest of the staff and students at MPI for making my stay a pleasant and instructive one.
Furthermore the helpful discussions and comments of colleagues at SINTEF (Foundation of Scientific and Industrial Research) Materials Technology and NTNU are gratefully acknowledged. In particular I would like to thank Dr. Katrin Nord- Varhaug and Dr. Jon Herman Ulvenspen whose works with titanium diboride became the starting point of mine, and Dr. Asgeir Bardal and Associate Professor Mari-Ann Einarsrud who also reviewed parts of this manuscript. I would also like to express my gratitude for the valuable discussions and the interest in my work shown by Dr. Dag 0vreb0 at Norsk Hydro Research Centre in Porsgrunn.
Some people deserve an extra thanks for their kindness and practical help: * Geir Sagmo at SINTEF taught me how to use the hot press * Egil Skybakmoen at SINTEF did the aluminium exposure experiments * Kjell Ramspskar at the Physics Department cut my hard samples * Kjell Muller at SINTEF did the elemental mapping with the SEM * everybody in the Electron Microscopy Group at NTNU for making Trondheim an interesting and inspiring place to work.
Finally, thanks to my family for care and support and to God for always looking after me.
Trondheim, April 1997
Gunnar Pettersen ii Ill
1. INTRODUCTION 1
2. THEORETICAL BACKGROUND - REVIEW OF PREVIOUS WORK 3
2.1 Aluminium electrolysis 3 2.1.1 The Hall-Heroult process 3 2.1.2 Inert, wetted cathodes 5 2.1.3. Energy consumption 6 2.1.4 Demands for a new cathode material 7
2.2 Titanium diboride - properties, production and use 7 2.2.1 Properties of TiB2 7 2.2.2 Production of TiB2 materials 12 2.2.3 Uses of TiB2 14
2.3 Sintering of TiB2 15 2.3.1 Introduction to sintering 16 2.3.2 Literature review 18 2.3.3 Concluding remarks 22
2.4 Degradation of TiB2 cathodes during aluminium electrolysis 23 2.4.1 Literature review 24 2.4.2 Possible mechanisms of aluminium penetration 28 2.4.3 Concluding remarks 33
3. EXPERIMENTAL 35
3.1 Materials selection and preparation 35
3.2 Microstructural characterisation 38 IV
4. RESULTS 41
4.1 The TiB2 powder 41
4.2 Hot-pressing 42
4.3 Microstructure of hot-pressed samples 44 4.3.1 Pure TiB2 44 4.3.2 TiBa with boron addition 52 4.3.3 TiB2 with titanium addition 57
4.4 Microstructure after exposure to liquid aluminium 70 4.4.1 Pure TiB2 70 4.4.2 TiB2 with boron addition 79 4.4.3 TiB2 with titanium addition 81
5. DISCUSSION 87
5.1 The TiB2 powder 87
5.2 Hot-pressing 90 5.2.1 Pure TiB2 90 5.2.2 TiB2 with boron addition 93 5.2.3 TiB2 with titanium addition 95
5.3 Aluminium exposure 100 5.3.1 Penetration of aluminium 100 5.3.2 Reactions with secondary phases 105 5.3.3 Degradation mechanisms 109
5.4 Further challenges 113
6. CONCLUSIONS 117
Appendix 1 Hot pressing parameters Appendix 2 Thermodynamical data and calculations SUMMARY
The introduction of an inert, wetted cathode in the aluminium electrolysis cell makes possible large savings in energy consumption during aluminium production. Titanium diboride (TiB2) holds special promise as a cathode material thanks to its low solubility in and high wettability by liquid aluminium. However, a limited lifetime during electrolysis has so far prevented industrial use of TiB2 cathodes. Degradation of the material is believed to be caused by penetration of aluminium into the cathode during electrolysis, and is closely linked to the impurities present in the TiB2 material after sintering. The aim of this work has been to achieve a better understanding of the sintering process and the mechanisms of aluminium penetration in TiB2. Detailed investigations of the microstructural development during sintering and penetration have provided the basis for such an understanding, using transmission electron microscopy (TEM) as the main characterisation tool.
Hot pressing at 5.0 MPa in an argon atmosphere at temperatures in the range 1790 - 1960 °C was employed to achieve high purity and high density TiB2 compacts. Increasing the sintering temperature increased the density from 90 to 96% in the samples sintered with no additives, but grain growth caused extensive microcracking at the highest temperature. In some samples 2 wt% of elemental boron or titanium was added to investigate the effect on the sintering behaviour. These elements did not dissolve in the TiB2, but caused formation of secondary phases in the grain junctions. Adding boron resulted in a slight increase in density and in formation of B4C. The reducing atmosphere set up by the graphite elements of the hot press is the origin of the carbon, and also removed the surface oxides present on the TiB2 powder. Evaporation of boron oxide is another possible mechanism of oxygen removal, both mechanisms depend upon an open network of pores to exist for gas transport to be possible. The absence of any grain boundary phase after sintering was proven by high resolution lattice imaging. The addition of titanium resulted in high densities, 99%, to be reached even at the lowest sintering temperature thanks to the introduction of liquid phase sintering. Due to the fast densification and closure of the pores the surface oxides have not been reduced by the sintering atmosphere, but has instead formed a Ti-O-rich liquid. Upon cooling this wetting liquid phase has crystallised into four different oxide phases located in the grain junctions of the sintered material. In VI
addition the liquid sintering has resulted in a strong segregation of iron to the grain boundaries. The equivalent of one monolayer of iron was found in a two nanometer wide region at the grain boundary, but lattice imaging still showed the grain boundary to be clean.
To investigate the aluminium penetration into the various TiB2 materials they were exposed to liquid aluminium at 980 °C for 24 hours. All the materials were penetrated in this test, but the amount and appearance of the penetrating aluminium depend on the sintering aid used. Due to the open porosity and microcracking the pure (pure = no sintering additives) TiB2 materials were easily penetrated and the pores were filled up with metallic aluminium. The increased density of the samples with boron addition prevented this rapid penetration. However, grain boundary penetration took place in both kind of materials and resulted in thin, typically 10 nm, wetting films of metallic aluminium at some grain boundaries. Other boundaries were clean, as judged from high resolution lattice images, or contained small lens-shaped aluminium particles. The difference in penetration behaviour between individual grain boundaries is believed to be caused by differences in grain boundary energy. No relation to segregants or grain boundary misorientation could be found, however. The orientation of the grain boundary plane and dewetting of thin films upon cooling of the samples are suggested as explanations for the observed microstructural development, but this needs further investigation for confirmation. The samples sintered with titanium addition suffered from extensive penetration in spite of their high density. Un like the other materials the grain boundaries became faceted and contained thicker films of metallic aluminium. This is believed to be caused by to the iron segregations increasing the solubility, and thus the rate of dissolution, of TiB2 in an iron-aluminium grain boundary phase. Finally all the secondary phases present in the grain junctions after sintering, except from the B4C phase, were found to react with the penetrating alu minium. The reactions were not, however, found to cause swelling and cracking of the materials as has been suggested by other authors. 1
The present work is motivated by the request for a wettable, inert material to replace the carbon cathodes currently used in aluminium electrolysis cells. This will in turn allow considerable improvements in terms of lowered energy consumption. Thanks to its inertness and excellent wettability by liquid aluminium, titanium diboride (TiB2) has for many years been considered the most promising choice of material for a new cathode. Despite the development work and many test runs performed by the aluminium industry, it has not yet been implemented in large scale production of aluminium, however. This is mostly due to the high manufacturing costs of TiB2 cathodes and the limited lifetime caused by the severe cracking and degradation often experienced during electrolysis.
Earlier attempts at sintering TiB2 compacts from a variety of powders and sintering additives have shown that a both pure and dense material is difficult to produce. This is believed to be a major problem when the material is to be used in aluminium electrolysis as high purity and low porosity are considered essential to avoid cathode degradation. From a review of literature in the field it seems to be generally agreed that penetration of liquid aluminium into the sintered TiB2 during electrolysis is part of the cause for the accelerated degradation. This penetration is believed to take place at the grain boundaries and through the pores of the material, and is related to the wettability of the system. Furthermore, the presence of various intergranular secondary phases is believed to play an important role in the degradation process in many TiB2 materials through reactions with aluminium.
To overcome the problems encountered so far when trying to implement the use of TiB2 as a cathode material, a better understanding of the mechanisms that control the sintering process and the cathode degradation during electrolysis is necessary. Detailed knowledge of the microstructural evolution of TiB2 during sintering and aluminium electrolysis is important to obtain such an understanding. In the present work transmission electron microscopy (TEM) has been the main characterisation tool due to its unique ability to collect both structural and chemical information from nanometer-sized areas. Special emphasis has been put on the characterisation of the grain boundaries as these are believed to be very important both during sintering, 2 Development of microstructure in titanium diboride ceramics aluminium penetration and cathode degradation. In many other ceramic systems TEM has proven the existence of thin intergranular films and grain boundary segregations which are important to the sintering behaviour and various physical properties. To the author's knowledge the present work is the first to investigate the grain boundaries of TiB2 at an atomic scale.
The experimental work includes sintering and exposure to liquid aluminium of various TiB2 materials. TiB2 samples, hot pressed without any sintering additives or with small additions of boron and/or titanium were made. The choice of sintering additives was based on the possibility that solid solubility would limit the formation of secondary phases. Also no previous reports on the effect of these additives have been published. The main objectives of the work have been the following:
-To give a detailed description of the microstructure of the sintered materials with emphasis at identification of secondary phases and the effect of sintering additives. -To identify any possible grain boundary films or segregations present after sintering. -To interpret the above in terms of possible sintering mechanisms.
-To describe the penetration behaviour of aluminium at the grain boundaries during exposure. -To identify chemical reactions between secondary phases formed during sintering and the penetrating aluminium. -To interpret the above in terms of penetration and degradation mechanisms during aluminium electrolysis.
The thesis is here presented in the form of a monography. It will later be published in parts as articles in scientific journals. Some of the results have earlier been presented in conference proceedings and as lectures and posters at various international conferences in the field of electron microscopy (Pettersen et al. 1995, 1995b, 1996, 1996b). 3
2. THEORETICAL BACKGROUND - REVIEW OF PREVIOUS WORK
This chapter provides a background for the present work. It starts with an introduction to the field of aluminium electrolysis, explaining the current interest for development of a new, wetted cathode material and describing the environmental stresses such a material must be able to withstand. Then the basic properties of TiB2 are presented along with some methods for manufacturing TiB2. Finally, a more detailed review of previous works and theories relevant for understanding the sintering process of TiB2 and the degradation mechanisms operating during aluminium electrolysis is given.
2.1 Aluminium electrolysis
Aluminium is the most abundant metallic element and makes up eight percent by weight of the earth’s crust. Due to its reactivity it is never found as a pure metal, but forms various compounds with oxygen and other elements. Metallic aluminium was first isolated by H.Davy in 1808, and was later produced by reducing aluminium chloride using potassium or sodium. It remained a very rare and expensive metal until the discovery of the Hall-HerouIt-process; in the court of Napoleon 3rd only the guest of honour was served on aluminium plates, the rest had to settle for plates of gold. C.Hall and P.Heroult independently discovered the electrolytic way of producing aluminium from alumina dissolved in a fluoride melt in 1886. Over the last century this process has been refined continuously, but the basic principles remain. Alternatives to the Hall-Heroult-process include electrolysis of chloride melts and carbothermic reduction, but have not been able to compete from an industrial point of view. The Hall-Heroult process is described in detail in several textbooks including Grjotheim and Kvande (1993), Grjotheim et al. (1982) and Sprlie and 0ye (1994), a short introduction is given below.
2.1.1 The Hall-Heroult process
The main features of a conventional aluminium electrolysis cell is shown in figure 2.1. It has carbon anodes intruding from the top and a carbon cell lining as the bottom and 4 Development of microstructure in titanium diboride ceramics sidewalls. On top of the carbon lining there is a pool of liquid alu minium , this is the acting cathode of the cell.
Figure 2.1 Schematic drawing of the main features of a conventional aluminium electrolysis cell.
The electrolyte floats on top of the aluminium pool, and is based on cryolite (Na3AlF6). Small amounts of other fluorides like aluminium fluoride (A1F3), calcium fluoride (CaF2), magnesium fluoride (MgF2) and lithium fluoride (LiF) are usually added to the electrolyte to improve properties such as increasing the electrical conductivity and decreasing the melting point. Alumina is feed into the electrolyte during electrolysis. The alumina level in the electrolyte is typically 2 to 4 weight percent, the cell temperature in the range 940 - 980°C and the cell voltage 4 - 4.6 V. The nature of the complex, ionic species that form in the electrolyte when alumina is dissolved is not precisely known, but the overall cell reaction is
2 A1203 (dissolved) + 3 C (s) = 4 A1 (1) + 3 C02 (g)
Aluminium containing species are reduced to aluminium at the aluminium-electrolyte interface whereas oxygen-rich species react with the carbon anode and form carbon dioxide gas, gradually consuming the anode. Chapter 2: Theoretical background - Review of previous work 5
Carbon is not well wetted by liquid aluminium. In order to have a stable cathode with a good physical and electrical contact between the carbon lining and the alumini um a thick pool of aluminium (10 to 50 centimetres) is kept at the top of the carbon lining as shown in figure 2.1. Strong magnetic fields arise from the high currents in an aluminium electrolysis cell and introduce waves/ripples in the aluminium surface. In order to avoid short circuiting the cell the anode to cathode distance (ACD) must be kept at a safe 4-5 centimetres. This causes a substantial ohmic loss in the electrolyte. Another major drawback of the carbon cathode is the swelling caused by a gradual penetration of sodium and electrolyte. This ultimately leads to failure, and limits the lifetime of the cathode to typically 3-5 years.
2.1.2 Inert, wetted cathodes
New inert and wettable electrode materials which can replace the carbon in the aluminium electrolysis cell have been sought for a long time. Such materials will allow new cell designs including vertical and sloped cathodes, see the reviews of Billehaug and 0ye (1980), Pawlek (1990) and McMinn (1992). A wetted cathode can be drained leaving only a thin film of alumi ni um at the cathode surface, an example is shown in figure 2.2. The ACD can then safely be reduced to below two centimetres, thereby lowering the cell voltage and reducing the energy consumption.
Figure 2.2 Schematic drawing of a possible cell design for an aluminium electrolysis cell with wetted, drained cathodes. 6 Development of microstructure in titanium diboride ceramics
If an inert, wetted cathode is combined with an inert anode one can think of an environmental friendly and very energy efficient closed cell requiring minim u m maintenance and producing only pure aluminium and oxygen. With today ’s carbon anode a theoretical 1.22 kg C02 is produced for each 1 kg of aluminium. In addition evaporation of fluoride gases from an open cell is a threat to the environment close to the production plant. Another environmental concern is the waste from fallured carbon cathodes which contains cyanide salts and large amounts of fluorides.
2.1.3 Energy consumption
A look at the energy consumption of an aluminium electrolysis cell shows a large potential for energy savings, see Grjotheim and Kvande (1993), Todd (1981) and Dorward (1983). Thermodynamic calculations give a theoretical value of 6.34 kWh/kg A1 for the production of aluminium from alumina for the Hall-Heroult process. Typical values for industrial cells have been lowered from 40 kWh/kg A1 a century ago to today ’s average of 14-15 kWh/kg Al. The best results reported are 12- 13 kWh/kg Al. Most of the electric energy supplied are in other words lost as heat to the surroundings. A part of this loss is due to the resistance of the carbon cathode itself. The larger part however, making up some two thirds of the heat loss, is the ohmic loss in the electrolyte. This is proportional to the ACD. Lowering the ACD from 4-5 centimetres in the conventional cell to 2 centimetres in a wetted, drained cell should therefore allow energy savings of 2-3 kWh/kg Al. An increased current density to maintain the heat balance and increased density of C02-bubbles in the electrolyte may reduce these nominal savings somewhat, but will increase the production capacity. The nominal global capacity for electrolytic aluminium production was approximately 22.6 million tons per year in 1995 (Pawlek 1995), of which approximately 5 % is in Norway. Thus an energy saving of 2 kWh/kg Al gives a possible global reduction in energy consumption of 45 TWh/year, or close to half the total Norwegian energy consumption. Chapter 2: Theoretical background- Review of previous work 7
2.1.4 Demands for a new cathode material
From the above description of the operating environment of an aluminiu m electrolysis cell the following demands must be fulfilled by a new cathode material: * wetted by liquid aluminium * chemically inert, i.e. low solubility in liquid aluminium and cryolite * high electrical conductivity * thermally stable up to 1000°C and resistant to thermal shock * satisfactory mechanical strength
2.2 Titanium diboride - properties, production and use
Billehaug and 0ye (1980) and Pawlek (1990) have reviewed earlier work done in search for an inert, wetted cathode material for use in aluminium electrolysis cells. The group of materials called RHM (refractory hard materials), and especially the borides and carbides of titanium and zirconium, are considered the most promising in view of the above stated demands. Of these titanium diboride (TiB2) has been most widely studied and is considered the most promising candidate. TiB2 was first proposed for use in aluminium electrolysis in a patent by Ransley (1958). A review of properties and production methods relevant to the use of TiB2 as a cathode material is given below along with some possible alternative uses.
2.2.1 Properties of TiB2
TiB2 is a covalently bond, brittle ceramic with high hardness and a high melting point. The crystal structure is hexagonal, AlB2-type, with alternating layers of titanium and boron as shown in figure 2.3. It belongs to the P6/mmm space group, and the lattice parameters are a = 0.303034 nm and c = 0.322953 nm according to JCPDS (35-741). The density of TiB2 is 4.52 g/cm3. The binary Ti-B phase diagram as reported by Murray et al. (1987) is given in figure 2.4. The melting point of pure TiB2 is very high, 3225°C, and according to the phase diagram the phase exists in a compositional range from 33.3 to 34.5 atomic percent titanium, i.e. stoichiometric or with a small 8 Development of microstructure in titanium diboride ceramics boron deficiency. Besides TiB2, two other binary phases are included in the phase diagram, TiB and Ti3B4.
Figure 2.3 The hexagonal unit cell of the layered TiB2 structure.
Weight Percent Boron 50 60 70 80 100
Atomic Percent Boron
Figure 2.4 Binary phase diagram for the titanium-boron system. From Murray et al.(1987). Chapter 2: Theoretical background - Review of previous work 9
Most engineering properties of TiB2 materials are not easily given as they in general depend upon testing methods and environment along with microstructural features such as grain size, porosity, secondary phases etc. The microstructure can be significantly varied by use of different raw materials and methods of production of solid bodies of TiB2. Some typical values for polycrystalline TiB2 are given below, when no other reference is given the data are from the review of Cutler (1992). Single crystal materials are rare and have directional properties in general. In light of the demands stated at the end of the previous chapter, the following properties are of special concern:
Wetting properties: The excellent wetting of TiB2 by liquid aluminium is of major importance for the potential use as a cathode material. Nord-Varhaug (1994) measured wetting angles of 10° and 20° respectively for two different TiB2 samples at 990°C in a vacuum of 2-10"8 bar. Other, earlier results indicating a poorer wetting are probably due to an oxide layer on the aluminium drop which was not removed. A large amount of a non-wetting secondary phase as e.g. in composite materials will cause the wetting to deteriorate.
Chemical properties: Pure TiB2 has a very low solubility in aluminium. Finch (1972) and Ransley (1965) have measured mutual solubilities of boron and titanium, and report a solubility product of approximately 110"8 (wt% Ti)-(wt% B)2 at temperatures of 900-1000°C. With this low solubility a simple calculation shows that with a cathodic current density of 1 A/cm2 and 100 % current efficiency the dissolution rate will be ~10"u m/s if the produced aluminium is saturated with titanium and boron from the cathode. In other words it will take approximately three years to dissolve one millimetre of the TiB2 cathode. During normal cell operation the TiB2- cathode is not supposed to be in contact with the electrolyte, but for short periods during start-up this may occur. Dorward and Payne (1983) reports the stability of high purity TiB2 to be satisfactory in an aluminium plus cryolite system. Skybakmoen et al. (1993) reports a dissolution of the grain boundary areas in cryolite, and attributes this to the possible presence of oxide impurities. Although pure TiB2 is practically inert, problems arise once secondary phases, especially oxides, are present in the TiB2 materials. These do not in general have the same inertness towards aluminium and bar), be when TiB cryolite, cause B the Photoelectron approximately Figure Figure 10 Besides a (Pastor, LOG P (BAR)
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Chapter 2: Theoretical background - Review of previous work 11
TiB2 - Fe (2 at% Ti) in the Ti-B-Fe-system with an eutectic temperature of 1320 °C. The Fe-TiB2 section given in figure 2.6 shows a solubility of eight mole percent TiB2 in liquid iron at 1450 °C.
The solid solubility of other elements in the TiB2 lattice is not precisely known in most cases. According to Castaing and Costa (1977) a high mutual solid solubility exists between the many AlB2-type diborides, including TiB2, at high temperatures, but experimental evidence is lacking in most cases according to Telle (1994). Zdaniewski (1985, 1987) have found some chromium (0-3 wt%) and minor amounts of other metals like iron, wolfram, aluminium, silicon and zirconium substituting titanium in sintered TiB2-bodies. Mroz (1994) has produced a wide range of solid solution (Ti,Zr)B2 alloys. The light elements hydrogen, carbon, nitrogen and oxygen are expected to have a very low interstitial solubility in the boron network.
Electrical conductivity: TiB2 has a very high electrical conductivity compared to other ceramics, it can even compete with metals. Typical electrical resistivity values are in the 0.09-0.015 pQ-m range at room temperature and increases at elevated temperatures and with increasing amounts of porosity and secondary phases. Comparison with corresponding values of 10-50 pQ-m for typical carbon cathode materials (Sprite and 0ye, 1994) shows that TiB2 is an excellent choice from an electrical conductivity point of view, capable of reducing the cathodic voltage drop substantially.
Thermal properties: The thermal and electrical conductivities are usually linked, and so the thermal conductivity of TiB2 is rather high, 40-120 W/m-K. Ausen (1993) studied thermal shock resistance by immersion in liquid aluminium at 800°C. No strength degradation occurred in the small specimens tested here, but this can become a problem in larger parts. According to Cutler (1992) diborides have an excellent thermal shock resistance thanks to the high conductivity. Dorward and Payne (1983) found cracking to be a problem when larger plates of TiB2 were subjected to sudden thermal shocks, while more gradual heating/cooling gave no evidential damage.
Due to the hexagonal structure a TiB2 crystal exhibits an anisotropic thermal expansion, the expansion along the c-axis of the crystal being the larger. Typical 12 Development of microstructure in titanium diboride ceramics thermal expansion coefficients are 6.6-10"6/K and 8.6T0' 6/K respectively in the 25- 1000°C region, and increases at higher temperatures. This thermal expansion anisotropy results in highinternal stresses upon temperature changes in coarse grained, polycrystalline samples of TiB2. Theoretical calculations and experimental investigation show that microcracking takes place when the grain size exceeds a critical size of approximately 15 pm, see Dorward and Payne (1983) and Tennery et al. (1981). As this microcracking seriously deteriorates the mechanical properties, it puts strict conditions on the processing to obtain the fine grained materials needed.
Mechanical properties: As most ceramics TiB2 is hard and brittle. Hardness values in the 20-33 GPa range are reported at RT (room temperature), decreasing significantly at elevated temperatures. A high Young’s modulus of 500-550 GPa (RT) decreasing to 460 GPa at 1000°C is expected for dense materials. In microcracked materials this value is lower, Baumgartner and Steiger (1984) found the Young ’s modulus of a heavily microcracked material to be only half that of a fine grained sample. Flexural strength values scatter strongly, but are typically in the 200-600 MPa range at RT, depending on the microstructure and testing method. No decrease in strength at 1000°C is reported. Compressive strength is a factor 3-10 higher. The fracture toughness is reported to be in the range 5-8 MPa-m1/2. Baumgartner (1984) and Dorward and Payne (1983) has measured strength in argon and liquid aluminium environment at 900-1000°C. A slight decrease in strength, 10-30 %, was found in the latter case, but the strength is still considered sufficient for the intended purpose. A more severe strength degradation after 24 hour exposure, up to 90 %, was found in some of the materials tested by Dorward and Payne.
2.2.2 Production of TiB2 materials
Powder production: TiB2 powders that are commercially available are usually made by carbothermic reduction of oxides followed by crushing and milling. The overall chemical reaction is:
Ti02 (s) + B203 (s) + 5C (s) = TiB2 (s) + 5 CO (g) Chapter 2: Theoretical background - Review of previous work 13
Alternatively B4C can replace B203 as boron source. Possible impurities arising from the carbothermic reduction or subsequent milling are discussed in chapter 5.1. Other ways of producing TiB2 include borothermic reduction where the carbon is replaced by boron powder: Ti02 (s) + 4 B (s) = TiB2 (s) + B202 (g)
Alternatively a different source of titanium can be used (Jiang and Rhine, 1993). In mechanical alloying elemental titanium and boron powders are milled together in a ball mill to form TiB2, see e.g. Park et al. (1994) and Radev and Klisurski (1994). Unfortunately this method will introduce impurities from the milling media. Finally some new methods for synthesising fine-grained, high-purity TiB2 powders without milling have been introduced recently, see chapter 2.3.2.
Densification: High purity, single crystal TiB2 rods 1 cm in diameter and 6 cm in length have recently been grown by Otani and Ishizawa (1994) by a floating zone method. Growing single crystals is not likely to become an inexpensive way of making TiB2 cathodes though. Sintering of TiB2 powder is the most common way of producing monolithic TiB 2 parts, and is presented in more detail in chapter 2.3. Being a faster process than sintering, SHS (Self-propagating High-temperature Synthesis) is yet a different approach to production of monolithic TiB2. Elemental powders of titanium and boron in a 1:2 ratio is mixed. Once ignited by an “electric match” the heat generated in the strongly exothermic reaction between elemental titanium and boron will make the reaction propagate spontaneously through the sample in a few seconds. By a strong, dynamic compaction some seconds after ignition Wang et al. (1994) obtained 99% dense TiB2 with some TiC contamination.
Coating: An approach different from the production of monolithic blocks by sintering of TiB2 powders is coating. A substrate, for instance a carbon block, is covered by a layer of TiB2, thus reducing the overall material cost if a dense, well-adherent coating can be obtained. CVD (Chemical Vapour Deposition) utilises hydrogen gas to reduce halide vapours to TiB2 according to the reaction:
TiCU (g) + 2 BC13 (g) + 5 H2 (g) = TiB2 (s) + 10 HC1 (g) 14 Development of microstructure in titanium diboride ceramics
By controlling the processing parameters (temperature, flow rate and partial pressures of gases) a slow growth of a dense TiB2 layer on a suitable substrate can be obtained, see Besmann and Spear (1975) and Capulto et al. (1985). By laser-induced CVD higher deposition rates are reached, different rate limiting mechanisms are discussed by Elders and Voorst (1994). Few results from exposure of CVD-coated materials to aluminium are reported. Becker and Blanks (1984) found a lifetime 50 times below what was expected from the solubility product during aluminium electrolysis with cathodes that had been CVD-coated with 700 pm of TiB2.
In Electrochemical deposition TiB2 can be deposited onto a carbon cathode in an electrolysis cell as current is run through a fluoride or chloride-fluoride melt containing titanium and boron ions. Depending on temperature, ion-concentrations, voltage and current densities either a smooth film or a loosely bound powder of TiB2 can be deposited, see Makyta and Utigard (1993) and Grjotheim et al. (1988) for details. Only thin films have been possible to grow so far, but high purity is reported and Wendt et al. (1992) claim an excellent stability of the deposition product when exposed to aluminium.
Various sputtering techniques (Riviere et al. 1991, Chen and Barnard 1995), PVD (physical vapour deposition, Knotek et al. 1992), evaporation techniques and plasma spray have also been used to coat substrates with thin films of TiB2 (Hollum and Ulvenspen, 1989). These materials are however not intended for use in aluminium electrolysis cells, and no testing for such an application is reported to the authors knowledge.
2.2.3 Uses of TiB2
Various other applications of TiB2 besides the use of monolithic blocks or coatings of pure TiB2 in aluminiu m electrolysis cells have been proposed. One idea that has been intensively studied recent years is to replace the carbon cathode with a TiB2-carbon composite where typically 20-90 % TiB2 particles are added to the carbon lining, or alternatively to coat the carbon with such a paste. The idea is that this will improve the wetting and conductivity and reduce the degradation of the cathode at a far less cost than a pure TiB2-cathode. The wetting is however not as good as with a pure Chapter 2: Theoretical background - Review of previouswork 15
TiB2 cathode and some degradation still occurs, meaning that the full potential for energy savings available for monolithic TiB2 cathodes can not be exploited. Some examples of development and pilot-plant operation of TiB2-carbon composite cathodes are found in Boxall et al. (1984), McMinn (1992), Liu et al. (1992), Watson and Toguri (1991) and Xue and 0ye (1994).
Due to its unique set of properties TiB2 has been used or proposed for use in a wide variety of applications. The high hardness, Young ’s modulus, thermal conductivity and melting point along with the chemical stability and wear resistance has made TiB2 an attractive material for use in cutting and extrusion tools, see Ulvenspen (1994). To further improve the fracture strength and toughness ceramic composites (TiB2 + ceramic) and cermets (TiB2 + metal) of TiB2 have been designed, see Smid and Kny (1988), Watanabe and Kuono (1982) and Jungling et al. (1994). Although decreasing the high temperature properties and inertness somewhat compared to pure TiB2 these materials can perform well. Flexural strengths of -1000 MPa along with toughness up to 14 MPa-m1/2 are reported by Jungling for a cermet sintered with a Ti-Fe-Cr-Ni binder phase. Surface coatings of TiB2 have been proposed to reduce wear, corrosion and oxidation. The light weight and high impact strength has made TiB2 an armour material in military clothing and vehicles (Cutler, 1992). The high melting point and electrical conductivity along with inertness towards aluminium has led to use as crucibles for inductive melting and evaporation of aluminium. Finally TiB2 is used for grain refinement in aluminium alloys.
2.3 Sintering of TiB?
The strong, predominantly covalent, bonding of TiB2 results in a low self-diffusion coefficient which makes TiB2 a particularly difficult material to sinter to high densities. This section begins with an introduction to the field of sintering, included for the physicist or electron microscopist who is not yet familiar with it, and puts emphasise on methods and mechanisms relevant to sintering of TiB2. For more details a standard textbook such as McColm and Clark (1988) is recommended. After the introduction a review of the literature in the field of sintering and sintering mechanisms of TiB2 is given. Finally an outline of the goals of this work is made. 16 Development of micro structure in titanium diboride ceramics
2.3.1 Introduction to sintering
Utilised for pottery in ancient civilisations sintering, or firing to obtain higher mechanical strength, is one of the oldest technologies known to man. The last decades have brought forth large advances in the understanding of the sintering process, and several new ceramic materials with much improved properties have been developed. During sintering a body formed from ceramic powder particles is heated to form strong bonds between the particles. This is accompanied by shrinkage of the body and a large increase in strength as the porosity is usually reduced from about 50% to only a few percent. The driving force for this densification is the reduction of surface and grain boundary energy when a solid body grows from the separate powder particles. As illustrated in figure 2.7 the result of the sintering process is determined by the outcome of the competition between various coarsening and densification mechanisms. The densification processes cause a net flux of atoms away from the contact point, thus the particle centres are moved closer and the porosity is reduced. The coarsening mechanisms result in neck growth without densification as the particle centres do not move.
Figure 2.7 Possible mechanisms of densification (1,2) and coarsening (3,4) during solid state sintering. (1) Bulk diffusion (2) Grain boundary diffusion (3) Surface diffusion (4) Evaporation-condensation
The mechanisms shown in figure 2.7 are the most important ones in most cases of solid state sintering. A different mode of sintering is named liquid phase sintering as most of the densification here relies on the presence of a liquid phase between the Chapter 2: Theoretical background- Review of previous work 17 solid powder particles. In the first stage of liquid phase sintering rearrangement of the particles into a closer packing causes significant densification. Another densification mechanism contributing to make liquid phase sintering faster than solid state sintering is the dissolution - precipitation mechanism. Ceramic material is dissolved into the liquid phase from the high-pressure, concave neck regions and reprecipitates at the low-pressure, convex surfaces. This causes a faceted grain shape with straight grain boundaries to develop during sintering. The final stage of densification in liquid phase sintering involves also the solid state mechanisms given above. The main disadvantage of liquid phase sintering is the incorporation of a secondary phase which deteriorates the high temperature properties and may cause corrosion problems unless it can be dissolved in the matrix phase.
Due to the high temperatures involved, grain growth is frequently found to occur during sintering. This is a problem as a larger grain size in general deteriorates the mechanical properties, and especially so for TiB2 in which microcracks form due to the high internal stresses set up by the thermal expansion anisotropy when the grain size exceeds ~15 pm. Grain growth also makes the final densification more difficult and slow as the larger grain size means longer diffusion paths for vacancies. In the early stages of sintering the pores reduces grain growth and are connected to form a continues, open network of porosity. When the density exceeds approximately 92-94 percent of the theoretical density, only isolated pores in the grain boundaries and grain comers are left (Kuczynski 1975, Coleman and Beere 1975). If the grain growthis too rapid in this last stage the grain boundaries can break away from the pores and leave them inside large grains, this may happen during exaggerated grain growth.
Several parameters can be played with to enhance or depress the individual sintering mechanisms, the effect of the most important ones are sketched here: The powder itself is usually the most important factor. A finer grain size speeds up all sintering mechanisms as transport paths become shorter, but is especially effective in increasing the contribution of grain boundary and surface diffusion. As seen below, submicron powders seem to be necessary for the successful solid state sintering of TiB2. Also the particle size distribution and particle shape is of interest as it influences the driving force for grain growth and the green density. Sintering aids, either deliberately added or as impurities arising from the powder production, often play crucial roles. Suitable additives are sought which can either 18 Development of microstructure in titanium diboride ceramics
enter the main phase in solid solution to increase bulk diffusivity or segregate to the grain boundaries or surfaces to control the diffusion and/or evaporation kinetics here. Also finely dispersed small, stable particles can be added at the grain boundaries to give a pinning effect which reduces grain growth. Finally liquid forming phases can be added to encourage liquid phase sintering. In all cases the effect of the sintering aid to the interesting properties of the final product has to be considered in order to choose an additive that is compatible with the desired application. Physical parameters include temperature, time, sintering atmosphere and pressure. Increasing the temperature speeds up all sintering mechanisms, but generally by different rates. Too long holds at hightemperatures might cause grain growth. A well defined sintering atmosphere, usually in the form of a vacuum, an inert gas or a reducing atmosphere, can in some cases be used to control the surface chemistry of the powders. Finally an external pressure can be applied to increase the driving force for densification and introduce plastic flow and diffusional creep as densification mechanisms. This is done either by hot uniaxial pressing (HP) or hot isostatic pressing (HEP) at pressures ranging from tens to hundreds of MPa. Both methods increase densification and allows lower temperatures and less sintering aids to be used, reducing the final grain sizes and content of secondary phases. The high costs and low productivity limits their applications to advanced materials which are not possible to compact by pressureless sintering. HP also limits the shape of the final products to simple geometries.
2.3.2 Literature review
Both solid state and liquid phase sintering have been employed for densification of TiB2 materials. Pastor (1977) in a review of early attempts on sintering of boride materials, TiB2 and ZrB2 in particular, concluded that high densities could only be reached by activating the sintering process with metallic additives like Fe, Ni and Cr. Telle (1994) in a review of boride and carbide ceramics attributed the difficulty in sintering of TiB2 to the high melting point and the comparatively high vapour pressure of the constituents causing coarsening by evaporation-condensation. Some more recent articles are reviewed below. Chapter 2: Theoretical background - Review of previouswork 19
Solid state sintering: The development of new powder production techniques giving purer and more fine grained TiB2 powders has made sintering to full density possible. Baumgarter and Steiger (1984) obtained densities of more than 99% during pressureless sintering of a high-purity, submicrometer-sized powder. The powder had been synthesised by hydrogen reduction of halogen vapour reactants in a plasma arc, following the same reaction as described for CVD-coating above. At low temperatures neck-growth without densification was observed. This was attributed to surface diffusion as no weight loss indicating evaporation was recorded. At temperatures around 2000 °C essentially full densification was reached without severe grain growth when 0.8 wt% of carbon was added. The densification was attributed to bulk diffusion of vacancies to the grain boundaries in the fine-grained (~1 pm) material. Rapid grain growth was found to occur at temperatures above 2100 °C. The levels of metallic impurity elements were below 100 ppm in these materials. The lack of success when sintering carbothermic powders was attributed to the onset of rapid grain growth before complete densification could be reached.
Finch et al. (1986) and Balk and Becher (1987) studied hot pressing and pressureless sintering of a fine-grained TiB 2 powder made in a reaction between titanium metal and BCI3. Densities up to 99% after hot pressing at 1600 °C and 96% after pressureless sintering at 2050 °C were obtained. The effect of the main impurities, oxygen, carbon, nickel and iron, on the sintering mechanism was studied. Oxygen in the form of surface oxides was found to have a deleterious effect during sintering. At the lower temperatures employed during hot pressing evaporation of B203 from the surface oxide is believed to be the cause of neck growth and coarsening. When the oxygen level exceeds 1.5 wt% serious grain growth took place before 90% density could be reached by hot pressing. At the higher temperatures employed during pressureless sintering most of the B2Q3 has already evaporated, thus reducing the effect of the evaporation-condensation mechanism during densification. Still high oxygen levels was found to cause substantial grain growth however, and surface diffusion in a Ti02 rich surface oxide left after the B203 evaporation was believed to be the reason for the coarsening. Densification was attributed to the grain boundary diffusion mechanism. Carbon additions can be used to reduce the oxygen level by evaporation of CO to allow successful sintering of high-oxygen powders. In addition the incorporation of 20 Development of microstructure in titanium diboride ceramics
small carbon particles is necessary to inhibit grain growth and reach high densities. The metallic impurities Fe and Ni were found to improve sinterability of the powders. With a Fe + Ni level of 0.1 wt% a maximum density of 94% was reached, whereas with 0.5 wt% Fe + Ni materials of 99% density were hot pressed. This might possibly be attributed either to the formation of small amounts of liquid phase or to improved grain boundary diffusion as these elements segregate to the grain boundaries.
Hill et al. (1989) has made a relatively fine-grained TiB2 powder (1 by 10 pm platelets) in a continuous carbothermic process. As no milling is necessary this powder has a low impurity content compared to ordinary carbothermic powders, only 0.3 wt% of oxygen and carbon and 200 ppm iron. The pure powder could not be sintered to high densities, but upon addition of 0.4 wt% of metallic additives as sintering aids the powder could be hot pressed to 99% density. Another 2.5 wt% CrB2 as grain growth inhibitor allowed even pressureless sintering to 98% density. Itoh et al. (1990) demonstrated the beneficial effect of a reaction between elemental boron and titanium taking place during sintering of a (TiB2 + 0.2 B) + 0.1 Ti powder. The powder was produced in a solid state reaction between TiN and boron and had a particle size of 0.5 - 2 pm. Hot pressing at 1800 °C gave a 98% dense single phase TiB2 material with little grain growth. Nothing was said about the impurity content or the sintering mechanisms.
Pressureless sintering of ordinary carbothermic TiB2 powders is shown by Kang et al. (1989) and Torizuka et al. (1995) to depend on the use of sintering aids. Kang obtained 95% density from a powder with a grain size of 0.8 pm by adding 0.5 wt% Fe as sintering aid and 10 wt% B4C to avoid abnormal grain growth. A TEM investigation showed an iron-rich phase to be present in the grain junctions and at grain boundaries. The existence of liquid iron at the grain boundaries during sintering was believed to be the reason for the improved densification. Torizuka et al. (1995) obtained 96% density during pressureless sintering by adding 2.5 wt% of SiC to a TiB2 powder of 2 pm grain size with a 1.5 wt% oxygen content. From SEM and TEM observations a reaction between the SiC and the surface oxide to form TiC and amorphous Si02 was confirmed. The existence of a liquid Si02 phase during sintering was proposed to be the reason for the improved densification. Chapter 2: Theoretical background - Review of previous work 21
Liquid phase sintering: As seen from some of the above examples small additions of metals are often used to "activate" the sintering process, giving a diffuse borderline between the solid and liquid sintering terms. Larger additions of metals are sometimes used to reduce the sintering temperature further and make composite materials with improved mechanical properties named cemented borides, cermets, hard metals etc. According to the reviews of Pastor (1977) and Telle (1994) 5 to 25 wt% of metals like Ni, Co, Cr and Fe or borides hereof are often used as binders. In contact with TiB2 they all form liquid phases at temperatures of 900-1300 °C which dissolves some TiB2 and wets the TiB2 phase. Liquid phase sintering to high densities at temperatures of 1400-1600 °C is thus possible, the liquid phase giving fast grain growth. In many cases brittle metallic boride phases form during cooling of the liquid phase and deteriorates the mechanical properties of these materials.
Jungling et al. (1994) showed that the formation of the brittle Fe2B phase when using Fe or Fe-Cr-Ni binders was due to C or B4C impurities, and was accompanied by Ti203 formation from the surface oxides. Sintering temperatures of 15-1700 °C was used to reach densities of 95-99% by pressureless sintering with 5-30 wt% binder content. The Fe2B formation could be avoided by a small addition of Ti to the binder, and resulted in materials with promising hardness-fracture toughness combinations at elevated temperatures.
Einarsrud et al. (1997) studied pressureless liquid phase sintering of TiB2 at 13-1700 °C with 1-5 wt% of Ni, NiB or Fe as sintering additives. A density of 94% was reached at 1500 °C with only a 1.5 wt% Ni addition. Higher sintering temperatures resulted in exaggerated grain growth which could only be suppressed by a carbon addition at the expense of reduced density. The exaggerated grain growth increased with increasing oxygen content, and was attributed to surface diffusion in a titanium oxide surface layer. The densification rate, as determined by the amount and content of sintering additives, and final density was found to be decisive for the level of oxygen in the final product. This was attributed to the mechanism of oxygen removal which was argued to be evaporation of B202 by the reactions
2 TiB2 (s) + 14 Ti02 (s) = 8 Ti203 (s) + 2 B202 (g) 22 Development of microstructure in titanium diboride ceramics
2 T1B2 (s) + 7 B2O3 (1) — T12O3 (s) + 9 B2O2 (g)
Evaporation of the surface oxides depends on the presence of open porosity. A rapid densification will cause the oxygen still not evaporated when closed porosity is reached to be trapped inside the sintered body.
2.3.3 Concluding remarks
From the above there seem to be agreement of the effect of various impurities and additions to the sintering behaviour of TiB2, yet uncertainties arise when the precise mechanism(s) of densification or coarsening are asked. Complete densification by solid state sintering is only possible by choosing high-purity powders with submicrometer grain size. The surface oxide of the powder should be minimised as it results in coarsening and encourages grain growth. Carbon or carbide additions have proven helpful in reducing the oxygen level and hence inhibiting grain growth. Small additions of metallic elements are very effective in improving the sinterability. Finally larger additions of metals like Fe, Ni and Cr are used to facilitate liquid phase sintering, but cause formation of intergranular metallic and oxygen-containing phases. Extensive grain growth seems to be the most important hindrance towards complete densification, it also seriously deteriorates the mechanical properties of the sintered body.
The materials sintered in the present study were intended for use in aluminium electrolysis cells, thus no intergranular secondary phases that could be attacked by aluminium were allowed. This puts strict conditions on the use of sintering aids. Simultaneously a low porosity was desired, thus hot pressing was applied to enhance densification. As the carbothermic powder available in this study could not be hot pressed to full density by itself a choice of two sintering aids was made, boron and titanium. The reason for choosing boron and titanium as additions was primarily that these two elements were the only ones already present in the system. The binary phase diagram B-Ti in figure 2.4 shows there to be a certain stoichiometric range in which the TiB2 phase can exist. Thus it was hoped that small additions of boron or/and titanium would dissolve in the TiB2 matrix during the sintering process and not form new secondary phases. Boron is a solid at the sintering temperatures used. It was Chapter 2: Theoretical background - Review of previous work 23 hoped to reduce the surface oxide by evaporation of B202, and possibly to pin grain growth. Titanium on the other hand is liquid at the sintering temperature and was hoped to improve densification through liquid phase sintering mechanisms. The combined addition of boron and titanium was hoped to give a reactive sintering contribution similar to that described by Itoh et al. (1990), but at far smaller levels of addition.
To the author's knowledge small additions of boron and titanium have not been tried as sintering aids before. Their influence on the sintering mechanism and sinterability of TiB2 are therefore of interest. Thorough investigations of the microstructure formed after sintering by optical and electron microscopy are hoped to contribute to a better understanding of the sintering mechanisms of TiB2. During the TEM-studies emphasis will be put at identifying any secondary phases present, and study the grain boundaries at nanometer scale. As proposed by many authors above oxide layers and grain boundary segregants are believed to play an important part during sintering of TiB2, both in terms of evaporation-condensation and diffusional transport behaviour. Thin intergranular oxide films with a composition different from the matrix phase have been observed by TEM in several different ceramics including Si3N4, A1203, sialons and Zr02 (Clarke 1987). These may form a liquid phase during sintering, and are known to be of great importance to both sintering behaviour and materials properties. No similar studies of grain boundary structure and chemistry in TiB2 at atomic level has been made until now.
2.4 Degradation of TiB? cathodes during aluminium electrolysis
During aluminium electrolysis the cathode surface is continuously exposed to the fresh, liquid aluminium that is being produced at a temperature of approximately 950 °C. If dissolution of TiB2 in this aluminium was the only mechanism to degrade the cathode, it would take approximately three years to remove the upper one millimetre (chapter 2.2.1), a very promising result. Unfortunately this does not usually seem to be the case. Many investigators have reported a much faster degradation, the rate depending on the microstructure of the TiB2 material in use. Several different mechanisms of degradation have been suggested to explain this behaviour. This chapter gives a short review of previous experiments with TiB2 in contact with liquid 24 Development of microstructure in titanium diboride ceramics aluminium as reported in the literature. Emphasis is put on microstructural evolution and degradation mechanisms. Then different possible mechanisms of aluminium penetration in TiB2 are discussed. Finally the current status is summarised and goals for this work set up.
2.4.1 Literature review
Ever since Ransley (1958) first proposed the use of TiB2 as an inert, wetted cathode for aluminium electrolysis, work has been going on to qualify this material for use in industrial cells. Most of this work has taken place in the laboratories of the larger aluminium producers and their suppliers, and are thus not always published in detail. Some reviews and patents from this work exists, however, along with publications from university groups.
The most common mode of failure reported from aluminium exposure tests is cracking and spalling of the TiB2 material after some time of electrolysis. This was observed already by Ransley (1962), working for British Aluminium, who attributed the cracking of certain TiB2 materials to poor control of the chemical composition. He claimed, however, (Ransley 1965) that neither of the elements B, C, N, O or Fe had any apparent deleterious effect during electrolysis as long as the content of each were kept below 1 % by weight. The density of these TiB2 material should be kept above 90 %, and successful tests with no cracking during five months of electrolysis were reported. No investigations of aluminium penetration or microstructure were reported. According to a review by McMinn (1992) materials developed by Reynolds Metals Company in the same period would often develop cracks, and eventually fail after typically 6 months. The cracking was due to a slow, intergranular corrosion by penetration of aluminium, this corrosion was believed to be caused by low levels of C, O and Fe impurities present at the grain boundaries. Probably the materials from British Aluminium suffered from the same degradation as their interest faded away after the initial optimism.
Gesing and Wheeler (1987) summarised work done by Alcan/Eltech in the mid 80's to evaluate different TiB2 materials. They found that the corrosion and cracking of TiB2 was caused by reactions between secondary phases and aluminium penetrating along Chapter 2: Theoretical background - Review of previous work 25
the grain boundaries. Harmful secondary phases were oxides present at grain boundaries which formed A1203 and carbon which reacted to AI4C3 in grain junctions. These reactions were both claimed to result in an increased volume of secondary phases, build-up of high stresses and finally cracking of the material due to these stresses. TiC was not found to react, whereas B4C reacted, but did not cause cracking. Increasing the density reduced the penetration rate, but it is suggested that failure occurs by the same mechanisms, only somewhat delayed. Only in high purity materials interconnected porosity did not lead to accelerated grain boundary attack. Unfortunately no further details from their experiments are reported.
Dorward and Payne (1983) of Kaiser Aluminium reports a comprehensive effort in development of TiB2 cathodes in a DOE-sponsored programduring the years 1975-83. They investigated several different materials, and gave the following specification for a suitable cathode material: the material should be well sintered with fine, equiaxed grains and a density of more than 98 % of theoretical density. Simultaneously high purity is needed; more than 99.7 % TiB2, no secondary phases in the grain boundaries and below 500 ppm of oxygen. Their best material showed no cracking and only a uniform dissolution in accordance with the equilibrium solubility in short term (39 days) electrolysis tests. From extrapolation of these results they predict a three year lifetime for the 6 mm thick TiB2 cathode plates. No detailed microstructural information concerning microstructure and aluminium penetration is given for this material. In another material the fracture mode was observed to change from transgranular to intergranular as aluminium penetrates the TiB2. Microprobe analysis from fractured and polished surfaces revealed the presence of metallic aluminium in grain boundary pores and at some of the grain boundaries (Dorward 1982, 1985). Cracking from thermal shock were often found to occur if the parts were not properly handled. Spalling upon cooling of partly penetrated parts were suggested to be the result of a reduction in the thermal expansion coefficient of the outer, penetrated layer.
The best TiB2 material for cathode application is probably that sintered from the expensive, fine grained powder produced by PPG industries in a gas-phase reaction (Baumgartner 1984, 1984b). After having acted as a cathode for 6 months, aluminium had only penetrated the grain boundaries to a depth of 400 pm in this very high purity material. STEM showed that virtually all grain boundaries within the penetration zone contained aluminium along with small amounts of Fe, Si and P not present prior to 26 Development of microstructure in titanium diboride ceramics electrolysis. Grain boundary diffusion was suggested to be the mechanism of penetration, Fe, Si and P being impurities from the pot metal and electrolyte following the aluminium. Small precipitates, 0.1 pm, containing aluminium were found at some grain boundaries. Aluminium penetration reduced the mechanical strength and fracture toughness of the material at 970 °C by 10-30 % depending on the grain size, and liquid metal embrittlement was argued to be the reason for the reduction. The aluminium penetration also changed the fracture mode from transgranular to intergranular.
Zdaniewski (1985, 1986, 1986b) has investigated microstructures and fracture in several TiB2 and ceramic composite materials in liquid aluminium environment. Differences in penetration behaviour between the materials are believed to be linked to impurities at the grain boundaries. Surface active elements segregating to the grain boundaries will alter the interfacial energy ratio and thus the penetration behaviour. No detailed study of interface chemistry to support this was reported, however. In a material with 17 wt% TiC cracking occurred as the TiC phase reacted to form AI4C3 and cause swelling of the grain boundaries. Segregations of Cr to the surface of a (Ti,Cr)B2 solid solution material caused microcracking and aluminium penetration. This was explained by tensile stresses set up by a decrease in unit cell volume along the surface due to the Cr-substitution.
Liquid phase sintered TiB2 with nickel as sintering aid were immersed in aluminium at 950°C for 100 hours by Finch and Tennery (1982). The density of these materials varied between 90 and 99 %, the residual nickel content were in the 0 to 10 % range and the oxygen level varied between 1 and 8 % by weight. Prior to exposure Ni3B and oxide-phases were found at the grain boundaries. In all materials the grain boundaries were readily penetrated by aluminium. Surprisingly, least cracking was found in a low density material with high oxygen content. No detailed study of the microstructure to explain this behaviour was given.
Wendt and Dermeik (1990) and Wendt et al. (1992) studied cathodic deposition of aluminium from a chloride electrolyte at 700 °C. The low temperature used with a chloride electrolyte reduces the solubility of TiB2 in aluminium. Sintered TiB2 materials dissolved at a rate ten times higher than the equilibrium solubility due to Chapter 2: Theoretical background - Review of previous work 27
grain boundary attacks and possibly pull-out of grains. Thin, high-purity TiB2 coatings prepared by electrochemical deposition of from FLINAK melts had very low dissolution rates. No details concerning the microstructure of these materials were given.
TiB2 coatings made by CVD were tested in aluminium by Becker and Blanks (1984) and Besmann et al. (1994). Films deposited at low temperatures were unstable. This was attributed to a high chlorine level causing intergranular corrosion and pull-out of individual grains. Deposition at temperatures above 1000°C yielded pure coatings that were reported to be stable for 20 weeks when soaked in aluminium at 1000°C.
Ulvenspen (1994) and Skybakmoen et al. (1993) investigated the stability of TiB2 materials HIP-ed from carbothermically produced powders during unpolarized exposure to a cryolite based electrolyte at 980 °C. Using SEM they found the cryolite to dissolve the grain boundaries in the surface area of their TiB2 materials, and cause high weight losses due to pull out of TiB2 grains. The dissolution of the grain boundaries was believed to be caused by the presence of oxide phases, but this was not proven. During electrolysis the electrolyte did not penetrate into the cathode although aluminium did. It was suggested that cryolite could stay in touch with the cathode in the early stages of electrolysis due to a better initial wetting, and that degradation was caused by the joint effect of cryolite leaching the grain boundaries and aluminium penetrating into the TiB2.
The most thorough microstructural characterisation of TiB2 materials reported is the TEM-work of Nord-Varhaug (1994) and Nord-Varhaug et al. (1992, 1993), summarised in (Nord-Varhaug 1996). A wide range of secondary phases, both metallic and non-metallic, present in triple junctions of HIP-ed TiB2 were identified, and their possible origins discussed. After unpolarized exposure to aluminium at 980 °C for 48 hours the aluminium had penetrated the materials to a depth of -150 pm. Many grain boundaries in the penetrated region contained a film of metallic aluminium with thickness up to 170 nm. Also three different oxide phases were identified as reaction products between the penetrating aluminium and secondary phases present after sintering. 28 Development of microstructure in titanium diboride ceramics
2.4.2 Possible mechanisms of aluminium penetration
The rate of penetration of a liquid metal into a polycrystalline, solid material generally depends heavily on the microstructure of the material and its chemical compatibility with the penetrating metal. In the case of liquid aluminium penetrating a sintered TiB2 body, no reaction occurs between the two phases, and the solubility of the solid phase in the metal is very low. The wettability of TiB2 by aluminium is good, however. Some possible mechanisms of penetration are suggested below. Which one will dominate in an actual case depends on the microstructure of that particular TiB2 sample.
Open porositv/microcracks As complete densification of TiB2 is hard to achieve without extensive use of sintering aids, most samples will contain some residual pores. If the density falls below 92-94 percent of the theoretical density, the material will exhibit open porosity, i.e. a connected network of pores through the structure. In such a sample penetration will be very fast as aluminium is free to run into the pore structure thanks to the good wettability. A similar situation occurs if extensive microcracking opens up paths for the penetrating aluminium. Due to the thermal anisotropy, this will be a problem in materials with a large grain size.
Reactions with secondary phases Most sintered TiB2 materials contain secondary phases originating either from the powder used for sintering, from sintering aids used to enhance densification or from the sintering atmosphere. As opposed to the TiB2 itself, most of these phases are chemically unstable in the presence of liquid aluminium. They will either be dissolved in the aluminium and transported away, or react with the aluminium to form new, solid phases. Examples of such behaviour were given in chapter 2.4.1 above. The reduction of the total Gibbs free energy resulting from these reactions contributes a driving force for the aluminium penetration.
The location of the secondary phases along with the chemical composition and the amount will decide their possible detrimental effects. Two kinds of microstructure will in particular be harmful as they can cause formation of open paths for penetration Chapter 2: Theoretical background - Review of previous work 29 in originally dense materials. If the grain junctions contain a phase which upon reaction with aluminium form a solid phase of increased molar volume, large internal stresses will build up. Eventually this may cause cracks to form, through which aluminium can penetrate further into the material. Another harmful microstructure is the presence of a continuous, wetting phase at all the grain boundaries. Such a thin film of a secondary phase is known to form in several ceramics, both during solid and liquid phase sintering. If this grain boundary phase reacts with or is dissolved by the liquid aluminium, a new fast path of penetration is found is found along the grain boundaries.
Penetration bv wetting of grain boundaries
An important reason for choosing TiB2 as cathode material was just the good wetting of the TiB2 surface by liquid aluminium. Unfortunately this may also make the material more prone to penetration by wetting of the grain boundaries as illustrated in figure 2.8.
Figure 2.8 Wetting of surface (left) and of a grain boundary (right). S = solid, L = liquid, V = vapour.
Wetting of a surface is governed by the Young equation
Ysv = Ysi + YivCosQ
The subscripts of the interfacial tensions, % denotes the solid, vapour and liquid phases respectively. Low values of the wetting angle 0 is equivalent to good wetting. For grain boundary wetting the interfacial tensions of interest are y ss, that of the solid-solid 30 Development of microstructure in titanium diboride ceramics interface, and y sl for the solid-liquid interface. The dihedral angle, 5, as defined in figure 2.8, is given by the equation
Yss = 2 y s! cos(S/2)
Based on this equation a classification into different possible microstructures as shown in figure 2.9 can be made (Smith 1964). For high grain boundary energies, Yss ^ 2 y$i, the dihedral angel is zero. For the liquid phase it is then energetically favourable to spread out and make up a thin film that wets all the grain boundaries. Increasing the dihedral angle, but keeping it below 60° (V3Ysi s Yss < 2 Ysi) forces the film to retract from the grain boundaries. It will still make up an interconnected network of liquid phase following the three-grain channels and grain junctions, however. For dihedral angles above 60° (y« < V3 Ysi) all the liquid will retract to form isolated islands in the multi-grain junctions. Consequently the low y si wanted to assure good surface wetting will also enhance grain boundary wetting.
As pointed out by Clarke and Gee (1992) the interfacial tensions in general depend on several factors including the crystallography of the interface, the temperature and any segregations of impurity elements to the grain boundary. A complete crystallographic and geometric description of an arbitrary grain boundary requires eight independent parameters to be determined. Three degrees of freedom describe the relative rotation of the two grains, two more is needed to give the orientation of the interface plane, and finally three degrees of freedom describe the relative translation of the two crystals. A complete determination is very time-consuming, but neglecting everything but the angle of misorientadon is in many cases found to be a fruitful simplification. The misorientation between two grains is defined to be the smallest rotation of one of the grains necessary to bring it into the same orientation as the other, the rotation axis being arbitrarily chosen. Figure 2.10 shows schematically the variation of y ss and 2 y si with grain boundary misorientation. The solid-liquid interface crystallography has only a small influence on the interfacial tension. For the solid-solid interface there is a larger effect of changing the misorientation between the two grains. Starting at zero misorientation one has no grain boundary at all, and consequently zero interfacial tension. For low-angle grain boundaries there is a linear dependence of the Yss on the Chapter 2: Theoretical background - Review of previous work 31
misorientation, whereas for higher misorientations the interfacial tension is almost a constant. Exceptions are seen as cusps in the graph, and are found at some special high-angle boundaries, like for instance twin boundaries, which has particularly low tensions. From figure 2.10 one can easily determine which grain boundaries should be wetted and which not by comparing the interfacial tensions. In the given example the low angle and some special boundaries are not wetted, whereas most high-angle grain boundaries are. Numerical values for the interfacial tensions relevant to grain boundary wetting of TiB2 by liquid aluminium have not been published. Thus precise predictions of which boundaries will be wetted at what temperatures and with a certain impurity level can not be made.
Figure 2.9 Wetting microstructure for different dihedral angles (Smith, 1964). Upper: Calculated, three-dimensional appearance of secondary phase. Lower: Experimental, two-dimensional microstructure in Cu-alloys. Dihedral angles (left to right): -80°, -50° and -0°.
The liquid-solid interfacial tension is in general more sensitive to temperature variations than the grain boundary tension (Clark 1987). Lowering the temperature may therefore cause dewetting of some grain boundaries as indicated in the second 32 Development of microstructure in titanium diboride ceramics graph of figure 2.10. As shown by Clarke there is an energy barrier to dewetting, some supercooling is required to nucleate a dewetted grain boundary from a wetted one.
Figure 2.10 Interfacial tensions as a function of misorientation at two different temperatures (T2 > T\). From Clarke (1987).
Variations in surface chemistry may critically alter the wetting properties of a surface (Clarke and Gee 1992). Impurity elements segregated to the grain boundaries must therefore be expected to have a significant effect on the interfacial tensions too, and thereby on the grain boundary wettability. Zdaniewski (1985,1987) supported this view, and demonstrated the presence of grain boundary segregations in TiB2 materials. Likewise the presence of secondary phases will alter the wettability.
The kinetics of grain boundary wetting differs from wetting of free surfaces (Clarke and Gee 1992). Penetration of a dense, polyciystalline material must be accompanied by an increase in volume. If the original grain boundary material can not be dissolved in the penetrating liquid, high elastic strains will build up due to the increased grain boundary volume. The resulting stresses will retard further penetration, but may eventually cause cracking of the solid body and consequently fast penetration. Also the application of external stresses may change the rate of, or even reverse, liquid penetration by wetting. If on the other hand the penetration is accompanied by dissolution of the matrix phase along the grain boundaries, no stresses will build up. The penetration velocity will however be limited by the solubility, the dissolution rate and the diffusion rate of matrix material in the liquid. Chapter 2: Theoretical background - Review of previous work 33
Diffusion of aluminium Diffusion of aluminium in TiB2 is another possible penetration mechanism, and was suggested by Baumgartner (1984). As bulk diffusion coefficients are normally orders of magnitude smaller than grain boundary diffusion coefficients, diffusion through the bulk TiB2 is not expected to be of any significance. This is in agreement with observations of penetrated samples, if aluminium is found it is always located to the grain boundaries and grain junctions. Unfortunately no diffusion coefficients for A1 in TiB2 are available, otherwise the penetration rate due to diffusion could be calculated. As for the wetting mechanism, the penetration rate by diffusion is expected to depend on grain boundary segregations as this will change the diffusion coefficient. Diffusional penetration would be most important at high temperatures and in fine grained materials.
2.4.3 Concluding remarks
All investigations seem to agree that TiB2 cathodes are well wetted by aluminium and thus represent a good choice for a wetted, inert cathode material. The biggest technological problem during electrolysis is premature failure of the cell due to cracking or spalling of the cathode. Cracking caused by thermal shocks can be avoided by careful preheating and cooling procedures. A more difficult problem to handle is the cracking caused by gradual penetration of aluminium into the cathode. It is generally agreed that microstructural features control this degradation. A high porosity level accelerates penetration as an open network of pores will then exist. The grain boundaries and grain junctions are other weak points in terms of penetration, TEM studies have shown aluminium to penetrate these leaving small particles or films of aluminium. Finally the presence of secondary phases in the TiB2 materials is frequently pointed to as a problem, reactions between some these and the penetrating aluminium have been identified by TEM.
Uncertainties arise when the detailed mechanism(s) of aluminium penetration and cathode degradation are asked. Differences in microstructure between the investigated materials opens up for different mechanisms to dominate in each case. Stresses set up either by an increase in grain boundary volume upon penetration or a volume increase of the secondary phases in grain junctions upon reaction with the penetrating 34 Development of microstructure in titanium diboride ceramics aluminium have been suggested to cause cracking. Reduction of fracture toughness due to liquid metal embrittlement or reduction of tensile strength due to a liquid aluminium phase at the grain boundaries have both been suggested to accelerate cracking. Another proposed mechanism of degradation, not involving cracking, is the fast dissolution of the cathode caused by pull-out of whole grains of TiB2. This mechanism has been linked to intergranular penetration of aluminium and to dissolution of the grain boundary material in aluminium or cryolite. No conclusive evidence for any one mechanism has been given.
It seems to be generally agreed that the penetration and degradation of TiB2 cathodes during electrolysis depend strongly on the microstructure. Yet a detailed microstructural characterisation is missing in most experiments, and cathode degradation mechanisms are often proposed without much support from such characterisation. Much focus has been put on the making, mechanical testing and aluminium exposure of yet new materials, and on practical aspects such as cell design and pilot scale tests. Further progress depends on a better understanding of which basic mechanisms of penetration and degradation are active during electrolysis, and calls for detailed microstructural investigations of the materials both prior to and after exposure to aluminium. The high spatial resolution obtained in modem transmission electron microscopes makes studies of structure and chemistry at Angstrpm (10'10 m) and nanometer (10"9 m) level possible. This allows the penetration mechanism of aluminium and the possible reactions between aluminium and the intergranular secondary phases to be studied in detail for the first time. 35
3.1 Materials selection and preparation
The powders All samples have been made from the same batch of carbothermically produced TiB2- powder, grade A, from the German supplier H.C.Stark. The chemical composition of this powder as determined by ICP-AES-analysis (Ion Coupled Plasma-Atomic Emission Spectroscopy) and Leco combustion analysis is given in table 1.1 together with the suppliers specification.
Table 3.1 Powder composition as measured at SINTEF Molab (upper 080695, middle 110893) and producers specification for the powder (lower).
Ti B B/Ti C N O Fe A1 Si Cr Zr Mo W wt% wt% mole wt% wt% wt% wt% wt% wt% wt% wt% wt% wt% 67.0 30.4 2.01 0.03 0.09 0.95 0.25 0.02 <0.1 0.02 0.36 0.52 0.07 67.2 30.0 1.98 0.06 0.06 0.57 0.18 <0.1 0.36 0.07 >66.5 >28.5 (1.90) <0.25 <0.25 <2.0 <0.25 other metallic impurities max. 0.1 wt%
In some samples small amounts of elemental titanium- and/or boron-powders were accurately weighed in and mixed with the TiB2-powder prior to hot pressing. The boron powder were from Merck (art. 12070), had a grain size below 10 pm and contained more than 96 % boron. It was mixed into the TiB2 powder by grinding in an agate (Si02) mortar. The titanium powder contained 99.7 % titanium (and less than 0.01 % iron) and was mixed with the TiB2 in a Turbula rotary mixer for two hours.
Hot pressing procedure Uniaxial hot pressing was performed in a hot press from Thermal technology GmbH, mod. HP50-70106. Some 3-4 grams of powder was placed in a graphite pipe with inner diameter of 10 mm and closed with graphite pistons. After assembling the sample in the hot press, the vacuum oven was evacuated to <0.5 mbar and refilled with 99.998 % argon gas three times. During sintering a continuous gas flow of argon was 36 Development of microstructure in titanium diboride ceramics let into the oven to provide a slight overpressure towards the surrounding atmosphere. The graphite heating element of the oven provided a reducing sintering atmosphere. The temperature was raised slowly to the sintering temperature of 1790, 1880 or 1960 °C in 2.5 hours, including a 15 minutes hold at 1200 °C. The sample was kept at this temperature for two hours, and then cooled slowly to room temperature in 3 hours. At the maximum temperature the pressure was kept at 50 MPa (420 kg). During sintering the temperature was continuously monitored with an optical pyrometer and recorded on a computer along with the pressing force and piston position. Temperature, pressing force and piston position data for two of the runs are included in appendix 1. The final shape of the samples after sintering was cylindric with a height of approximately 10 mm.
Aluminium exposure The hot pressed samples were exposed to liquid aluminium at a temperature of 980°C for 24 hours. Each sample were placed in an alumina crucible containing about 40 grams of pure aluminium. To reduce the thermal shock the samples were preheated prior to the immersion in aluminium.
Specimen preparation for microstructural investigations Thin slices (0.5-1 mm) were cut from the hot pressed samples with a water cooled diamond saw. Specimens for optical microscopy (OM) and scanning electron microscopy (SEM) were made by grinding and polishing the cut sections to obtain shiny, mirror-like surfaces. Grinding started on a cloth sprayed with a suspension of 45 pm diamond particles and proceeded in steps with finer particle sizes. After a final polishing step with 3 pm diamonds the final, mirror-like surface seen in the micrographs was obtained. A few samples were also etched in a 10:1 mixture of HC1 and HN03 to reveal grain boundaries and microcracks during OM (figure 4.2.c).
Transmission electron microscopy (TEM) specimens were prepared from the cut slices. The same grinding procedure as described above, with the exception of the final polishing step, was used until a thickness of approximately 100 pm had been reached. Discs with a diameter of 3 mm were cut with an ultrasonic disc cutter (Gatan, Model 601). Then a dimple grinder (Gatan, Model 656) was used to make a dimple in the centre part of the disc, giving a final thickness off approximately 20 pm. Chapter 3: Experimental 37
The final thinning step has been performed by argon ion-milling. Two different mills have been used, a standard mill (Gatan, Model 600 Dual Ion Mill) and a low-angle mill (Baltec RES 010) milling from one side only. Typical milling parameters for the Gatan mill were an incidence angle of the ions of 12°, an acceleration voltage of 4.5 kV and a total probe current of 1.0 mA. Corresponding parameters used with the Baltec mill were 6°, 6 kV and 4 mA until perforation and then a shorter period at 6°, 2 kV and 2 mA for polishing. TiB2 powder was placed on a "holey" or "lacey" carbon film for TEM investigation.
Artefact from TEM specimen preparation? Figure 3.1 shows a boron nitride particle and an amorphous oxide phase at the rim of a pore in a sintered TiB 2 specimen. Similar amorphous phases have been found in pores in all the specimens investigated except from the less porous samples sintered with titanium addition. Chemical analysis by EDS and PEELS show oxygen, titanium and boron to be important elements in all these amorphous particles (Pettersen et al. 1995). In addition other elements like argon, carbon, molybdenum, iron, chromium, nickel and aluminium have been detected in varying, but generally smaller, amounts in the different specimens.
If the amorphous oxide phase had been a real feature of the sintered samples, it would be an important observation, and should be taken into consideration when discussing the sintering mechanism of TiB2. It has however not been included amongst the results and discussions in the rest of this thesis as it is believed to be an artefact from the ion milling. Redeposition of material sputtered from the specimen itself or from the specimen holder on the back surface is known to take place during single-sided ion milling (Fisher and Angelini, 1985, Bentley 1992). Double-sided milling eliminates the surface contamination. However, the pores in a porous material can still be contaminated by redeposits that are not removed by the low-angle ion milling. This is what is believed to be the origin of the amorphous oxide phases described above. In the thinnest parts of the specimen, where the specimen itself provides no shadowing, all the pores are free from the amorphous oxide. The chemistry of the amorphous phase also strongly suggests it to be a sputtering artefact. Boron and titanium are the major elements of the sample itself and are thus always present, in the exposed samples traces of aluminium were also found. Some oxygen is available in the 38 Development of micro structure in titanium diboride ceramics vacuum chamber of the ion mill, additional oxidation can take place once the specimen is removed from the mill. In the Gatan mill a stainless steel cathode and a molybdenum specimen holder were used. Sputtering from these explains the presence of metallic elements in specimens prepared in the Gatan mill. These elements were replaced by carbon in the amorphous oxides of the specimens prepared in the Baltec mill. This is in agreement with the redeposition mechanism as the specimen holder in this mill was made from graphite. Finally the same amorphous oxide phase was found alongside metallic aluminium in the pores of samples that had been exposed to liquid aluminium. This suggests that the amorphous oxide was formed upon ion milling rather than during sintering as it would be expected to react to form aluminium oxide in the presence of liquid aluminium.
Figure 3.1 TEM-micrograph from ion-milled TiB2 specimen. The dark, amorphous phase at the rim of the pore is believed to be an artefact from the ion milling.
Attempts to prepare TEM-specimens without the ion milling step to test the validity of the redeposition mechanism have not been successful.
3.2 Microstructural characterisation
Density measurements The hot pressed samples were weighed, and their height and diameter accurately measured. The density was then calculated from the formula p = m/(7td 2h/4). The accuracy of the calculated values are within ± 0.2 %. Chapter 3: Experimental 39
Optical microscopy (OM) An optical microscope (Reichert MeF3) was used for measurements of grain sizes and low magnification microstructural characterisation. Polarised light was used to reveal the grain structure, see figure 4.2.a. The pores show up as bright spots when using polarised light and dark spots without polarisation, figure 4.2.b. The average number of grains intersected by a straightline was used to calculate the average grain size.
Scanning electron microscopy(SEM) SEM of the TiB2 powder was performed with a JEOL JSM 6400 microscope. The powder was placed on a conductive tape during imaging. The aluminium exposed samples were studied in a JEOL JSM 840 microscope. Elemental mapping of aluminium x-ray signals was carried out with a LINK EDS system attached to this microscope.
Transmission electron microscopy (TEM) TEM was the main characterisation tool in this study and different microscopes were used to explore different features of the micro- and nanostructure.
General analytic electron microscopy (AEM) including imaging of grain structures and identification of phases by electron diffraction methods and element sensitive spectroscopic methods was performed with two different microscopes. These were a Philips CM30 microscope operating at 300 kV acceleration voltage and a JEOL JSM 2000 FX 200 kV microscope. The Philips microscope was equipped with an energy dispersive x-ray spectrometer (EDS) from ED AX and a parallel electron energy loss spectrometer (EELS) from Gatan for chemical analysis. To avoid carbon contamination the specimen was kept at a temperature of -170 °C during EELS- acquisitions. The JEOL microscope was equipped with an ultra-thin window EDS detector from NORAN for detection of light elements (Z>5). The elemental distribution in some exposed specimens was also studied by electron spectroscopic imaging (ESI) in a Zeiss EM 912 microscope operating at 120 kV with an fi in column energy filter. 40 Development of microstructure in titanium diboride ceramics
High resolution electron microscopy (HREM) imaging was performed with a JEOL JSM 4000 EX microscope operating at 400 kV. The point resolution of this microscope is 1.7 A and the tilting range ±20°. In addition some HREM imaging was done with the Philips CM30 microscope which has a point resolution of 2.3 A. A VG HB 501 STEM (scanning TEM) was used for spatially resolved elemental analysis from grain boundary areas. The field emission gun gave an electron current of 1 nA in a 1 nm largeprobe (FWHM), and EDS spectra were recorded with a NORAN system.
The grain boundary misorientations were determined from Kikuchi patterns obtained with the Philips CM30 microscope. The patterns were digitised with a CCD camera and analysed with CHANNEL 2.5 software after adjusting the software parameters for use with TEM. More details on this method and examples of its applications can be found in Hpier et al. (1994) and Bardal et al. (1995). 41
This chapter first introduces the TiB2 powder which has been used and summarises the most important results from the hot-pressing experiments. The last two sections give a detailed description of the microstructure of the different hot-pressed materials, ranging all the way from the millimetre level as seen by optical microscopy to atomic resolution analysis in the transmission electron microscope. First the microstructure that arises after hot-pressing is presented, then the changes arising from 24 hours of exposure to liquid aluminium.
4.1 The TIB? powder
The shape, size and phase distributions of the powder were briefly investigated by SEM and TEM. SEM shows the particle size distribution to be broad. Particles range in size from fine debris smaller than 1 pm in "diameter" to the largest particles of 10- 20 pm, the most typical size being 2-3 pm as seen in figure 4.1.a. The powder particles are seen to have an irregular, angular shape after the milling. TEM shows each powder particle to be a single crystal with no grain boundaries inside. Along the comers there is a high dislocation density from the milling process, figure 4.1.b. Besides the TiB2 phase TEM shows the powder to contain a minor amount of BN- particles, figure 4.1.c. These were identified by EELS, further details of the BN phase are given in chapter 4.3.1. No other phases were found. 42 Development of microstructure in titanium diboride ceramics
Figure 4.1 Size and shape of TiB2 powder particles as imaged by (a) SEM and (b) TEM. (c) TEM micrographof BN-particles present in the powder.
Hot pressing was used to obtain dense materials with low amounts of secondary phases. Various temperatures and amounts of sintering additives were tried in search for the optimum combination of density and purity. A summary of hot pressing Chapter 4: Results 43
conditions and some key results from eight of these hot pressing trials are given in table 4.1. The microstructure of these samples is presented in detail in the following chapters.
Table 4.1 Results from eight different hot pressing trials. The theoretical densities for boron-containingmaterials arecalculated in chapter 5.2.2.
Sample Hot pressing Density Theoretical Relative Average* Aluminium composition temperature (-W- 1%) density-TD density grain size penetration (wt%) (°Q (g/cm3) (g/cm3) (%TD) (pm) (SEM) TiB2 1790 4.04 4.52 90 9.6 Yes TiB2 1880 4.21 4.52 93 11 Yes TiB2 1960 4.33 4.52 96 21 Yes TiB2 + 2.0Ti 1790 4.46 4.52 99 12-19** Yes TiB2 + 2.0Ti 1880 4.46 4.52 99 17-21** Yes TiB2 + 2.0B 1880 4.28 4.44 96 10 No TiB2 + 2.0B+ 2.0Ti 1880 4.28 4.44 96 11 No TiB2 + 0.9B+ 2.0Ti 1880 4.13 4.48 92 11 No
* The largest grain sizes found were typically twice the average value. ** Inhomogeneous structure, see figure 4.12. The lower value is for the central regions, while the larger value is measured in the outermost regions.
Based on the microstructural investigations it is found convenient to divide the samples into three different groups in the subsequent chapters. The members of each group show similarities in microstructure and resistance to aluminium penetration, and are discussed together. Optical microscopy micrographs showing typical features of each group are given in figures 4.2,4.8 and 4.12 respectively. The first group consists of the three samples with no boron or titanium additions. These are hereafter referred to as pure TiB2, although it might be argued that there are some impurities present according to the chemical analysis in table 3.1. The second group consists of the three samples with only boron or boron and titanium additions, and the third group are the two samples with only titanium addition. 44 Development of microstructure in titanium diboride ceramics
4.3 Microstructure of hot-pressed samples
4.3.1 Pure TiB2
In accordance with table 4.1 the OM investigations of the TiB2 samples sintered without additives show the grain size to increase and the porosity level to decrease when the sintering temperature is raised. The micrographs in figure 4.2 also show that no secondary phase particles large enough to be resolved by optical microscopy are present in the pure TiB2 samples. The sample sintered at the highest temperature is found to contain a substantial amount of transgranular microcracks and pores that are located inside the grains. In the two other samples most pores are located in the grain junctions and grain boundaries, and no microcracking is observed.
(continued next page) Chapter 4: Results 45
Figure 4.2 Microstructure as seen by optical microscopy of pure TiB2 samples hot pressed at various temperatures. a) Polarised light. Left: 1790°C, middle: 1880°C and right: 1960°C. b) Unpolarised light, 1880°C. Only pores visible. c) Etched, unpolarised light, 1960°C.Notice transgranular microcracking.
The OM investigations showed no differences in microstructure throughout one sample. Thus the TEM-specimens that were all taken from 1/2-1 mm below the outer surface of the hot-pressed samples do represent the bulk material well. As TEM- investigations showed all three samples to have a similar" microstructure, only the results from the intermediate sintering temperature, 1880 °C, are reported in detail. The larger-grained sample sintered at 1960 °C in addition show extensive intergranular and transgranular microcracking (figure 4.3.b) as expected from the OM investigations. 46 Development of microstructure in titanium diboride ceramics
Figure 4.3 a) Microstructure of pure TiB2 as seen in TEM. b) Transgranular (left) and intergranular (right) microcracks in large grained sample sintered at 1960 °C.
As expected from the OM-investigations no large amounts of secondary phases are present. The major impurity phase is found to be boron nitride, BN. The BN-particles are usually located in grain junctions or at grain boundaries, see figure 4.4.a. In some cases they are also trapped inside large TiB2-grains, this is most frequently seen in the 1960°C-sample (Pettersen et al. 1995b). The BN particles are shaped as hexagonal Chapter 4: Results 47
plates. When tilted edge on (i.e. with the basal plane parallel to the beam direction) they appear rod-like with typical dimensions 0.2 pm times 1 pm, and are easily recognised from the strong diffraction contrast with high local strains and cracks as seen in figure 4.4.a. Microdiffraction patterns from systematic tilting series proved the crystal structure to be that of hexagonal BN with lattice parameters a = 2.5 A and c = 6.7 A. A [21T 0]-zone axis pattern is shown in figure 4.4.b.
The EELS-spectrum from the BN particles, figure 4.4.d, show boron and nitrogen to be the only elements present. The ELNES (Energy Loss Near Edge Structure) part of the boron edge contains four sharp, separated peaks which distinguish this phase from the other boron-containing phases found in these investigations. Along with the small plasmon peak at 7.5 eV in the low-loss spectrum, 4.4.c, this can be used as a ‘fingerprint ’ for a fast identification of BN. Quantification of the EELS-spectra gives a B/N-ratio of 1.0 ± 0.15. A marked orientational dependence of the ELNES of the boron and nitrogen edge is seen when comparing spectra acquired with the electron beam parallel to the c-axis and normal to it as demonstrated in figure 4.4.d and e. This dependence has earlier been discussed in detail by Leapman et al. (1983). Except from a very few occurrences of boron carbide no secondary phases besides BN were found in the pure TiB2 samples. The characteristics of the boron carbide phase is described in detail in the next section.
To look for possible nanometer sized phases, crystalline or amorphous, the grain boundaries of the pure TiB2 material were studied by HREM. Grain boundaries were chosen at random and tilted to make the grain boundary plane parallel to the electron beam. This step is important to avoid false interpretations of the images as illustrated in figure 4.5. The width of a grain boundary film will be underestimated, or the film may no longer be visible at all, if the grain boundary plane is not properly aligned within a few degrees from the beam direction. Finally, while keeping the boundary parallel to the electron beam, the specimen is tilted to allow simultaneous lattice imaging in both the grains that make up the grain boundary. This was done looking at the kikuchi bands from the two grains simultaneously. The small unit cell of TiB2 allows only the (0001), (1100) and (llOl)-planes to be resolved without difficulty, and thus severely restricts which crystal orientations can be used for lattice imaging. These strict conditions on the orientation of both the grain boundary and the two 48 Development of microstructure in titanium diboride ceramics
so 100 Energy Loss (eV) Energy Loss (eV)
Figure 4.4 BN particles in pure TiB2 sample, a) Located in a grain junction (left), at a grain boundary (middle) and inside a large TiB2-grain (right). b) Microdiffraction pattern from  zone axis of hexagonal BN. c) EELS low-loss spectrum from BN. d-e) EELS-spectra from BN showing the orientation dependence of the fine structure. Electron beam parallel to  and normal to . Figure 4.5 Amorphous phase between two grains in polycrystalline, pure TiB2 as revealed by lattice imaging. The phase is believed to have formed by oxidation of a microcrack, a) Imaged edge on. b) Same position, but tilted approximately 5°. Notice the apparent disappearance of the film in the thicker (upper) part in the tilted image. 50 Development of microstructure in titanium diboride ceramics crystals, along with the limited tilting range of the specimen holder, imply that a large fraction of grain boundaries cannot be imaged satisfactorily. This makes the HREM investigations time consuming.
A total of approximately thirty grain boundaries in the pure TiB2 material were investigated by lattice imaging. Of these 90% were found to be without any grain boundary phase as shown in figure 4.6. The lattice planes from the two grains join at the grain boundary, and leave no room for any secondary phases. The remaining, few boundaries contain a thin film as shown in figure 4.5. The lack of diffraction contrast from this phase tells that it is amorphous. The authenticity of the film as a real feature of sintered TiB2 is doubted. The amorphous film is rather believed to have formed from a clean boundary after the sintering process by a mechanism of intergranular cracking and oxidation of the cracked surfaces. This is in agreement with the similarity in appearance between the surface oxide and the film as seen in figure 4.5. Also the EDS analysis from the film showed no difference in chemistry from the pure boundaries. With the setup used for the EDS measurements, oxygen could not be detected due to absorption in the detector window.
Figure 4.6 HREM micrograph from grain boundary in pure TiB2. Viewed edge on, no grain boundary phase present. Chapter 4: Results 51
Some of the grain boundaries that had been studied with HREM were later examined in a VG HB 501 STEM. EDS-spectra show a segregation of iron to both the clean grain boundaries and the ones with a thin film as described above. Scanning a 1 by 20 nanometer large area parallel to the grain boundary gave an iron content of four weight percent, whereas only one weight percent was recorded in the bulk TiB2 away from the grain boundary. Figure 4.7 shows EDS-spectra from the grain boundary imaged in figure 4.6 and from the TiB2 grain some nanometres away from the boundary. The true iron level is somewhat lower than given above as a background signal from the VG microscope containing small iron and copper peaks should be subtracted from both spectra. No significant differences in chemistry between the clean grain boundaries and the one with a film were found using EDS. Rapid build-up of carbon contamination during radiation made EELS analysis of light elements impossible. Using a conventional TEM and a larger electron probe (FWHM of ~20 nm), the grain boundary segregation of iron is hardly visible. Only a minor increase can be detected due to the strongly localised segregation and the large spot size. EELS with this spot size could not detect any differences between the grain boundaries and the bulk TiB2. In addition to the iron small amounts of zirconium and molybdenum were found uniformly distributed throughout the grains and grain boundaries, see figure 4.27.
Figure 4.7 EDS-spectra showing segregationof iron to the grain boundary imaged in figure 4.6. a) 1 nanometer width at the grain boundary, b) Same, but some nanometres away from the grain boundary. 52 Development of microstructure in titanium diboride ceramics
4.3.2 TiB2 with boron addition
Figure 4.8 shows the microstructure of the TiB2 + 2.0B sample, as seen in an optical microscope. The other two samples in the group looked very much the same. Compared to the pure TiB2 sample sintered at the same temperature, figure 4.2, the grain size and shape is seen to be the same, the porosity is lowered, and a new, grey phase has showed up. This new phase is located in grain junctions and at some grain boundaries.
Only the sample hot pressed with 2 wt% boron addition has been examined by TEM as the OM and SEM studies showed no differences in microstructure or aluminium penetration for the two related samples. The TEM-specimens were made from 1/2-1 mm below the outer surface of the sample, thus representing the bulk of this material as no inhomogenities were seen in the OM investigations.
TEM shows the only major difference between this material and the pure TiB2 to be the large amounts of boron carbide, B4C, found after boron addition, see figure 4.9.a. This phase is identical to the grey phase from the OM-investigations. The B4C-phase is found in many grain junctions and has a rounded shape. A size of Vi - 2 pm is the most usual, although some larger B4C-particles, up to 5-10 pm, have been found both by OM and TEM. The stacking faults seen in figure 4.9.b is a characteristic of the phase, the twinning plane being in most cases a (llOl)-plane, see figure 4.9.d. Selected Area Electron Diffraction (SAED) patterns from systematic tilting series has shown the crystal structure to be trigonal with lattice parameters a = 5.6 A and c = 12 A, in agreement with JCPDS (35-798). The EELS-spectrum in figure 4.9.f shows only boron and carbon to be present. Only weak EDS signals from the surroundings are detected due to the absence of heavier elements. The B4C-phase has a large solid solution range allowing C/B-ratios from 0.10 to 0.32 according to the binary phase diagram, and is sometimes referred to as B i3C2. Quantification of the EELS-spectrum in figure 4.9.f gives a C/B-ratio of 0.25 ± 0.06. The high uncertainty is due to the closeness of the two edges. Background modelling for the carbon-edge is made difficult by the EXELFS (extended energy-loss fine structure) from the boron-edge, a reference sample of known composition is needed for higher accuracy. As in the pure Chapter 4: Results 53
Figure 4.8 OM micrographs from TiB2 sample hot pressed with 2 wt% boron addition at 1880 °C. (a) Polarised light, (b) Unpolarised light. Notice the grey B4C phase. 54 Development of microstructure in titanium diboride ceramics
300 400 Energy Loss (eV) Energy Loss (eV)
Figure 4.9 The B4C-phase in TiB2 hot pressed with boron addition. a) Low magnification micrograph showing B4C particles located in grain junctions, at a grain boundary and inside a larger TiB2 grain. b) The stacking faults are a characteristic of the B4C phase. c) SAED pattern from the  zone axis of B4C. d) SAED pattern from [2 T T 0] zone axis of B4C with stacking faults, e-f) EELS-spectra from the B4C-phase. Chapter 4: Results 55
TiB2 sample some BN-particles were found in grain junctions and grain boundaries, sometimes together with B4C as seen in figure 4.9.b. The B4C-phase is far more abundant, however.
HREM from fifteen randomly chosen grain boundaries showed that nor in this material were secondary phases present at the grain boundaries. A lattice image from a clean boundary is shown in figure 4.10. Only one grain boundary had an amorphous film, but this boundary was believed to be microcracked as explained in 4.2.1. HREM from the TiB2 - B4C interface shows this too to be clean as seen in figure 4.1 l.a. The HREM micrograph in figure 4.1 l.b shows the B4C-phase with a high density of stacking faults.
Figure 4.10 HREM micrograph of clean grain boundary in TiB2 hot pressed with boron addition. Viewed edge on.
The chemistry of the grain boundary area prior to aluminium exposure was investigated in a conventional TEM with a probe size of ~20 nm. The result of this was similar to that of the pure TiB2 sample above. A minor increase in iron concentration at the grain boundary, some 0.2 weight percent higher than in the bulk TiB2 100 nm away from the grain boundary, was detected with EDS. EELS did not show any detectable differences at this level. 56
=T- Chapter 4: Results 57
4.3.3 TIB2 with titanium addition
While the microstructure of the two other groups was found to be homogeneous throughout the whole cross-section of the samples, the titanium addition has resulted in an inhomogeneous microstructure. From the low-magnification OM micrograph, figure 4.12.a, three markedly different regions are found as one moves from the surface to the centre of the sample. The outermost 0.5 mm has a relatively high porosity as compared to the inner regions, and the largest grain sizes are found here. Most of the pores are trapped inside the large TiB2 grains. No secondary phases large enough to be detected by optical microscopy are found in the outermost region. In between the outer layer and the central area an intermediate 0.5-1.0 mm broad region is found. Here areas with a mixture of small TiB2 grains, 2-5pm in diameter, and a secondary phase is found along with other areas of more average sized TiB2 grains, see figure 4.12.b. The porosity in the intermediate region is low.
The central region of the TiB2 + 2.0Ti sample reveals yet another microstructure. This innermost and largest part of the sample has a radius of 3-4 mm. The porosity is low, and the grain size is smaller than in the outermost region, although still somewhat larger than in the corresponding samples sintered without titanium addition. The TiB2 grains have a faceted appearance with straight grain boundary lines as seen in figure 4.12.b. A black secondary phase is found in the grain junction, sometimes extending into the grain boundaries as needle-shaped particles.
As the OM-investigations of these samples revealed an inhomogeneous microstructure, care was taken to prepare TEM-specimens from well defined regions of the sample. The outermost region has not been investigated by TEM.
Central region: One TEM-specimen was made from the central region, 3.5 mm below the outer surface. In this specimen several new oxide phases were identified. These phases fill the grain junctions and in some cases extend into the grain boundaries, their shape and size show them to make up the black intergranular phase found during OM investigation. Most grain junctions contain not only one phase, but a mixture of different oxides. An example is given in figure 4.13.a. In addition to the oxides some 58 Development of microstructure in titanium diboride ceramics smaller iron-containing particles are found in three grain channels. No BN or B4C were found in the materials hot pressed with titanium addition. A detailed description of all the phases in the central region is given below.
Figure 4.12 OM micrographs from TiB2 hot pressed with 2 wt% titanium addition at 1790°C. a) Low magnification, polarised light. The three regions with different microstructures are shown, b) Details from the intermediate (left) and central (right) regions. Chapter 4: Results 59
Trigonal Ti203, as seen in figure 4.13, is the most frequently occurring phase in the central region of the samples sintered with titanium addition. The size of the Ti203 particles are typically 1 pm. Systematic SAED tilting series show the crystal structure to be trigonal (Al203-like) with lattice parameters a = 5.1 A and c = 13.7 A, in agreement with the structure reported in JCPDS (10-63). A [21 T 0] zone axis pattern is included in figure 4.13.b. EDS shows only titanium and oxygen to be present, figure 4.13.C. The EELS spectrum in figure 4.13.d verifies this, and quantification gives an atomic ratio Ti/O = 0.66 ± 0.10, in perfect agreement with the stoichiometric 2/3.
Ti(ZrAl)02 is the second most frequent of the oxide phases found, and is always found along with a pure titanium oxide phase, see e.g. figure 4.13, 4.14 and 4.16. The Ti(ZrAl)02-phase seem to be the one of the oxide phases that has the lowest interfacial tension towards TiB2 as it is usually this phase that is found in the comers of the grain junctions and sometimes extending into the grain boundary as seen in figure 4.14.a. SAED-pattems obtained from systematic tilting series, figure 4.14.b, show the crystal structure to be base centred cubic (bcc) with lattice parameter a = 10.2 ± 0.3 A. The EDS- and EELS-spectra in figure 4.14.C shows titanium, zirconium and aluminium to be present along with oxygen. Quantification of the EDS-spectrum gives the ratio 80±3:7±1:12±2 (Ti:Zr:Al) between the metallic elements, whereas a Ti/O-ratio of 0.40±0.06 was found from the EELS-spectrum. No binary, ternary or quaternary phase containing oxygen and titanium and with a similar crystal structure could be found in the JCPDS files. 60 Development of microstructure in titanium diboride ceramics
Energy Loss (eV) Energy Loss (eV)
Figure 4.13 Grain junction with a mixture of oxides as found in the central region of the sample hot pressed with titanium addition at 1790 °C. a) TEM bright field. T = Ti203, Z = Zr(Ti)G2, TZA = Ti(ZrAl)G2. b) SAED pattern from [21 10] zone axis of the trigonal Ti203 phase. c) EDS spectrum from the trigonal Ti203 phase. d) EELS spectra from the trigonal Ti203 phase. Chapter 4: Results 61
Al Zr JU) 5.
100 150 400 SOI Energy Loss (eV) Energy Loss (eV) Figure 4.14 From the central region of the sample hot pressed with titanium addition. a) Ti(ZrAl)02 phase between two TiB2 grains. b) SAED patterns from systematic tilting series of the Ti(ZrAl)02 phase. c) EDS spectrum from the Ti(ZrAl)G2 phase. d) EELS spectra from the Ti(ZrAl)Q2 phase. 62 Development of microstructure in titanium diboride ceramics
A Zr(Ti)02 phase is frequently found as small particles embedded in the larger trigonal Ti203-particles, and is believed to be a variant of cubic zirconia, Zr02 (JCPDS 27-997). The Zr(Ti)02 are typically 0.1 pm in size, an example is given in figure 4.15. They show a strong bright field contrast from defects and/or polycrystallinity, due to this the crystal structure of this phase could not be determined completely. Figure 4.15.b indicates there to be at least one fourfold axis with d-values of 5.1 A. The strongest reflections are in perfect agreement with the fee structure with a = 5.1 A of cubic zirconia. The dissolution of some titanium in the zirconia lattice may be responsible for the reduced symmetry of the phase making also the (OOl)-type reflections visible. Chemical analysis by EDS and EELS shows zirconium, titanium and oxygen to be the only elements present. Quantification of the EDS-spectrum, figure 4.15.c, gives a Zr/Ti-ratio of 4.0 ± 0.3, but parts of the titanium signal could be from the surrounding Ti203-phase. The EELS-spectrum, figure 4.15.d, closely resembles that of Zr02 with the addition of a weak Ti-edge.
Cubic Ti203 was only found once, see figure 4.16. It is surrounded by the Ti(ZrAl)02-phase, and the interface towards this is faceted. The crystal structure was determined by systematic tilting to be bcc with a lattice parameter a = 9.5+0.2 A, see figure 4.16.b. The stronger reflections alone are consistent with a face centred cubic (fee) structure with half the lattice parameter. A close relation to the surrounding Ti(ZrAl)02-phase is indicated by the faceting and an orientation relationship. The  zone axis of the cubic Ti203 phase (figure 4.16.b) is parallel to the  zone axis in Ti(ZrAl)02, and simultaneously the strong (222)-reflections in the first coincide with the (231)-reflections in the later. The EDS- and EELS-spectra from this phase match the ones from trigonal Ti203 perfectly, no other elements are present and the stoichiometry is the same. Chapter 4: Results 63
5.00 10.00 15.00
Ep=I3.5 eV 100- Ep=24.5 eV a. 10-
Energy Loss (eV) Energy Loss (eV)
Figure 4.15 From the central region of the sample hot pressed with titanium addition. a) Particles of Zr(Ti)02 (dark) embedded in Ti202. b) Fourfold SAED zone axis pattern from the Zr(Ti)02 phase. c) EDS spectrum from the Zr(Ti)02 phase. d) EELS spectra from the Zr(Ti)02 phase. 64 Development of microstructure in titanium diboride ceramics
Ep=24 eV o 50 -
0 100 1 =5 20- Energy Loss (eV) o- 10-
Energy Loss (eV)
Figure 4.16 From the central region of the sample hot pressed with titanium addition. a) Grain junction with cubic Ti203 (inner) and Ti(ZrAl)02 (outer). b) SAED-pattems from systematic tilting of the cubic Ti2G3 phase. c) EELS-spectra from cubic Ti203 phase. Chapter 4: Results 65
Fe-rich particles like the ones shown in figure 4.17 is the only phase, besides the oxides described above, found in the central part of the sample sintered with titanium addition. They appear as thin bands, typically 10-20 nm wide, in three grain channels as seen in figure 4.17.b, and may form somewhat larger particles when they end up in multi grain junctions as seen in figure 4.17.a. EDS-analysis indicates that this might be iron, maybe with a small chromium content, but due to the surrounding TiB2 an iron-titanium or iron-boron phase can not be ruled out.
Figure 4.17 Microstructure from the central region of the sample hot pressed with titanium addition, a) Multi grain junction with Fe-rich phase and three oxide phases, b) Three grain channel with Fe-rich phase, c) EDS-spectrum from iron-rich phase. Titanium signal probably from overlapping TiB2. 66 Development of microstructure in titanium diboride ceramics
Intermediate region: Another TEM-specimen from 1.0 mm below the sample surface was made to investigate the intermediate region in figure 4.12. By looking at the same ion-thinned TEM-specimen both by TEM and OM, the secondary phase in this region is proven to be Ti(CNO). Figure 4.18 is from one of the small grained areas with TiB2 and Ti(CNO). As expected from the OM-investigations the fraction of Ti(CNO)-phase is large in this area. The Ti(CNO) particles occupy the grain junctions and some grain boundaries. Their size is typically 1 pm, some larger ones measuring up to 5 pm. SAED-pattems from systematic tilting show the Ti(CNO)-phase to have a fee crystal structure with lattice parameter a = 4.30 ± 0.05 A. Figure 4.18.C shows the  zone axis. The binary phases TiC, TiN and TiO all have the same fcc-structure with lattice parameters in the range 4.19 - 4.33 A. The Ti(CNO)-phase is therefore believed to be a solid solution of these. An orientation relationship to the TiB2 is found where TiB2 II  j^cNO) and (0001)T®2 H (111)t<(cno ) as shown in figure 4.18.d. This gives a semicoherent interphase boundary between Ti(CNO) and TiB2 which is often reflected by dislocation networks and faceting at the interface, see figure 4.18.b. The EELS-spectrum in figure 4.18.e shows only C, N, Ti and O to be present. Quantification of this spectrum gives an atomic ratio of 1:1 between titanium and the sum of the light elements and a 1:2:2 ratio between C:N:0, thus making Ti(C0.2No.40o.4) a proper designation for this mixed phase. Still within the inhomogeneous intermediate region, but outside the small grained clusters, smaller amounts of the Ti(CNO) phase is present in grain junctions along with the oxide phases found in the central region of the specimen. Chapter 4: Results 67
Energy Loss (eV) Energy Loss (eV)
Figure 4.18 The Ti(CNO)-phase in the intermediate region of a TiB2 sample hot pressed with titanium addition. a) Between the grains in a small grained region. b) Faceted interface towards TiB2. c) SAED pattern from  zone axis. d) Orientation relationship between Ti(CNO) and TiB2 neighbour. e) EELS-spectrum from Ti(CNO) phase. 68 Development erf microstructure in titanium diboride ceramics
Lattice imaging shows the grain boundaries both in the central and intermediate region to be without any grain boundary phase prior to aluminium exposure. Ten randomly chosen grain boundaries from each region were successfully imaged in HREM, an example from the central region is given in figure 4.19.
Figure 4.19 Lattice image showing a clean grain boundary in the central region of a TiB2 sample hot pressed with titanium addition. Chapter 4: Results 69
Some of the grain boundaries from the HREM study were later examined in a VG HB 501 STEM. EDS-spectra show a strong segregation of iron to the clean grain boundaries. Scanning a 2 by 20 nanometer large area parallel to the grain boundary gave an iron content of 12 weight percent, far above what was recorded in the samples sintered without titanium addition. Minor amounts of nickel and chromium was found to follow the iron. The next scan was done 5-10 nanometres away from the grain boundary (some drift occurred during acquisition) where the iron concentration had fallen to two weight percent. Finally a background level of only 0.2 weight percent iron was recorded in the bulk TiB2 50 nm away from the grain boundary. Figure 4.20 shows EDS-spectra from the grain boundary imaged in figure 4.19 and from the TiB2 grain 50 nanometres away from the boundary. Even with a much larger probe size, ~20 nm, in a conventional TEM the increase in iron concentration at the grain boundary is clearly visible in the EDS-spectra. The strong and narrow segregation zone of iron detected is roughly equivalent to a single monolayer of iron located at the grain boundary. EELS showed no other difference between the grain boundary and the bulk TiB2 than the uneven iron distribution when using a 20 nm probe.
• Figure 4.20 EDS-spectra from the grain boundary imaged in figure 4.19 showing a strong segregation of iron to the grain boundary, a) At the grain boundary, b) Fifty nanometres away from the grain boundary. 70 Development of microstructure in titanium diboride ceramics
4.4 Microstructure after exposure to liquid aluminium
4.4.1 Pure TiB2
Cross-section specimens from all the three samples hot pressed without sintering additives (pure TiB2) at different temperatures and exposed to liquid aluminium were studied by SEM. They all showed strong penetration of aluminium. Backscattered electron images (BEI) and elemental mapping images using the AIK X-ray signal were recorded from 150 x 150 (xm2 large areas. In addition recordings of the complete X- ray spectrum from selected pores and grain boundaries were performed. Figure 4.21 a and b are BEI and AIK images from the sample hot pressed at 1880°C. Most of the pores are seen to contain aluminium, this is true even at depths of 300 pm below the exposed surface. Point analysis proves that aluminium are present in the pores even several millimetres away from the surface. The three samples hot pressed at different temperatures all showed similar results with regards to aluminium penetration.
Figure 4.21 TiB2 hot pressed at 1880°C and exposed to aluminium at 980°C for 24 hours, a) SEM BEI image, b) SEM-mapping of AIK X-ray signals from the same area. Chapter 4: Results 71
Only the sample hot pressed at 1880°C was investigated by TEM after exposure as SEM revealed complete penetration in all three Al-exposed samples. TEM-specimens were made from depths of 10 and 100 pm below the exposed surface respectively. As expected from the SEM investigations aluminium is found in many grain junctions, an example of which is given in figure 4.22.a. The aluminium is found to fill up most grain junctions completely, forming single crystals typically one or a few microns in size. SAED, EDS and EELS show this phase to be metallic, fee aluminium with a lattice parameter of 4.05 A. A fast identification is possible from the characteristic, sharp plasmon peak at 15 eV in the low-loss region of the EELS-spectrum, figure 4.22.b. Besides the aluminium a few alumina particles (a-Al203) were found in grain junctions and grain boundaries in the outermost specimen. This phase was never found in the specimen taken from 100 pm below the exposed surface, and is thus believed to have formed from surface oxides. The alumina phase is described in detail in chapter 4.4.3 below.
Prior to aluminium exposure boron nitride was the only important impurity phase found in this material. After exposure no BN was found in the outermost sample, but at 100 pm depth both BN and aluminium nitride, AIN, were found. Chemical analysis showed the BN to be contaminated by aluminium in some areas, while in other areas only AIN and no BN were found. The AIN particles were often found to be coexisting with metallic aluminium in grain junctions and positioned in between the aluminium and the TiB2 matrix as seen in figure 4.22.a. The shape and size of the AlN-particles is comparable to that of the BN particles in the unexposed samples. The crystal structure was determined by systematic tilting to be hexagonal with lattice parameters a = 3.15 ± 0.05 A and c = 4.95 ± 0.05 A, in agreement with the structure reported in JCPDS (25-1133). A SAED  zone axis pattern is shown in figure 4.22.C. Chemical analysis show aluminium and nitrogen to be the only elements present, EDS- and EELS-spectra are given in figure 4.22.d and e. 72 Development of microstructure in titanium diboride ceramics
o 600 -
6 400 - Ep=21.5 eV
Energy Loss (eV) Energy Loss (eV)
Figure 4.22 Microstructure of pure TiB2 sample after aluminium exposure. a) Grain junctions with metallic aluminium and two, small AIN particles. b) Low loss EELS-spectrum from the aluminium phase. c) SAED pattern from  zone axis in AIN. d) Ultra-thin window EDS-spectrum from the AIN phase. e) EELS-spectra from the AIN phase. Chapter 4: Results 73
Some of the grain boundaries in the pure TiB2 materials have been penetrated by aluminium during the 24 hours exposure. Although they look unaffected at a first glance, a closer investigation of the grain boundaries at high magnifications reveal changes to have taken place from the clean grain boundaries of the unexposed materials. Three distinctly different appearances of the grain boundary penetration are found at different grain boundaries in the same sample after alu minium exposure. Some contain a thin film of aluminium, some contain small lens-shaped particles of aluminium while in the last group of grain boundaries no aluminium phase is found. Figure 4.23 serves to illustrate this behaviour. The three grain boundaries leading into the common three-grain-channel exhibit each of the three different appearances of penetration behaviour described. All three situations were commonly observed in both the outermost specimen and the specimen from 100 Jim below the exposed surface.
Figure 4.23 Microstructure of pure TiB2 sample after aluminium exposure. a, b) Bright field micrographs from the same three-grain-junction. Taken at slightly different specimen tilt angles to obtain satisfactory contrast for all three grain boundaries, c) Schematic drawing of the microstructure showing three different kinds of grain boundary penetration behaviour: thin film of A1 (1), lenses of A1 (2) and no A1 (3). 74 Development of microstructure in titanium diboride ceramics
The thin films of aluminium found at some boundaries have a fairly constant width along the grain boundary. The width is usually in the 5-30 nm range, in more rare cases up to 100 nm have been measured. The films often extend several micrometers in length, i.e. from one grain junction to another. The HREM micrograph in figure 4.24 show the solidified film to be single-crystalline, and the lattice plane spacing of 2.03 A is equal to that of the aluminium (200) planes. The width of this particular film is found to be 6 nm.
Figure 4.24 HREM micrograph showing a thin, intergranular aluminium film on a grain boundary in a pure TiB2 sample after aluminium exposure. Chapter 4: Results 75
The grain boundaries with lens-shaped aluminium particles are somewhat less frequent than the two other kinds. Figure 4.25.a is from an inclined grain boundary in a rather thick specimen. A rather high density of lenses are seen, the shoal of "fishes" would have been even more dense if all the lenses had been visible at the same specimen tilt. The lenses are typically 30-50 nm long and 10-20 nm wide in this case. In the areas between the lenses the grain boundary is clean. Examples of intermediate stages between the film of equal thickness and the lenses are also found as illustrated in figure 4.25.b. The small size of the films and lenses makes systematic tilting and diffraction inconvenient for phase identification. EDS- and EELS-spectra are easily mixed with signals from the surrounding TiB2 grains, but show a high aluminium content. The EELS low loss spectrum reveals sharp plasmon peaks at 15 eV which proves this to be metallic aluminium. Electron spectroscopic imaging (ESI) showed metallic aluminium to be the only element present in the lenses, figure 4.25 .c and d. The weak signals due to the low inelastic scattering cross-section at the AIK edge (1560 eV) makes the three window technique rather time consuming. However, direct plasmon imaging of the sharp 15 eV Al-plasmon using a 6 eV energy slit gives excellent contrast at very low exposure times. HREM imaging of the lenses also confirmed them to be metallic aluminium.
Finally some grain boundaries reveal no signs of a grain boundary phase even after careful tilting and HREM imaging. Thus these grain boundaries look just like the clean grain boundaries found prior to the aluminium exposure. A HREM micrograph from such a grain boundary is included as figure 4.26. Chemical analysis by EDS using a spot size of 30 nm (FWHM) reveal the presence of a very small amount of aluminium even at these "clean" grain boundaries. In figure 4.27 a spectrum acquired from the grain boundary is compared with one obtained under the same conditions a few hundred nanometres away from the grain boundary. No aluminium signal is detected away from the grain boundary. The spectrum from the grain boundary region shows 0.3 weight percent of aluminium to be present, or roughly equal to half a monolayer of alu min ium at the boundary region. In addition to the aluminium weak signals from iron, zirconium and molybdenum are seen in both spectra. Quantification of the amount of each of these impurities gives results in accordance with the chemical analysis in table 3.1. 76 Development of micro structure in titanium diboride ceramics
Figure 4.25 Lenses of aluminium at grain boundaries in pure TiB2 after aluminium exposure. a) Tilted grain boundary with a high density of lenses. b) Intermediate stage between smooth film and lenses. c) ESI of grain boundary with lenses of aluminium. Three window technique with a slit width of 70 eV at the AIK-edge. Recording time 3x30 seconds. d) As c), but single acquisition only. From the metallic aluminium plasmon peak (15 eV) with a 6 eV slit width. Chapter 4: Results 77
Figure 4.26 HREM micrographfrom a grain boundary in pure TiB2. No aluminium phase is present after 24 hours of alumi nium exposure.
0.70 1.40 2.10 2.80 3.50 4.20 4.90 5.60 6.30
Figure 4.27 Comparison of EDS spectra from a "clean" grain boundary and from the neighbouring TiB2 grain. Pure TiB2 exposed to aluminium for 24 hours. 78 Development of microstructure in titanium diboride ceramics
The results from a study of the relationship between grain boundary misorientation and aluminium penetration are given in table 4.2 and figure 4.28. All the grain boundaries in three different small regions, all half a millimetre apart, were studied. In addition some more randomly chosen ones were analysed. The misorientation was found by indexing Kikuchi-pattems using the Channel software. Based on careful tilting and examination at high magnifications and supported by EDS analysis each grain boundary was classified into one of the categories "clean", "lenses" or "film" depending on the kind of penetration behaviour observed. Some grain boundaries failed to fit into this classification as they showed different behaviour in different parts of the boundary. For instance one end the grain boundary could have a continuous film of aluminium while only discrete lenses or no aluminium at all was found some micrometers away at the other end. In table 4.2 such boundaries are placed in the "mixed" category, whereas in figure 4.28 they are represented by one mark in each category.
Table 4.2 Number of clean grain boundaries and grain boundaries with lenses or films of alumini um in three different areas. Pure TiB% sample exposed to liquid aluminium at 980 °C for 24 hours. TEM-specimen from 100 jam depth.
Area # Clean # Lenses # Film # Mixed E Grain boundaries 1 8 1 1 2 12 2 9 1 7 5 22 3 5 6 21 6 38 Others 6 1 9 1 17 Total 28 9 38 14 89 Chapter 4: Results 79
Aluminium penetration vs. misoMentation
Figure 4.28 Aluminium penetration of grain boundaries as a function of grain boundary misorientation in TiB2 sample after 24 hours aluminium exposure. The film at the "0° grain boundary" is penetration of a transgranular microcrack.
4.4.2 TiB2 with boron addition
Cross sections from the three samples hot pressed with boron additions (table 4.1) were prepared for and investigated by SEM in the same way as the pure TiB2 samples. SEM mapping images failed to reveal any aluminium X-ray signals in any of these samples, an example is given in figure 4.29 for the TiB2 + 2 wt% B sample. As the mapping technique is time consuming, and therefore gives a limited resolution in terms of detectability, additional spot-measurements at selected positions were performed. These confirmed the picture from the mapping as no aluminium signals could be found in any of the investigated pores. This result is opposite to what was found in the samples without boron additions reported above, and shows that less, if any, aluminium has penetrated the samples hot pressed with boron additions. 80 Development of microstructure in titanium diboride ceramics
Figure 4.29 TiB2 with two weight percent boron addition hot pressed at 1880°C and exposed to aluminium at 980 °C for 24 hours, a) SEM BEI image, and b) SEM mapping of AIK X-ray signals from same area.
The sample hot pressed with 2 weight percent boron addition only was investigated by TEM to supplement the SEM observations of the exposed samples. TEM-specimens were made from depths of 10 and 80 pm below the exposed surface respectively. Unlike what was expected from the SEM investigations aluminium was found in both these specimens. The total amount of alu minium that had penetrated this sample was significantly smaller than in the pure TiB2 samples, however. No larger particles of aluminium-rich phases or pores filled with aluminium could be found.
In both specimens some of the grain boundaries were found to contain thin films of metallic aluminium similar to those described above in the pure TiB2 samples. However, the fraction of penetrated grain boundaries was smaller and no lenses of aluminium were found. In the innermost specimen, 80 pm below the exposed surface, most of the grain boundaries did not show any sign of aluminium penetration at all. Unlike what was the case for the pure TiB2 sample EDS-investigation with a 30 nm (FWHM) spot failed to detect traces of aluminium at these grain boundaries.
Both the B4C and BN phase that were present in the unexposed samples are still found intact after aluminium exposure. No reactions with the small amounts of aluminium that have penetrated during the 24 hours of exposure were recorded. Figure 4.30 shows a B4C-particle located at a penetrated grain boundary. The aluminium film follows the TiB2-B4C interface, no aluminium is present inside the B4C particle. Chapter 4: Results 81
Figure 4.30 Grain boundary with aluminium film and a B4C particle in TiB2 sample sintered with boron addition and exposed to liquid aluminium for 24 hours. The dark spots are carbon contamination from EDS recordings.
4.4.3 TiB 2 with titanium addition
Cross section specimens from the samples sintered with titanium additions and exposed to liquid aluminium at 980 °C for 24 hours were prepared for SEM investigations. SEM imaging along with aluminium mapping and X-ray spot analyses revealed penetration of aluminium into the grain junctions all the way into the central parts of the samples. This was the case for both samples irrespective of original hot pressing temperature. The SEM results are similar to those reported in the pure TiB2 case above.
Two TEM-specimens were made from the sample hot pressed at 1760 °C, one from close to the exposed surface (10 |im below) and one somewhat deeper into the sample (100 um). As the outermost two millimetre of the sample had been cut away prior to exposure, the TEM-specimens were from the two innermost regions of the inhomogeneous microstructure described in chapter 4.3.3. The penetration of aluminium after 24 hours of exposure is found by TEM to be much stronger in the samples sintered with titanium addition than what was the case for the other samples 82 Development of microstructure in titanium diboride ceramics reported above. In addition the penetrated grain boundaries have become faceted. Details of the penetration and of reaction products formed from the secondary phases present are reported below. None of the original oxide phases found in the central region of the unexposed samples were present after aluminium exposure.
In both specimens alumina (a-Al203) was found to be the dominating phase in the grain junctions along with some metallic aluminium. Figure 4.31.a gives an example where polycrystalline alumina with grain sizes of approximately two hundred nanometres fills up most of a grain junction. In between the central alumina particles and the surrounding TiB2 metallic aluminium is found. The shape of the grain junctions filled with alumina is similar to that of the oxide phases in the unexposed material. The crystal structure of the alumina phase was determined from SAED- pattems from systematic tilting series to be trigonal with lattice parameters a = 4.8 ± 0.05 A and c = 13.0 ± 0.1 A. This is in agreement with the JCPDS (10-173) data for a-Al203. A  zone axis SAED pattern is given in figure 4.3l.b, the extinction rules for the trigonal phase take away most of the reflections. Ultrathin window EDS and EELS show oxygen and aluminium to be the only elements present, figure 4.3l.c and d.
Figure 4.32 shows a small grained region with a microstructure similar to the intermediate region prior to exposure. The Ti(CNO) phase present in the grain junctions prior to exposure has been replaced by a mixture of three phases; metallic aluminium, alumina and Ti(CNO), as seen from the magnification in figure 4.32.b. The aluminium seem to have started to dissolve the Ti(CNO) phase preferentially along certain crystallographic directions, leaving a network of aluminium and Ti(CNO). Some alumina particles are also found in the same grain junction. Both Ti(CNO) and A1 has a fee crystal structure, and the difference in lattice parameters is only a few percent as seen from the slight splitting of the reflections in the  zone axis SAED pattern in figure 4.32.C. This pattern was made by including both phases inside the diffraction aperture, and shows them to have the same crystallographic orientation after solidification of the aluminium. Chapter 4: Results 83
£25- O 20 - .9 15-
o 10 -
50 100 Energy Loss (eV) Energy Loss (eV) Figure 4.31 Microstructure of TiB2 sample hot pressed with 2 wt% titanium addition and exposed to aluminium for 24 hours at 980 °C. a) Grain junction filled with a-Al203 particles and aluminium b) SAED-pattem from  zone axis in a-Al203. c) Ultrathin window EDS-spectrum from a-Al203-phase. d) EELS-spectra from a-Al2G3. 84 Development of microstructure in titanium diboride ceramics
Figure 4.32 Microstructure in the reaction zone between penetrating aluminium and the Ti(CNO)-phase. From the intermediate region of a TiB2 sample hot pressed with 2 wt% titanium addition and exposed for 24 hours. a) Low magnification overview. b) Details from (a). Bright areas are aluminium, darker areas Ti(CNO). c) SAED pattern from  zone axis showing both metallic aluminium and Ti(CNO) to be present simultaneously and with the same orientation.
While the grain boundaries in the pure TiB2 samples and the ones sintered with a boron addition were either unaffected or covered by a thin film or lenses of aluminium after exposure, the grain boundaries in the samples with titanium addition have suffered a much more violent attack during exposure. Some examples are given in figure 4.33. All the grain boundaries are penetrated by aluminium. The thickness of the aluminium layers varies from a few nanometres up to a few hundred nanometres. Chapter 4: Results 85
In addition to the increased thickness of the aluminium film, most of the grain boundaries have become faceted upon exposure. The faceting follows specific crystal planes in the TiB2 grains, namely the basal plane (0001) and the perpendicular prismatic (1100) plane. This is clearly seen from the HREM micrograph in figure 4.34.
Figure 4.33 Faceted grain boundaries with various levels of faceting and various thicknesses of metallic alumi nium films. TiB2 hot pressed with 2 wt% titanium after 24 hours of aluminium exposure. 86 Development of microstructure in titanium diboride ceramics
Figure 4.34 HREM micrograph from faceted grain boundary in TiB2 sintered with 2 wt% titanium addition and exposed to aluminium for 24 hours. The facets are very sharp, and the faceting planes in TiB2 are the basal plane (0001) and the prismatic plane (1100). 87
As concluded in the literature review a high purity and high density material is considered essential to avoid serious degradation of titanium diboride cathodes during aluminium electrolysis. The previous chapter has given a detailed description of the micro- and nanostructure of various TiB2 materials, including the effect of minor boron and titanium additions, both after hot pressing and after subsequent exposure to liquid aluminium. In the following sections these results are discussed and interpreted with emphasis on explaining the microstructural evolution during sintering and understanding the mechanisms of aluminium penetration and cathode degradation during exposure.
5.1 The TiB? powder
A large part of the impurities in the final hot-pressed materials originate from the carbothermic powder production. In this section the formation of impurities during powder production is discussed with references to thermodynamical calculations, chemical analysis and TEM-observations of microstructure. The chemical analysis of the powder, table 3.1, show oxygen and some metallic elements to be the main impurities along with only trace amounts of carbon and nitrogen. The B/Ti-ratio is not significantly different from 2.0, i.e. close to the boron-rich limit of the narrow compositional range for TiB2. The phase diagram of boron-titanium, figure 2.4, gives the equilibrium range for the TiB2 phase as 1.90 < B/Ti < 2.00 in the temperature range from 500 °C up to -2200 °C. SEM and TEM investigations show the grain size of the powder to be 2-3 pm, and show some BN particles to be present along with the TiB2.
The impurities in the powder must originate either from the raw materials used in the carbothermic reduction, from side-reactions during the reduction or from the final milling process to produce the fine-grained powder. The main, overall reaction of the carbothermic reduction is
Ti02 (s) + B203 (1/g) + 5 C (s) = TiB2 (s) + 5 CO (g) 88 Development of microstructure in titanium diboride ceramics
Additional by-products are possible. The formation of these depends on the reaction temperature, the atmosphere and the relative amount of the reagents. Thermodynamical data and calculations relevant to the carbothermic powder production of TiB2 are included in appendix 2. These cover selected reactions in the system consisting of (B203, Ti02, C, TiB2, CO, BN, B4C, TiC, TiN, N2, 02). The results are summarised in table 5.1. All activities were set equal to 1 during these calculations. Taking into account the reduced partial pressure of CO in the reduction oven will lower the temperature limits for all reactions.
Table 5.1 Possible by-products from the carbothermic reduction process. All activities set equal to 1 during calculations.
B2C>3/Ti02 < 1 B203/Ti02 = 1 B203/Ti02 > 1 C deficit residual oxides C balance, TiC (T>1600K) only TiB2 (T>1600K) B4C (T>1900K) P(N2) = 0 C balance, TiN (1500K
The competition between carbide and nitride formation depends on the partial pressure of nitrogen in the reaction atmosphere. Reducing the nitrogen pressure will favour the formation of carbides at the expense of nitrides. With an excess of Ti02 the carbide (TiC) is the stable product at typical carbothermic reduction temperatures (above 2000°C) when an excess carbon is present. Only in a limited temperature range well below this the nitride (TiN) will be preferred to the carbide provided that the nitrogen pressures is above at least 3'10"2 atm.
Excess B203 results in formation of B4C or BN. B4C is preferred to BN at low nitrogen pressures, the limiting pressure being l'lO"4 atm N2 at 1900 K increasing to 1.3 10"' atm N2 at 2500 K. If the amount of carbon is not sufficient to form B4C free boron will be stable at low nitrogen pressures. The TEM-observation of BN-particles in the powder along with the high B/Ti-ratio from the chemical analysis indicates a Chapter 5: Discussion 89 slight excess of B203 and presence of nitrogen during powder production. The BN is then formed according to the reaction
B203 (1/g) + C (s) + N2 (g) = 2 BN (s) + 3 CO (g)
The volume fraction of BN-particles is calculated to be 0.32 volume percent assuming the whole nitrogen content of 0.09 weight percent to be bound as BN. TEM revealed no C-containing phases in the powder. Along with the very low carbon level from the chemical analysis, this indicates a well balanced carbon level in the carbothermic reduction.
No residual oxide particles were found in the TiB2-powder. The oxygen impurity must therefore be caused only by surface oxidation of the TiB2 upon exposure to air as expected from the thermodynamical calculations. Once a surface oxide layer is formed it will it will act as a barrier and protect the material against further oxidation. Milling increases the oxygen content as a smaller grain size increases the total surface area. As seen from SEM the typical grain size of the powder seem to be 2-3 pm, but a lot of submicron particles are present along with a few large ones (—10 pm). The chemical analysis gives an average oxygen content of 0.75 wt%, corresponding to 2 volume percent of oxides if the oxygen is present as Ti02 and B203 in a 1:1 ratio. Assuming a spherical shape and uniform grain size the volume fraction of the surface oxide can be expressed as 6t/d, where t = oxide layer thickness and d = particle diameter. This gives a tentative surface oxide layer thickness of 66-100 A, close to the approximately 50 A reported by Nord-Varhaug (1994). A BET-value of 0.83 m2/g for the surface area was measured. This is in agreement with a grain size in the 2 pm range when assuming a model of uniformly sized spheres giving BET = 1.33/d (m2/g) when d is the diameter measured in microns.
Typical impurities introduced during milling are metals like iron (from steels), tungsten (WC, hard metals), zirconium (Zr02) etc. depending on what materials are used as milling media and container, see e.g. Nord-Varhaug (1996). Some of these impurities can be removed by chemical cleaning procedures such as an acid wash, no such treatments have been performed in the present work, however. Chemical analysis of unm illed powder from the same producer showed only trace levels of metallic impurities other than Zr. Thus any metallic impurities in the milled powder 90 Development of microstructure in titanium diboride ceramics except from zirconium are believed to originate from the milling process. As shown in table 3.1 these are mainly Fe and Mo along with traces of W and possibly Cr and Ni. All of these elements are possible components in a tool steel which according to the producer was the only material in contact with the powder besides a plastic. No zirconium-rich phases were found in the powder. The zirconium is therefore believed to be dissolved in the TiB2 lattice in agreement with the observations of Mroz (1994) of a wide solid solubility of (Ti,Zr)B2. Zirconium may have been an impurity element in the titanium oxide raw materials used.
Thermodynamical calculations for selected reactions in the system of TiB2, surface oxides and BN along with B- or Ti-additions in a reducing atmosphere are included in appendix 2. In these calculations the activity of all the reactants are set equal to 1.
5.2.1 Pure TiB2
As expected both the density and the grain size of the final compacts increased as the sintering temperature was raised in the three hot pressing trials with no sintering additives (table 4.1). Only moderate grain growth has taken place at the two lowest temperatures. The grain size distribution is still broad, especially at the lowest sintering temperature where grains in the 2-3 pm range still exist along with the larger ones. At the highest sintering temperature the grain growth takes place much faster, and the smallest grains have disappeared. The average grain diameter is now twice that obtained at the lower temperatures. The fast grain growth has resulted in trapping inside the larger TiB2 grains of pores and BN particles that were earlier present in the grain junctions only. As most of the pores are now located inside the grains and not in the grain junctions, a further increase in sintering temperature or time will only result in grain growth without much further densification.
Microcracking of the TiB2 grains was found to be a problem only at the highest sintering temperature. This supports the predictions of Tennery et al. (1981) that the thermal anisotropy causes microcracking upon cooling when a critical grain size of Chapter 5: Discussion 91
-15 pm is exceeded. The microcracking was found to deteriorate the mechanical strength significantly, as demonstrated when the material was cut for TEM specimen preparation. It is also believed to ruin the positive effect of the reduced porosity and enhance the aluminium penetration during electrolysis. The optimum sintering temperature was therefore concluded to be in the range between 1880 and 1960 °C. A temperature of 1880 °C was chosen as the starting temperature in the rest of the hot pressing experiments.
The HREM studies of hot pressed TiB2 show the grain boundaries to be clean after sintering. No oxide phases are present, thus the surface oxide layer of the TiB2- powder has been removed during sintering. The surface oxides can be removed in two ways in a reducing atmosphere; by evaporation or carbothermic reduction. Evaporation of Ti02 is slow, but B203 has a low melting point, 450°C, and high vapour pressures at the sintering temperatures used. Lamoreaux et al. (1987) have calculated the high-temperature vaporisation of B203, see figure 2.5. B202 and B203 are the most important gaseous species under reducing conditions (Po2 = 10'15 bar) with partial pressures in the 10"1 - 10"2 bar range at temperatures approaching 2000K. No data on possible mixed boron-titanium-oxides are available.
Carbothermic reduction of surface oxides is possible thanks to the reducing atmosphere produced by the graphite oven. The Boudouard-equilibrium C + C02 = 2 CO is set, and contributes a transport mechanism for carbon from the graphite oven and oxygen from the surface oxide of the powder. The surface oxides can then be reduced to TiB2 following the same overall reaction as for the powder production from the same oxides. Also excess oxides, B203 or Ti02, can be reduced to B4C and TiC respectively, following the same reactions as described for excess oxides during powder production. No carbides indicating such reactions were found however. A combination of the above mechanisms, in which boron oxide evaporates, then is reduced to boron or boron carbide and finally reacts with the titanium oxide to form TiB2, could also be possible.
Unfortunately neither the exact nature and amount of different oxide species in the surface oxide layer of the TiB2-powder nor the nature of the oxide species evaporating during sintering have been identified. Thus it is not possible to conclude which of the 92 Development of microstructure in titanium diboride ceramics above mechanisms, evaporation or carbothermic reduction, dominates the oxygen removal. Both mechanisms do depend on a connected network of pores to exist in the early stages of densification as pointed out by Einarsrud et al. (1997). This network provides the path through which the gaseous species, whether carbon- or boron- oxides, can be transported during oxygen removal. If the densification is very fast, some of the oxygen will be trapped inside the sintered body.
The lack of secondary phases, oxides or other phases, at the grain boundaries after sintering in a reducing atmosphere raises the question of what mechanisms are important during sintering of TiB2. Unlike many other ceramics, in which a liquid oxide film of nanometer thickness is believed to aid the densification by enhancing diffusion and dissolution-precipitation, the later stages of TiB2 sintering must take place without such a film. Due to the immediate surface oxidation of the TEM- specimens in air similar observations from the pores can not be done. However, the same mechanisms that are responsible for removing the grain boundary oxide are believed to remove the oxides from the pores too. As shown in the literature review the existence of a surface oxide layer during sintering was proposed by most authors to explain the coarsening and grain growth. This conclusion seems reasonable in the earlier stages of sintering when an oxide layer is obviously present. As the removal of oxygen proceeds during sintering the importance of this layer to surface diffusion and evaporation diminishes, until at some stage it is completely removed. Thus oxygen is not believed to be responsible for grain growth in the final stages of solid state sintering in this experiment. No experiments to investigate the effect of a thicker initial oxide layer on the densification behaviour were performed.
The iron content of the TiB2 powder was ~0.2 wt%, and the EDS analysis showed a weak segregation to the grain boundaries after sintering. Most of the iron seem to be dissolved in the TiB2 grains, however, as was also observed by Zdaniewski (1985). As reported by several authors (chapter 2.3.2) even small amounts of iron are effective in enhancing densification. This must be ascribed to an increase in grain boundary diffusivity due to the iron segregation when there is not enough iron present to form a liquid grain boundary film. Due to the low total content of iron and the limited degree of segregation, the effect of iron to the sintering behaviour of the pure TiB2 samples is not believed to be very important. No segregations of other metallic elements to the Chapter 5: Discussion 93
grain boundaries were found. Except from Zr the content of these elements was very low, and their presence were not detected by TEM.
The BN particles are expected to decompose into free boron and nitrogen in an inert atmosphere if a sufficiently low nitrogen pressure can be established. This was not found to be the case in this work, the TEM investigation proved the existence of small BN particles in the sintered compacts. If all the nitrogen detected in the LECO analysis, table 3.1, is due to BN the volume fraction of this phase would be ~0.3 %. This seem reasonable in light of the TEM investigations, although no quantitative results concerning the volume fraction of different phases were derived.
5.2.2 TiB2 with boron addition
The addition of 2 wt% elemental boron prior to sintering was not found to cause any large changes in the sintering mechanism of TiB2. Comparing the microstructure with that of pure TiB2 sintered at the same temperature the grain size and shape is found to be the same, and solid state sintering is still believed to be the mechanism of densification. The added boron was not dissolved in the TiB2 matrix, but has formed B4C particles located in the grain junctions and at some grain boundaries. The porosity is significantly reduced, from 7% to 4%, when compared to the pure TiB2. The high density of small, solid B4C particles in the grain junctions and grain boundaries may have an inhibiting effect on grain growth by pinning the grain boundary movement. Kang et al. (1989) reported the addition of B4C particles to reduce grain growth during sintering with iron additions. Further sintering experiments at higher temperatures is necessary to confirm whether a boron addition has the same ability to inhibit grain growth by formation of B4C. The simultaneous addition of boron and titanium was hoped to enhance densification by a contribution of reactive sintering as demonstrated by Itoh et al. (1990). As seen from table 4.1 this was not found to be the case with the modest additions made in these experiments. OM showed a microstructure similar to that found with only the boron addition, no TEM studies of these two samples were performed. 94 Developmentof microstructure in titanium diboride ceramics
HREM shows that the surface oxides have been removed during sintering as was also the case when sintering without any additions. The evaporation of boron oxides may have increased at expense of the carbothermic reduction if elemental boron converts the titanium oxide to TiB2 and B203 according to the reaction
3 Ti02 (s) + 10 B (s) = 3 TiB2 (s) + 2 B203 (1/g)
This reaction is thermodynamically favourable at any temperature, and some of the boron addition may be consumed by it, but not all. Even if all the oxygen was bound as Ti02 and reacted according to the above reaction, half the boron would be left when the reaction had completed. The excess boron reacts with carbon to form boron carbide according to the net reaction
4 B (s) + C (s) = B4C (s)
The formation of a significant amount of B4C is confirmed both by TEM and light microscopy observations. The carbon content of the powder is well below that needed to complete this reaction, thus most of the carbon must originate from the graphite oven. It is then transported into the sample as carbon-oxides, following the Boudouard-equilibrium. This is the same mechanism as described above for carbothermic reduction of the surface oxides during sintering, and depends on the presence of an open network of pores. B4C exists in a wide compositional range, the carbon content ranging from 10 to 23 percent by weight according to the phase diagram. The original 2 weight percent boron addition gives rise to a 4 percent fraction of B4C by volume. Even if some of the boron is lost due to vaporisation, a significant amount is left in the sintered sample. This is in agreement with the microstructural observations, and shows that no significant amount of elemental boron has been able to enter into the TiB2 grains and change the stoichiometry of this phase. To calculate the true porosity of the TiB2 with B4C as secondary phase, the low density of B4C (2.52 g/cm3) compared to that of TiB2 (4.52 g/cm3) has to be taken into account. Assuming all the added boron has formed B4C, a theoretical density of 4.44 g/cm3 is obtained for the composite material. This is the value used to calculate the porosity of the samples with boron addition in table 4.1. Chapter 5: Discussion 95
As expected from the thermodynamic calculations the addition of boron did not affect the BN phase. Also the distribution of iron was unchanged from the pure TiB2 samples, which is expected from the lack of change in sintering mechanism upon boron addition.
5.2.3 TiB2 with titanium addition
As described in chapter 4 the microstructure of the samples hot pressed with a two weight percent titanium addition is significantly different from that of the other samples. The porosity is reduced to only 1%, and OM shows the TiB2 grains to have grown larger and obtained a more faceted shape when titanium is used as a sintering addition. Also a mixture of oxide phases is found in the grain junctions after sintering. These microstructural changes are all consistent with a liquid phase mechanism contributing to the sintering process. The oxide phases in the grain junctions must be formed in a reaction between the added titanium and the surface oxides of the TiB2 powder. The reducing sintering atmosphere has a very low partial pressure of oxygen, and can not be responsible for the formation of oxide phases. The titanium-oxygen phase diagram, figure 5.1, verifies the existence of a liquid titanium-oxygen phase at temperatures above 1700-1800 °C. Upon cooling from the hot pressing temperature, this liquid solidifies to form the cornered oxide phases observed in between the TiB2 grains and sometimes penetrating into the grain boundaries. Dissolution and reprecipitation of TiB2 in the intergranular liquid Ti-0 phase is believed to be responsible for the increased growth rate and the faceting of the TiB2 grains during sintering. 96 Development of microstructure in titanium diboride ceramics
Weight Percent Oxygen
Atomic Percent Oxygen
Figure 5.1 Binary phase diagram for the titanium-oxygen system. From Murray and Wriedt (1987).
The samples sintered with titanium addition were the only ones which contained oxide phases in the final, sintered material. In the other samples the surface oxides did evaporate as boron- or carbon-oxides through an open network of porosity in the early stages of sintering, leaving a final material with low oxygen content. With titanium added a few volume percent of liquid phase do form in the early stages of sintering due to melting of the elemental titanium and dissolution of the surface oxide in the titanium melt. Rearrangement of the TiB2 grains in this liquid phase causes a high initial densification rate. This rapid initial densification is believed to remove the open porosity present during the early stages of densification in the other materials. The transport of oxide vapours out of the material is stopped when the network of pores is closed. The same concept of a densification rate controlled loss of oxygen was proposed by Einarsrud et al. (1997) in liquid phase sintered materials with nickel as sintering aid. The rapid initial densification is confirmed by the densification curves given in appendix 1. Although unfortunately to inaccurate to probe any finer details in the densification behaviour due to the poor resolution, these measurements Chapter 5: Discussion 97 show that the onset of densification occurs more strongly and at a lower temperature when titanium is used as sintering aid. Also a higher final density is reached when these materials are compared to the ones sintered without a liquid phase present.
Some thermodynamic data and calculations relevant to the sintering of TiB2 with additions of metallic titanium are included in appendix 2. Because of the rapid densification and closure of pores no reducing atmosphere has been included, Pco = 0. Boron-containing surface oxides are thermodynamically unstable in the presence of titanium and react to form titanium oxides according to
2 B203 (1) + 5 Ti (1) = 2 TiB2 (s) + 3 Ti02 (1)
A further reduction of the titanium dioxide to oxides containing more titanium is possible:
3 Ti02 (1) + 3 Ti (1) = 2 Ti203 (1) + 2 Ti (1) = 6 TiO (1)
The final product depends on the relative amount of titanium and oxygen present. Assuming no oxygen has evaporated during sintering, an atomic ratio Ti/O = 0.75 results from the surface oxide and the titanium addition. This implies the formation of a mixture of Ti203 and TiO during cooling of the liquid oxide phase, and agrees with the TEM-investigations showing Ti203 to be the dominating oxide phase. From XRD (X-Ray Diffraction) measurements Ti203 is also known to form during liquid phase sintering of TiB2 with Fe and Ni additives (Jungling et al. 1994, Einarsrud et al. 1997).
In addition to the expected trigonal Ti203 phase, TEM investigations have identified the oxides Zr(Ti)02, Ti(ZrAl)02 and cubic Ti203. The elements in brackets are minor constituents only, and the uncertainties related to the EDS- and PEELS-quantification leaves the exact stoichiometry of the phases uncertain. Two of the oxide phases identified, Ti(ZrAl)D2 and cubic Ti203, could not be found in the ASTM-files. They have probably formed during cooling of the intergranular, titanium-rich liquid phase, and are likely to be metastable phases as they have not been reported in the literature earlier. The cubic Ti2Q3 phase was only observed once, and was completely surrounded by the Ti(ZrAl)D2 phase. This along with the interface faceting, the orientation relationship and the similarity in crystal structure described in the previous 98 Developmentof microstructure in titanium diboride ceramics chapter indicates a close relationship between the two phases. The cubic Ti203 phase was found to be unstable under electron irradiation at high accelerating voltages. The structure was determined in a 200 kV-microscope, and later investigations at 300 and 400 kV accelerating voltage caused the phase to transform into nanocrystallites of unknown structure. Both the Zr(Ti)02 and the Ti(ZrAl)02 phase contain considerable amounts of zirconium, whereas no zirconium-rich phases are present in the samples sintered without titanium addition. The zirconium must originate from the TiB2 powder which contained some minor amounts of zirconium, 0.36 weight percent, dissolved in the TiB2 matrix. As the matrix is dissolved in and reprecipitated from the intergranular liquid phase, the liquid is enriched in zirconium. Upon cooling zirconium-rich phases form from this liquid. The origin of the aluminium-content of the Ti(ZrAl)02 phase is not understood as no significant amounts of aluminium were detected in the chemical analysis of the powder.
An increased segregation of iron to the grain boundaries along with formation of iron- rich ribbons and particles in grain junctions is observed when the samples sintered with titanium addition are compared with solid state sintered samples. The increased iron content must be caused by the same scouring mechanism from the TiB2 matrix as proposed for the zirconium. This is supported by the EDS measurements which show there to be no iron left in the TiB2 matrix close to the grain boundaries. The high iron concentrations found at the grain boundaries along with the formation of ribbons of iron in the three-grain channels show that the iron phase is wetting TiB2 even better than the oxide phases, which are only found in connection with multi grain junctions. Both the strong segregations to the grain boundaries and the Fe-rich ribbons in the three-grain channels are believed to contribute to the improved sinterability of the samples with titanium addition.
Addition of two weight percent titanium changes the overall B/Ti-ratio from 2.00, as found by analysis of the powder, to 1.94. This composition is still within the equilibrium area of the TiB2 phase at the sintering temperature; 1.90 < B/Ti < 2.00. From this it was hoped that the titanium addition would be incorporated in the TiB2 matrix. This has not been found to happen. The reason is believed to be the surface oxides present on the TiB2 powder which reacts with the titanium and forms thermodynamically stable intergranular titanium oxide phases. To form single phase Chapter 5: Discussion 99
TiB2 materials by titanium addition, the oxygen content must first be significantly reduced. Due to the limited accuracy of PEELS quantification, small changes in the Ti:B stoichiometry would still not be easily detected. No reaction of titanium and TiB2 to form other borides, TiB or TigB^ were observed.
The above discussion has mostly been concerned with the central and largest part of the samples sintered with titanium addition. As opposed to the homogeneous microstructure found in samples sintered without titanium addition, the liquid phase sintering has given an inhomogeneous microstructure as described in chapter 4.2.2. The outermost layer, about half a millimetre in thickness, has a higher porosity level, an increased grain size and less secondary phases compared to the inner zones. This may be caused by an evaporation process taking place only at the surface; the liquid oxides evaporate from the outermost area, but the fast densification prevents evaporation from the bulk. An intermediate zone, another 1/2 - 1 millimetre wide, makes up the transition between the outer, porous zone and the central region. This zone has a low porosity indicating that less evaporation has taken place from here during sintering. The most frequently occurring secondary phase in this region is the fee Ti(CNO)-phase, which is not found elsewhere. As stated in chapter 4.2.2 this phase is probably a solid solution of the binary phases TiC, TiN and TiO, all with similar lattice structures. The TiN-part is believed to be formed during a reaction between the liquid titanium and BN:
2 BN (s) + 3 Ti (1) = TiB2 (s) + 2 TiN (s)
This reaction is thermodynamically favoured at all temperatures, and explains the lack of BN particles in the samples sintered with titanium addition. The solubility of nitrogen in liquid titanium at the sintering temperatures used is ~2 at% according to Massalski (1990). This is high enough to allow diffusional transport of small amounts of nitrogen to the outer zone, forming solid TiN only upon cooling. 100 Development of microstructure in titanium diboride ceramics
5.3 Aluminium exposure
5.3.1 Penetration of aluminium
As revealed by the SEM and TEM investigations all the investigated TiB2 materials were penetrated by aluminium during the 24 hours exposure test in liquid metal. The degree and nature of penetration were found to depend on the sintering temperature and the sintering additives used. Both the pure TiB2 samples and the ones sintered with titanium additions were heavily penetrated. SEM investigations showed aluminium to be present in the pores and grain junctions even two millimetres below the exposed surface. The samples sintered with boron addition gave no sign of aluminium penetration when examined by SEM.
Pure TiB2 and TiB2 sintered with boron addition: An obvious reason for the fast and strong penetration of the pure TiB2 samples is their relatively low density which implies that an interconnected network of pores is present after the sintering. Only the most dense one of the pure TiB 2 samples has a density for which closed porosity is assured, but this sample was microcracked due to the extensive grain growth. Thus all the three pure TiB 2 samples provide easily accessible open paths for penetration of liquid aluminium, and penetration may occur rapidly thanks to the excellent wettability of the system.
The open porosity does explain the fast and strong penetration of aluminium in the pure TiB2 samples. When the porosity is closed, as in the samples with boron addition, the speed and amount of penetration is, as expected, significantly reduced. TEM analysis reveals a second mechanism of aluminium penetration, not dependent on the presence of a network of pores, to be active. Both in the pure TiB2 samples and in the sample sintered with boron addition, aluminium was found to penetrate through the grain boundaries. This grain boundary penetration was found both close to the exposed surface and at depths of 80-100 pm below the surface. As described in chapter 4.4 the grain boundary wetting is not uniform. Some grain boundaries are wetted by thin, uniform films of aluminium, typically 10 nanometres in width, while an approximately equally large part of the boundaries contain no such film. Grain Chapter 5: Discussion 101 boundaries with discrete lenses of aluminium are also found. These variations are frequently observed at adjoining grain boundaries, sometimes even between different parts of the same grain boundary. Thus a difference in grain boundary wettability must cause the variations, not the local availability of aluminium.
As argued in chapter 2.4.2 the tendency for grain boundary wetting is governed by the magnitude of the grain boundary energy, y ss, and its solid-liquid interface counterpart, y si. Three important factors influence the grain boundary energy: the grain boundary chemistry, the grain boundary crystallography and the temperature. As shown in the literature review, grain boundary impurities are believed to play an important part both during sintering and aluminium penetration. Especially iron and oxygen impurities have been focused on. The present investigation could not detect any differences in grain boundary chemistry between the individual grain boundaries in the pure TiB2 or TiB2 plus boron samples. The HREM investigations showed no secondary phase to be present at the grain boundaries prior to aluminium exposure. Except from a small segregation of iron, equivalent to a quarter of a monolayer at the grain boundary, no difference in chemistry between the bulk material and the grain boundary regions was found. Any difference in chemistry between the individual grain boundaries in the pure TiB2 would have to be below the detection limits of TEM. Such differences are not anticipated to arise from the processing methods used to make the TiB2 materials. Thus it is not believed that the grain boundary chemistry can account for the observed differences in wetting behaviour between neighbouring grain boundaries.
To explore the effect of the grain boundary crystallography on the wetting properties, the misorientation of several neighbouring wetted and non-wetted grain boundaries was measured. As shown in figure 4.28 no relation between misorientation and wettability could be found, only a random set of points. This result was somewhat unexpected. Both experimental investigations of sintering behaviour of other ceramics such as Si2N4 (Pan et al. 1995) and theoretical considerations of grain boundary energies as function of misorientation (chapter 2.4.2) predicts a dependence. Low- angle and a few special high-angle grain boundaries were expected to have a low energy and should thus not be wetted. The fraction of non-wetted grain boundaries found (-0.5) is also significantly higher than the expected fraction of low-angle grain boundaries in a random polycrystal (-0.1 with an angle below -15°, Verhoeven 1975). 102 Development of micro structure in titanium diboride ceramics
The experimental data indicates that the actual fraction of low-angle boundaries is even lower than expected. Possibly, the grain growth during sintering has reduced the number of low-angle boundaries.
Although believed to be an important one, the angle of misorientation is only one out of eight parameters needed to describe the grain boundary crystallography completely (see chapter 2.4.2). Thus the lack of relationship between misorientation and wetting behaviour does not exclude the grain boundary crystallography as an explanation. A complete determination of all grain boundary parameters including the small atomic displacements at the interface is a difficult and time-consuming task. Only for special boundaries careful, HREM imaging combined with computer image simulations can give the complete answer, but this has been beyond the scope of the present work. The present results do however point at the inclination of the grain boundary plane as a possible important factor. The frequent observations of "mixed" grain boundaries do support this. "Mixed" means one part of the grain boundary is wetted while a neighbouring segment of the same boundary with a different orientation of the grain boundary plane is not. No detailed and systematic study of grain boundary plane orientation versus penetration has been performed yet.
In general the wettability increases as the temperature is increased. This is believed to be the explanation for the observed formation of lenses of aluminium at some grain boundaries. Some of the grain boundaries that were originally wetted by a thin film of aluminium at the exposure temperature of 1000 °C are believed to have dewetted as the samples were cooled to room temperature. The aluminium films have either retracted to the pores and grain junctions or formed the discontinuous lenses of aluminium at the grain boundaries. This is supported by the EDS results from the non- wetted grain boundaries which show a minor amount of aluminium, equivalent to approximately one monolayer at the grain boundary, to be present even here. As all the investigated grain boundaries in the pure TiB2 samples contain either metallic aluminium in the form of films or lenses, or traces of aluminium in the grain boundary lattice, the possibility exists that they were all wetted at the exposure temperature of 1000 °C. The reason why only some of the grain boundaries are dewetted during cooling is believed to be differences in grain boundary crystallography as described above. Chapter 5: Discussion 103
The term grain boundary wetting was deliberately used to describe the grain boundary penetration mechanism. The metallic films of several nanometres thickness found upon cooling prove the existence of a liquid aluminium grain boundary phase during high temperature aluminium exposure. The term grain boundary diffusion was used by Baumgartner (1984) to characterise a penetration behaviour similar to the one described here. No grain boundary film of aluminium was found, but "grain boundary precipitates" similar to the lenses of the current investigation and alu min ium EDS- signals from all the grain boundaries were reported. Diffusion of single aluminium atoms in the grain boundary lattice is an alternative to the dewetting mechanism in explaining the aluminium traces found at the clean grain boundaries. It does not apply to the 10 nm wide films however, and do not explain the formation of discrete lenses. The small number of intergranular microcracks left from the sintering process excludes these cracks as an important origin of penetration for all but the most coarse grained sample.
The sample sintered with boron addition differed from the pure TiB2 samples mainly in the amount of penetration and not in mechanism. The increased density explains the reduced speed of penetration as no open pore structure is present. The total absence of aluminium at most boundaries in the innermost specimen is believed to change into a structure similar to that of the pure TiB2 samples when the exposure time is prolonged.
Unlike for the pure TiB2 samples the SEM investigations failed to detect any aluminium penetration in the samples sintered with boron additions. Detailed investigations by TEM did, however, show some of the grain boundaries to contain thin films of aluminium as discussed above. The failure of SEM in detecting the grain boundary penetration of aluminium is due to the limited spatial resolution of the chemical information available by this bulk specimen technique. The large volume which emits X-rays (~pm3) compared to the thin (-10 nm) grain boundary films of aluminium makes detection of aluminium penetration a hard task. Only when larger pores filled with aluminium are present the penetration can be easily detected by SEM. This finding illustrates the importance of the high spatial resolution offered by TEM when investigating interface behaviour. It also questions the correctness of earlier observations of aluminium penetration behaviour in TiB2 or the lack of such 104 Development of microstructure in titanium diboride ceramics penetration, as many of these observations are based on methods with a poorer spatial resolution.
Titanium addition: In spite of the high density obtained during hot pressing the samples sintered with titanium addition experienced extensive penetration of aluminium during the 24 hours exposure. The microstructural investigations showed the grain junctions to contain aluminium and alumina to a depth of at least two millimetres behind the exposed surfaces. The porosity level of only one percent after sintering is well below that of the other samples in this study and leaves only a few, closed pores. Thus the mechanism of aluminium penetration through a network of open porosity does not contribute in the samples with titanium addition.
The grain boundaries of the titanium-added samples all contained aluminium after exposure. The thickness of the aluminium films was in general larger than in the other samples, some of them measuring up to a few hundred nanometres in width. Also the aluminium-TiBa interfaces had become faceted following the basal and prismatic planes of the TiB2 crystals. The reason for the increased amount of grain boundary penetration and the faceting is believed to be the change in grain boundary chemistry arising from the liquid phase sintering process. As shown in chapter 4.3.3 the liquid phase sintering has markedly increased the amount of iron segregation to the grain boundaries prior to aluminium exposure. Phase equilibrium studies in the TiB2-Fe system by Ottavi et al. (1992) have shown a considerable solubility of TiB2 in liquid iron at temperatures above a eutectic at 1320 °C. This is supported by the studies of Ghetta et al. (1992) who also showed the wetting of TiB2 by iron to be good. In addition they found the dissolution of TiB2 in iron to cause faceting of the TiB2 surface. The faceting of the TiB2-aluminium interfaces in the present investigation is thus believed to be caused by dissolution (and possibly some reprecipitation) of the TiB2 in a liquid aluminium-iron mixture. The presence of iron in this liquid is believed to increase the solubility of TiB2, which is otherwise very low in pure aluminium. The faceting of some TiB2 grains upon aluminium penetration was also observed by Nord-Varhaug (1994). This was also interpreted as being caused by dissolution of TiB2 in aluminium, but no detailed explanation or study of the grain boundary chemistry prior to exposure was performed. The increased width of the Chapter 5: Discussion 105 grain boundary films of aluminium in the samples sintered with titanium addition can also be understood from the same arguments. The assumed increased solubility in the liquid phase has enhanced the dissolution rate of the TiB2 grain boundaries and the ability of titanium and boron species to be transported away with the liquid. As the old grain boundary material is removed, this gives the possibility to form wider grain boundary films of aluminium without introducing high stresses in the material.
5.3.2 Reactions with secondary phases
When aluminium penetrates into the grain boundaries and grain junctions of the TiB2 materials, it will get in contact with any secondary phases left here during the powder production and sintering process. These phases will either be stable towards liquid aluminium, or they will react with aluminium and form new phases and/or be dissolved in the liquid aluminium. Possible equilibrium phases from the aluminium- rich part of the binary phase diagrams between aluminium and the most common elements in the TiB2 materials are included in table 5.2. From ternary phase diagrams (Petzow and Effenberg 1990-93) the aluminium containing ternary phases AlgB4C7 and Ti4oAl36024 are the ones of most interest. In addition several non-equilibrium phases are reported in JCPDS.
The solubility in aluminium of boron, titanium and the most important impurity elements in TiB2 are given in table 5.2. These data are valid for binary systems, and will be modified when more elements are present simultaneously. During normal cell operation conditions, with temperatures slightly below 1000 °C, none of the elements carbon, nitrogen or oxygen have any significant solubility in the aluminium. Thus they will be present only as solid carbide-, nitride- and oxide phases. On the other hand boron and metallic elements like titanium and iron have solubilities in the in the 1-10 atomic percent range. 106 Development of microstructure in titanium diboride ceramics
Table 5.2 Solubilities of various elements in aluminium for binary systems. Data from Massalski (1990).
Element Solubility in aluminium (at. %) Binary phase solid liquid liquid Tm(°C) (-660°C) (~660°C) (~1000°C) B <0.025 0.055 -2.5 A1B2 980 Ti 0.7 0.1 -1.7 Al3Ti -1350 C -0.005 -0.05 A14C3 -2200 N <1E-11 -IE-11 -0 AIN -2800 0 <3E-8 <3E-8 -0 ai2o3 2054 Fe 0.03 0.9 -9 FeAl3 -1160
Thermodynamical calculations predict many of the secondary phases to react with liquid aluminium. Relevant data are given in appendix 2. Due to the significant solubility of boron and titanium in aluminium, these elements are expected to dissolve in the aluminium during the reactions if a sufficient amount of aluminium is present. To account for this solubility in the calculations the end product of the reactions are taken to be boron or titanium dissolved in alu minium and not the A1B2 or Al3Ti phases. The activity of the dissolved species are set equal to 1 during the calculations. Using a lower activity value of the dissolved species or using A1B2 and Al3Ti as the final products will make the reactions more attractive energetically.
Pure TiB2: As earlier stated the TiB2 phase itself is practically inert towards aluminium, only a very low solubility product of 11 O'8 (wt%) is found (Finch 1972). This is supported by the increase in free energy of approximately 100 kJ/mole associated with the formation of aluminides from TiB2 and aluminium.
BN was the only secondary phase found in the materials sintered without any sintering additions. After exposure AIN was found in the same material. The similarities in shape and position between the BN and the AIN particles support the assumption that AIN has formed from the BN particles. The proposed reaction is Chapter 5: Discussion 107
A1 (1) + BN (s) = AIN (s) + B(diss)
This reaction is thermodynamically favourable. No A1B2 or other phases containing aluminium and boron were found throughout the specimen, supporting the assumption that the boron has dissolved in the liquid aluminium. When dissolved in aluminium, the boron may be transported far away from its origin by diffusion in the liquid phase. Only upon cooling of the material the solubility in aluminium will decrease, and eventually the A1B2 phase may precipitate if the boron content is above -0.025 at% which is the solid solubility limit in aluminium. The lack of A1B2 inside the sample after exposure indicates that most of the boron from the BN phase has diffused out of the sample and into the aluminium pool. This idea was also put forward Bardal et al. (1994) to explain the formation of A1B2 particles at the TiB2-aluminium interface upon cooling in wetting experiments of aluminium on TiB2. No AIN phase formed inside the TiB2 sample was found in those studies, however (Nord-Varhaug 1994). Not all the BN particles had finished the transformation to AIN after 24 hours of exposure, some were only contaminated with aluminium. This is attributed to the limited amount of alu minium that has penetrated into the material during this short time and a possible slow kinetics of the reaction. Fujii et al. (1993) reports the same reaction to take place in wetting experiments of BN by liquid aluminium, lowering the wetting angle slowly from 135° to zero in two hours.
Boron addition: B4C was still found to be present after exposure of the samples sintered with boron addition. Small amounts of aluminium were found surrounding the B4C, but the interior of the particles were not contaminated by aluminium. Possible reactions between the carbide and aluminium are
4 Al (I) + 3 B4C (s) = ALA (s) + 12 B (diss)
8 Al (1) + 7 B4C (s) = AlgBA (s) + 24B (diss)
The first reaction, formation of alu minium carbide, is associated with a slight increase in Gibbs free energy. This is in agreement with the observed lack of reactivity. If the formation of A1B2 is included, the thermodynamics strongly favour the reaction, 108 Development of microstructure in titanium diboride ceramics however. This reaction was reported to take place when B4C impurities are present by Gesing and Wheeler (1987), but no details were reported to prove the existence of the two new phases. The relative amount of aluminium and B4C present could be decisive as to whether the reaction takes place. The limited solubility of boron in aluminium implies the formation of A1B2 when only small amounts of aluminium are present. A14C3 is also known to form in the reaction between elemental carbon and aluminium at temperatures above 860 °C (Eustathopoulos et al. 1974). This is in agreement with thermodynamical calculations.
For the second reaction, formation of A18 B4C7 , data for the Gibbs free energy are missing. Several reports show this reaction to take place, however. Torvund et al. (1994) reported the formation of a phase containing aluminium, carbon and boron in a wetting study of B4C by aluminium. This was believed to be the A18 B4C7 phase which is also predicted by the ternary phase diagram (Petzow and Effenberg 1990-93). A slow decrease in wetting angle with time indicates that the kinetics of the reaction is slow. Pettersen and Bardal (1996) in a TEM-study identified the A18 B4C7 phase as a reaction product after aluminium exposure in a ZrB2 material with B4C as the original secondary phase. This material had a higher porosity than the presently studied TiB2 material, and was thus readily penetrated by large amounts of aluminium. Yet only parts of B4C had reacted to A18 B4C7 after 24 hours, showing the kinetics of this reaction to be slow. The slower penetration of the denser TiB2 material along with the small amounts of aluminium present in the interior of the material after 24 hours and the slow reaction kinetics must be the reason why no A18 B4C7 was found in the present study. During longer exposures, as e.g. a realistic cell test of some duration, all the B4C particles present are expected to react to A18 B4C7 .
Titanium addition: Several oxide phases, of which Ti203 was the dominating, were present in the grain junctions of the TiB2 materials sintered with titanium addition. After aluminium exposure these materials were found to be heavily penetrated, and metallic aluminium along with alumina were the dominating secondary phases in the grain junctions. The amount of alumina in the triple junctions resembles that of the oxides after sintering. The following reaction is proposed to have taken place: Chapter 5: Discussion 109
2 A1 (1) + Ti203 (s) = A1203 (s) + 2 Ti (diss.)
This reaction is thermodynamically feasible, and results in a reduction in the Gibbs free energy of 120 kJ/mole at 1000 °C assuming a titanium activity of 1. No Ti-Al phases were found, supporting the assumption that the titanium is dissolved in the aluminium during the reaction and transported away in the liquid. Besides trigonal Ti203 smaller amounts of the oxide phases Ti(ZrAl)02, Zr(Ti)02 and cubic Ti2Q3 were present in the grain junctions prior to sintering. These phases are believed to have reacted in a way similar to the usual trigonal Ti2G3 phase as no other reaction products than A1203 were found. Both for the Ti02 and the Zr02 phase the reaction into alumina is thermodynamically favourable. The formation of alumina as a reaction product during aluminium electrolysis with TiB2 cathodes containing oxide impurities has been suggested to be an important degradation mechanism by several authors. Actual identification of alumina was reported by Nord-Varhaug (1996), in this case the alumina had formed from a Si02 phase present in the grain junctions after sintering. The formation of ternary oxide phases such as Ti4oAl36024 should be possible according to Petzow and Effenberg (1990-93), but no indications of the formation of such phases were found in the present investigation.
Only parts of the Ti(CNO) phase found in the intermediate zone had reacted with the aluminium. As expected from thermodynamical calculations (appendix 2) some alumina had formed, whereas no aluminium carbide or -nitride phases were found.
5.3.3 Degradation mechanisms
Although no actual macroscopic degradation of the TiB2 materials were observed during the 24 hours exposure test, some possible causes of failures in long term tests can be discussed with reference to the observations made.
Swelling of secondary phase particles present in the grain junctions upon reaction with aluminium and subsequent stress-corrosion or cracking due to build-up of internal stresses has been proposed as a possible degradation mechanism by several authors, see chapter 2.4. Table 5.3 gives the molar volume of most relevant phases, and is used to predict whether a certain reaction results in swelling or not. 110 Development of microstructure in titanium diboride ceramics
Table 5.3 Density and molecular weight data for selected phases from "Handbook of Chemistry and Physics" (1992). Molar volume was calculated from these.
Phase Molar Density Molecular volume (g/cm3) weight (cmVmole) (g/mole) A1 9.99 2.70 26.982 A1B2 15.24 3.19 48.6 AI4C3 61.02 2.36 144 AlgB4C2 145.1 2.366* 343.2 AIN 12.57 3.26 40.99
P 25.72 3.965 101.96 Al3Ti ? ? 128.8 B (cryst.) 4.62 2.34 10.81 B4C 21.93 2.52 55.26 BN 11.03 2.25 24.82 B2O3 28.42 2.46 69.92 C (graphite) 5.22 2.3 12.0 Ti 10.64 4.5 47.88 TiB2 15.44 4.50 69.50 TiC 12.15 4.93 59.89 TiN 11.86 5.22 61.89 TiQ2 18.75 4.26 79.88 TiO 12.95 4.93 63.85 Ti2Q3 (trig.) 31.25 4.6 143.76
* Calculated value from JCPDS Chapter 5: Discussion 111
Considering only the solid phases of the reactions, not taking into account the liquid aluminium and the elements that are dissolved in this liquid, the following trends are found: * C (+290%) and TiC (+67%) are both very harmful if they react to form AI4C3. * BN (+14%) and TiN (+6%) give a moderate swelling upon reaction to AIN. * B4C (-7% / -5%) shrinks upon reaction to AI4C3 / AlgB4C7 . * The oxides TiO (-34%), Ti203 (-18%), Ti02 (-9%) and B203 (-10%) all shrink when they react to form alumina. The percentage given in brackets for each phase gives the increase (+) or decrease (-) in volume due to the proposed reaction. These calculations show that with respect to swelling any excess carbon else than that in the form of B4C must be avoided, whereas oxide phases are no threat to the material.
According to this the only harmful phase present in the currently studied materials is the BN phase. The TEM-studies gaveno evidence of cracking due to swelling of this or other secondary phases. For the BN phase this can be explained by a combination of the relatively small increase in molar volume associated with AIN formation, the small amount of BN present, the small particle size and the incomplete reaction. The B4C phase was not found to react in this study. As discussed in chapter 5.3.2 a slow formation of AlgB4C7 is known to take place in similar materials, but no evidence of swelling or cracking was found (Pettersen and Bardal 1996). This is in agreement with the calculated volume reduction following the reaction. Gesing and Wheeler (1987) also found B4C to react without causing any swelling.
After exposure alumina and some aluminium were found in the grain junctions that had previously been filled with titanium oxides. This is in agreement with the calculated shrinkage of the oxide leaving some space for metallic aluminium to solidify upon cooling, and will not cause build-up of mechanical stresses. Finch and Tennery (1982) reports no cracking during aluminium penetration of the material with the highest oxygen content (8 wt%). This supports the view that the presence of oxides do not cause failure of TiB2 materials due to swelling. The opposite view was taken by Gesing and Wheeler (1987). They claimed stresses to build up and stress corrosion to take place due to the volume increase associated with the reaction of oxides into alumina and A1B2 or Al3Ti. No evidence for the actual formation of these phases were given, however. In the present work neither A1B2 nor Al3Ti were found 112 Development of microstructure in titanium diboride ceramics as reaction products in any materials, and are thus not included when calculating the volume change associated with the reactions. No elemental carbon or TiC was found in the presently investigated materials. Both Gesing and Wheeler (1987) and Zdaniewski (1986b) reported formation of AI4C3 and cracking of materials containing C and TiC as secondary phases respectively. This agrees well with the current calculations.
Liquid metal films at the grain boundaries due to aluminium penetration is another possible, and from the above observations more probable, cause of cathode degradation for the currently investigated materials. The existence of such films is believed to reduce the mechanical strength of the grain boundary. This makes the material more vulnerable to any mechanical stresses and impacts occurring during normal electrolysis operation than an unpenetrated material. A reduction in the strength of TiBz upon immersion in liquid aluminium has been reported by both Dorward and Payne (1983) and Baumgartner (1984). This was accompanied by a transition in fracture mode from trans- to intergranular at the penetrated grain boundaries, an observation strongly suggesting a mechanical weakening of these by the aluminium.
A more dramatic possible effect of the grain boundary penetration is the complete dissolution of the grain boundaries causing whole grains of TiBa to detach from the cathode surface. This has been suggested as a degradation mechanism by several authors in materials with high grain boundary impurity levels (see chapter 2.4.1). Although not directly bringing about cracking and failure of the material such a mechanism will result in a strongly accelerated dissolution of the cathode. The present investigation points at iron as an impurity which could be responsible for increased grain boundary dissolution. The samples with the strongest iron segregation experienced both faceting of the grain boundaries and increased thickness of the aluminium films at the grain boundaries. Other impurities or surfactants segregated to the grain boundaries were not found in the present study. Chapter 5: Discussion 113
5.4 Further challenges
Several fundamental questions relevant to the possible use of TiB2 as inert cathode in an aluminium electrolysis cell still remain to be answered, some of which are outlined below.
What can be doneto improve the sintering process? The present work provides a starting point for further investigations of the sintering behaviour of TiB2. A more thorough theoretical study of the thermodynamics of the Ti-B-C-O-N system, including the effect of changes in temperature and sintering atmosphere, should be performed to optimise the firing scheme and atmosphere during sintering. Experimentally the production of a well defined starting powder with higher purity and smaller grain size is the single most important step towards a better understanding of the sintering process. Only then the individual effect of various sintering parameters such as sintering aids, impurities, temperature, pressure and atmosphere can be effectively pinpointed. In addition to the microstructural evolution the chemical composition, densification rate and evaporation during sintering should be more carefully monitored. Of special interest is the effect of varying the oxygen-level and the B/Ti-ratio of the powder. Also the effect of adding small amounts of carbon or metallic elements such as iron should be studied.
Does aluminium penetratea perfectly clean grainboundary? The present work has shown penetration to occur even at grain boundaries where no secondary phases are present. To answer whether this penetration is inherent to the TiB2-aluminium system, or only a result of changes in the chemical environment due to grain boundary segregations, more work is needed. The experimental challenge remains that of manufacturing a significantly purer material and precisely determine the grain boundary chemistry. A different approach is to model the penetration behaviour mathematically, a difficult task as input parameters such as grain boundary structure and interface energies are not known.
How do chemistry, crystallography and temperatureaffect penetration? Some insight in the effect of all of these parameters have been gained in the present study. An extension of the present TEM-work to include the orientation of the 114 Development of microstructure in titanium diboride ceramics grain boundary plane may explain the difference in penetration behaviour between adjacent grain boundaries. Iron was found to have a detrimental effect, increasing both dissolution and penetration. The effect, both to the sintering and the penetration behaviour, of other grain boundary segregants should be studied in a similar way. Exposure tests at different temperatures and subsequent quenching to freeze the high temperature penetration microstructure, might give an answer to how the temperature affects the aluminium penetration in TiB2. In particular it can tell whether the observed lenses and clean boundaries are a result of dewetting or the actual high- temperature microstructure.
How is wettability influenced by impurities/surfactants and crystal orientations? A different approach to help answering the above questions is to make well controlled surface wetting studies on TiB2. The effect of various impurities in the aluminium melt to the wetting angle could help determining their influence on the interfacial energies, thus guiding the choice of sintering aids to be used to reduce or promote grain boundary penetration. Using single crystal TiB2, the effect of varying the surface plane orientation on the wettability could be investigated.
What is the longterm effect of aluminium penetration? No actual degradation of the materials besides the penetration was found after 24 hours exposure to liquid aluminium. An extension to longer duration (weeks, months and years) is necessary to deter mine whether the microstructure obtained after 24 hours is an equilibrium structure or not. If the width of the aluminium films have stabilised and no aggressive grain boundary dissolution is experienced, the current materials may work satisfactory in spite of the penetration. Hence it may not be necessary to aim at a material with "no penetration" if this is proven difficult or impossible.
What additional stresses are introduced by polarised exposure? The present aluminium exposure experiments did not include polarisation of the material, exposure to a cryolite-based electrolyte or actual metal production in an industrial cell. Gesing and Wheeler (1987) compared different exposure test methods and showed the polarised and drained cathode tests to give a more severe degradation ChapterS: Discussion 115 than only unpolarized exposure for most materials. The survival of the material in such a test has to be the ultimate test for its possible use as an inert cathode. 116 Development of microstructure in titanium diboride ceramics 117
The present work has been concerned with two problems commonly encountered when carbon cathodes are to be replaced with titanium diboride (TiB2) cathodes in aluminium electrolysis cells. These are the difficulty of sintering high purity and high density TiB2-parts, and the degradation of the cathode during electrolysis. The TiB2 parts studied were hot pressed from carbothermically produced TiB2 powder. In some cases small amounts of titanium or boron was used as sintering aid. Detailed microstructural investigations have been used as the basis to understand the densification and degradation processes taking place. The main conclusions from this work are summarised below.
Sintering process Densities in the range 90-99 % of the theoretical density of TiB2 were reached during hot pressing depending on the temperature and sintering aid used. Undesired grain growth at high temperatures limits the useful densification as microcracking occurs during thermal cycling when the grain size exceeds ~15 pm. Solid state sintering takes place when no excess titanium is present. Boron addition gives a slightly increased density as a new phase is formed and fills up parts of the pores. A two percent titanium addition results in liquid phase sintering and consequently high density of the final product even at reduced sintering temperatures.
Secondary phases arise from impurities in the original TiB2 powder, from sintering aids added to improve densification and from the sintering atmosphere. Without sintering aids small boron nitride (BN) plates at some grain boundaries and grain junctions was the only secondary phase present in significant amounts. BN plates were also present in the initial TiB2 powder, and have probably formed as a by product during the carbothermic powder production. Addition of boron results in the formation of a boron carbide phase, B4C, usually found as submicrometer particles in TiB2 grain junctions, in addition to the BN. The carbon content of the B4 C phase originates from the graphite parts in the hot-press which creates a reducing CO- atmosphere. In samples sintered with titanium as a sintering aid several different metal-oxide phases form. These have been liquid at the sintering temperature, and have then solidified in the grain junctions during cooling. The oxygen originates from 118 Development of microstructure in titanium diboride ceramics the surface oxide layer always present at TiB2 powder. It reacts with the titanium and is trapped inside the rapidly densifying material. In the more slowly densifying samples without titanium addition an open network of pores is present in the earlier stages of densification. The surface oxygen then has enough time to evaporate as boron- and/or carbon-oxide gases through the pores, resulting in a material with very low oxygen content.
High resolution electron microscopy (HREM) proves the grain boundaries of sintered TiB2 to be atomically clean. No oxide phase, neither amorphous nor crystalline, is found at the grain boundaries in any of the materials. Segregations of iron impurities to the grain boundaries have been found, this effect is amplified by the liquid sintering process. The equivalent of one monolayer of iron at the grain boundary is found in a nanometer wide grain boundary region when titanium is used as sintering aid. An iron-rich phase forms as thin bands in the three-grain-channels in this material.
Aluminium exposure Scanning (SEM) and transmission (TEM) electron microscopy show that all the investigated materials suffer from aluminium penetration during a 24 hour unpolarised exposure test in pure aluminium at 980 °C. The extent of the penetration is found to depend on the microstructure. The open porosity and microcracking of the TiB2 samples sintered without any additives has provided a fast path of penetration as the pores are filled with liquid aluminium. The reduced porosity after boron addition has slowed down the penetration in these samples where SEM was unable to detect any penetration at all.
Grain boundary penetration was shown by TEM to take place in all samples. The penetration behaviour of a specific grain boundary is governed by its energy and thus depends on grain boundary chemistry, grain boundary crystallography and the temperature. In the pure TiB2 samples and the ones with boron addition thin wetting films of metallic aluminium were found at some boundaries. Other grain boundaries in the same area had no wetting film and were shown to be clean by HREM except for some occasional lenses of aluminium. Weak segregations of aluminium were still found however. This difference in penetration behaviour is believed to be caused by the grain boundary crystallography, although no direct relation to the grain boundary Chapter 6: Conclusions 119 misorientation alone could be found. Grain boundary dewetting upon sample cooling is suggested as the mechanism for the formation of lenses. The grain boundaries in the samples with titanium addition are more heavily penetrated than grain boundaries in the other materials, and the TiB2 grains develop facets at certain crystal planes. This is believed to be due to the iron segregations increasing the solubility, and thus the rate of dissolution, of TiB2 in the iron-aluminium grain boundary phase. Possible cathode degradation mechanism encouraged by the grain boundary penetration are reduced mechanical strength of the grain boundaries due to the liquid films and detachment from the surface of grains that become completely undermined by aluminium.
All of the secondary phases formed during hot-pressing are thermodynamically unstable in contact with liquid aluminium, but the kinetics of penetration and reaction delay some transformations. Parts of the boron nitride has finished a transformation to aluminium nitride, whereas no reaction was observed between boron carbide and aluminium in 24 hours. The oxide phases formed upon titanium addition react completely and form a-alumina. All these reactions contribute to the driving force of the aluminium penetration. No signs of stresses due to swelling of the secondary phases in the grain junctions upon reaction with aluminium were found, however. Such a mechanism has been suggested by other authors to explain cathode degradation by cracking, but was not supported neither by current calculations nor by experimental observations. 120 Development of microstructure in titanium diboride ceramics 121
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190.0 - 1600
188.0 1200 » 186.0
800 3 184.0
180.0 300 Time (min)
------Piston position (mm) — — Pressing force (kg) ...... Temperature (C)
"o -l 188.0 1£2. "□^ 1200 3 CD « 186.0 ?! O CD 184.0 ? o
Time (min) Figure Al l Hot pressing parameters recorded during sintering of pure TiB2 (upper) and TiB2 with 2 wt% titanium addition (lower). Temperature and pressure are kept equal in the two cases as seen from the graphs. The pressing piston position reflects the degree of densification as a reduction in piston position is equivalent to a corresponding reduction in specimen height. The initial rise in the piston position curve is due to thermal expansion of the system during heating and masks the true densification behaviour of the samples. The poor resolution of the piston position measurements makes a detailed description of the densification behaviour during sintering impossible. Appendix 2: THERMODYNAMICAL DATA AND CALCULATIONS
This appendix provides thermodynamical data and calculations that are used as a background for the discussion in chapter 5. First the standard Gibbs free energy of formation values, AGf°, at various temperatures for a selection of phases relevant to the present work on TiB2 are included. All data are taken from Barin I. "Thermodynamical data of pure substances" (1989, VCH, Weinheim). Melting points taken from the same source are included when available.
The calculations only include a limited number of reactions: The possible reactants and reaction products considered are included below. Also only the standard Gibbs free energy change of the reactions, AG°, are given. Thus the calculations are only valid when the activity of all the reactants and products equals 1. This is the case for the solid and most of the liquid phases (except from dissolved species). Corrections to include the partial pressures must be made when vapour phases are present, and will depend on the actual sintering atmosphere.
The calculations are divided into three parts:
The carbothermic production of TiB2 (Chapter 5.1) -Covering possible by-products when one of the reactants are in excess. -Reactants: B203, Ti02, C, N2, 02 -Products: TiB2, CO, BN, B4C, TiC, TiN, Ti02, B203, C
Sintering of TiB2 (Chapter 5.2) -Including the use of boron and titanium as sintering aids. -Reactants: B203, Ti02, BN, C, B- and Ti-additions -Products: B4C, TiC, CO, TiB2, B203, Ti02/Ti203/Ti0, TiN
Reactions during aluminium exposureof TiB2 (Chapter 5.3) -Reactants: Al, TiB2, BN, B4C, TiO, Ti203, TiC, TiN -Products: AIN, A14C3, A1203, A1B2, TiAl Gibb's free energy of formation, AGf, for selected compounds (kJ/mole)
Temperature TiB2 TiB B203 BN B4C TiO Ti203 Ti02 TIN TiC CO (K) (C) (cryst) (cryst) (cryst/liq/gas) (cryst) (cryst) (cryst/liq) (cryst/liq) (cryst/liq) (cryst) (cryst) (gas)
298 25 -320 -160 -1193 -229 -71 -513 -1434 -889 -309 -181 -137 500 227 -317 -159 -1139 -211 -70 -494 -1376 -852 -290 -178 -155 1000 727 -308 -157 -1019 -166 -69 -446 -1241 -763 -243 -173 -200 1300 1027 -302 -155 -954 -140 -68 -419 -1164 -709 -215 -169 -226 1500 1227 -297 -153 -912 -122 -66 -401 -1113 -674 -196 -166 -244 1600 1327 -295 -152 -891 -113 -66 -393 -1087 -656 -187 -165 -252 1700 1427 -293 -151 -870 -105 -65 -384 -1062 -639 -178 -163 -261 1800 1527 -291 -150 -850 -96 -64 -376 -1038 -621 -169 -162 -269 1900 1627 -288 -148 -829 -87 -63 -366 -1013 -604 -159 -160 -278 2000 1727 -285 -147 -809 -79 -62 -359 -987 -586 -150 -158 -286 2100 1827 -282 -145 -788 -70 -62 -352 -961 -568 -140 -156 -294 2200 1927 -279 -143 -768 -62 -61 -345 -940 -552 -130 -154 -303 2300 2027 -276 -140 -748 -53 -60 -339 -919 -537 -120 -152 -311 2400 2127 -271 -137 -735 -55 -332 -899 -522 -110 -149 -319 2500 2227 -264 -133 -726 -33 -46 -326 -878 -508 -100 -147 -327
Melting, decomposition and boiling temperatures: Tm (K) 3193 2500 723 2743 2023 2115 2130 3223 3290 Td (K) 2600 Tb (K) 2339 Carbothermic powder production
Standard Gibbs free energy c range of selected reactions (kJ/mo le):
Basic: Excess B203: Excess Ti02: Excess C/N/O:
Ti02+B203+5C B203+7/2C B203+3C+N2 TI02+3C Ti02+2C+N TIB2+5/202 TIB2+3/2C TIB2+3N Temp. vs. vs. vs. vs. vs. vs. vs. vs. (K) (C) TiB2+5CO 1/2B4C+3CO 2BN+3CO TiC+2CO TiN+2CO Ti02+B203 TiC+1/2B4C TIN+2BN
298 25 1077 747 324 434 306 -1762 104 -447 500 227 899 639 252 364 252 -1674 104 -395 1000 727 474 385 87 190 120 -1474 101 -267 1300 1027 231 242 -4 88 42 -1361 99 -193 1500 1227 69 147 -64 20 -10 -1289 98 -143 1600 1327 -8 102 -91 -13 -35 -1252 97 -118 1700 1427 -89 55 -123 -46 -61 -1216 98 -95 1800 1527 -165 11 -149 -79 -86 -1180 97 -70 1900 1627 -245 -37 -179 -112 -111 -1145 97 -45 2000 1727 -320 -80 -207 -144 -136 -1110 96 -23 2100 1827 -396 -125 -234 -176 -160 -1074 95 2 2200 1927 -474 -172 -265 -208 -184 -1041 95 25 2300 2027 -546 -215 -291 -237 -205 -1009 94 50 2400 2127 -609 -250 -308 -265 -226 -986 95 75 2500 2227 -665 -278 -321 -293 -246 -970 94 98 Sintering
Stand ard Gib os free energy change of selected reactions (kJ/mole):
No addition: Boron addition: Titanium addition: 1/2* 2B203+7C Ti02+3C 4B+C TiO2+10/3B B203+5/2Ti 3Ti02+Ti Ti203+Ti 2BN+3Ti Ti+C Temperature vs. vs. vs. vs. vs. vs. vs. vs. vs. (K) (C) B4C+6CO TiC+2CO B4C TiB2+2/3B203 3/2Ti02+TiB2 2TI203 3TiO TiB2+2TiN TiC
298 25 747 434 -71 -226 -461 -201 -105 -709 -181 500 227 639 364 -70 -224 -456 -196 -106 -686 -178 1000 727 385 190 -69 -224 -434 -193 -97 -628 -173 1300 1027 242 88 -68 -229 -412 -201 -93 -592 -169 1500 1227 147 20 -66 -231 -396 -204 -90 -567 -166 1600 1327 102 -13 -66 -233 -388 -206 -92 -556 -165 1700 1427 55 -46 -65 -234 -382 -207 -90 -544 -163 1800 1527 11 -79 -64 -237 -373 -213 -90 -533 -162 1900 1627 -37 -112 -63 -237 -365 -214 -85 -519 -160 2000 1727 -80 -144 -62 -238 -355 -216 -90 -506 -158 2100 1827 -125 -176 -62 -239 -346 -218 -95 -492 -156 2200 1927 -172 -208 -61 -239 -339 -224 -95 -477 -154 2300 2027 -215 -237 -60 -238 -334 -227 -98 -463 -152 2400 2127 -250 -265 -55 -239 -319 -232 -97 -448 -149 2500 2227 -278 -293 -46 -240 -300 -232 -100 -431 -147 Standard Gibbs free energy of formation for reaction products (kJ /mole)
Temperature Al AIB2 AIB12 AI4C3 AIN AI203 TiAl TI3AI (K) (C) (cryst/liq) (cryst) (cryst) (cryst) (cryst) (cryst) (cryst) (cryst) 298 25 0 -149 -272 -196 -287 -1582 -73 -140 500 227 0 -148 -276 -188 -266 -1519 -72 -135 1000 727 0 -143 -279 -163 -212 -1361 -68 -121 1100 827 0 -141 -278 -154 -200 -1328 -66 -115 1200 927 0 -138 -277 -144 -188 -1295 -64 -108 1300 1027 0 -136 -275 -134 -176 -1262 -62 -102 1400 1127 0 -274 -125 -165 -1229 -60 -95 1500 1227 0 -273 -116 -153 -1196 -58 -89 Tm=933 fd=1673 Tm=2423 Tm=2500 Td=2790 Tm=2327 Td=1733 Td=1613
Aluminium exposure Standard Gibbs free energy change of selected reactions (kJ/mole): No addition: Boron addition: Titanium add ition: TIB2+2AI BN+AI 3B4C+4AI 3B4C+10AI 3TIO+2AI TI203+2AI 3TIC+4AI TiN+AI vs. vs. vs. vs. vs. vs. vs. vs. (K) (C) AIB2+TIAI AIN+B AI4C3+12B AI4C3+6AIB2 AI203+3TI AI203+2TI AI4C3+3TI AIN+Ti
298 25 98 -58 17 -877 -43 -148 347 22 500 227 97 -55 22 -866 -37 -143 346 24 1000 727 97 -46 44 -814 -23 -120 356 31 1100 827 99 -43 53 -793 -17 -113 362 34 1200 927 102 -40 60 -768 -14 -105 369 37 1300 1027 104 -36 70 -746 -5 -98 373 39 1400 1127 -34 76 1 -91 379 41 1500 1227 -31 82 7 -83 382 43