The crystallization of wollastonite-type, (Ca, Mg, Fe)SiO₃, in the system of Al₂O₃-CaO-Fe₂O₃-K₂O-MgO-MnO₂-Na₂O-P₂O₅-SiO₂-TiO₂ has been studied over the temperature range of 1148-1348 K. A linear growth rate of surface crystallization of the wollastonite-type has been observed for the glass system. The crystal growth rate was measured as a function of temperature in a platinum crucible contact regime where heterogeneous crystal has been found to occur. The beginning of the crystallized-layer growth was preceded by a considerable induction period. The activation energy barrier for crystallization of the wollastonite-type at large undercoolings was determined to be 222.20 kJ•mol^-1.

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The migration of Ca ions during thermal decomposition of Ca-deficient hydroxyapatite (Ca-def HAp) involving the formation of β-TCP was discussed in the present paper. The migration of Ca ions at high temperature was examined by studying the thermal behavior of Sr ions in Sr-substituted Ca-def HAp using transmission electron microscopy (TEM) and energy dispersive X-ray spectroscopy (TEM-EDX). Substitutions of Ca ions in Ca-def HAp by Sr ion were carried out by hydrothermal treatment in Sr(NO₃)₂ solution. The composition of samples obtained by hydrothermal treatment was identified as Ca_9.27Sr_0.27(HPO₄)_0.46(PO₄)_5.54(OH)_1.54. In addition, it was considered that Sr ions would be in Ca(II) site in Sr-substituted Ca-def HAp lattice. Since the thermal behavior of Sr-substituted Ca-def HAp was almost the same as that of Ca-def HAp, the migration of Sr ions in Sr-substituted Ca-def HAp during thermal decomposition could be regarded as the same as that of Ca(II) ions in Ca-def HAp. By TEM and TEM-EDS analysis, it was revealed that Sr ions were detected only in β-TCP phase after thermal decomposition. Therefore, these results suggested that β-TCP phase would be formed by the priority migration of Ca(II) ions in Ca-def HAp.

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Influence of heat treatment at 1000-1500°C for 24 h on phases, microstructures and mechanical properties was studied for laminated SiC satin or aluminosilicate plain fabric/mullite matrix composites. Composites were produced by a polymer impregnation and pyrolysis method (PIP) using a mullite precursor. A mullite precursor solution of the mixtures of Si(OC₂H₅)₄ and Al(NO₃)₃ was impregnated into the green layered fabric/mullite powder (filler) composite and was pyrolyzed at 500°C in air. This polymer impregnation and pyrolysis process was repeated 10 times. The SiC fabric (25-27 vol%)/mullite matrix composite with a relative density 72-75% showed weight loss after heat treatment in an Ar atmosphere (oxygen partial pressure ~0.75 Pa) at 1300-1500°C. This weight loss was due to the formation of volatile SiO gas resulting from the pyrolysis of mullite matrix (3Al₂O₃•2SiO₂(s)→3Al₂O₃(s)+2SiO(g)+O₂(g)) and active oxidation of the SiC fiber (SiC(s)+O₂(g)→SiO(g)+CO(g)). The heat-treated SiC fabric/mullite matrix composites showed an elastic deformation in the initial stage of the stress-displacement curve, followed by pseudoductility. The strength (186 MPa) after the 1000°C heat treatment decreased to 17 MPa after the 1500°C heat treatment. The energy of fracture was 1-4 kJ/m² in the heat treatment temperature range from 1000 to 1500°C.On the other hand, the aluminosilicate fabric (25-30 vol%)/mullite matrix composite with a relative density 71-73% showed no weight loss after heat treatment at 1000-1500°C in air. The porosity of the composite decreased with increasing heat treatment temperature because of the densification of the mullite matrix. This composite also showed pseudoductility after the heat treatment. The strength for the composite was slightly enhanced by increasing the heat treatment temperature.

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Submicrometer-sized α-SiC powder was mixed in an aqueous solution at pH 5.0 with the following Al₂O₃ (1-2 vol% against SiC)-Y₂O₃ (0.94-1.6 vol%) sintering additives: sample A-Al₂O₃ powder plus Y₂O₃ powder, sample B-Al₂O₃ powder plus Y₂O₃ powder and polyacrylic acid dispersant, sample C-Al₂O₃ powder plus Y³⁺ ion adsorbed and sample D-Y³⁺ ion adsorbed. The prepared suspensions were consolidated by filtration through a gypsum mold. Green compacts were densified to a 96.2-99.3% relative density at 1800-1950°C under a pressure of 39 MPa for 2 h in an Ar or reduced atmosphere. The sinterability of SiC became higher in the following order: sample D<sample C<sample A≈sample B. The sinterability of samples A and B was comparable to that of SiC hot-pressed with Al₂O₃. Dense SiC with Y₂O₃ or Al₂O₃ plus Y₂O₃ showed the following excellent mechanical properties: average four-point flexural strengths of 565-719 MPa, fracture toughness of 5.0-6.2 MPa•m^1/2, Weibull modulus of 5.4-11.4 and Vickers hardness of 19-22 GPa. Suppression of grain growth in SiC enhanced the mechanical properties.

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A simple process was developed to produce transparent MgO ceramics. The Existence of chloride ion was one of the necessary conditions to produce highly transparent MgO ceramics. In this process, magnesium carbonate was initially precipitated by means of dropping magnesium chloride solution to a sodium carbonate solution. This magnesium carbonate solution was kept for 18h at 35°C and resulted in magnesium carbonate basic. 1 to 10 mol% of ammonium chloride was added to magnesium carbonate basic. And the powders were calcined at 900°C for 4 h. Doping of chloride ion resulted in the collapse of hard agglomerates. Sintering was carried out at 1600°C for 3 h in an oxygen atmosphere.

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Tests were performed on Al₂O₃-Y₂O₃ based compounds, which are the most widely used sintering additive for Si₃N₄, with the objectives of evaluating the high-temperature viscosity characteristics of Al-Y-Si-O based glass and clarifying the relationship between high-temperature viscosity and the high-temperature strength of Si₃N₄. The viscosity of the glass at 1600°C was found to be about 8-50P, while the activation energy for high-temperature viscosity was about 9-37 kJ/mol. Properties of glass such as viscosity are governed by the composition of the glass; for compositions within the glass forming region, it is thought that the network structure, hardens and viscosity becomes high in compositions having a substantial amount of SiO₂, or in compositions where the ratio of Y₂O₃/Al₂O₃ (close to the eutectic composition) is 1.0. For compositions outside the glass forming region, the influence of the precipitation of the crystal phase is considered to become more prominent. The room temperature strength of Si₃N₄ is not governed by the properties of the glass phase, but by micro structural factors such as the residual amount of pores. However, the properties of the glass phase in the sintering process are considered to determine structural factors such as the density of sintered bodies and the morphology of Si₃N₄ grains.The high-temperature strength of Si₃N₄ is dependent on the high temperature viscosity characteristics of the glass phase, and this was seen to impart high high-temperature strength to Si₃N₄ for which glass with high high-temperature viscosity would exist at the grain boundaries. For Si₃N₄ with similar viscosity characteristics at high temperatures, high-temperature strength is dependent on the size of the Si₃N₄ grains, seen to result in superior high-temperature strength for Si₃N₄ with large Si₃N₄ grains. Correlation between the rate of change in 1300°C strength and the activation energy of viscosity for Si₃N₄ was observed when evaluated in high temperature strength tests conducted at slow cross-head speeds of 0.1mm/min to 0.5 mm/min, and it was found that the rate of change in the 1300°C strength tended to become smaller for Si₃N₄ compositions with lower activation energy.

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The purpose of the present experiments is to consider the relations between coloration and crystalline phases of lime-glaze, dolomite-glaze and lime-barium-glaze containing 15 mass% ferric oxide and 15 mass% phosphate. The result of the experiments showed that red color under hard oxidizing fire only appeared in the glaze with MgO, Fe₂O₃ and P₂O₅. As a result of X-ray diffractometry, the colour of red was exposed by magnesioferrite-crystal. However, magnesioferrite-crystal can exist as whitlockite-type crystals. The hard oxidizing fire iron red glaze is one of the crystalline glazes which is colorized by magnesioferrite during cooling process from firing. In addition, magnesioferrite crystal formation is closely related to the presence of whitlockite-type crystals.

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The purpose of this paper is to interpret the enhancement of T_c and J_c by the replacement of cations and changing the excess oxygen content for Bi-2223 system by the replacement of Bi and Ca by Pb and Y, respectively, that is Bi_2.20-zPb_zSr_1.86Ca_2.03Y_0.05Cu_2.86O_y. The relations between lattice constant, average valences of Bi+Pb and Cu, excess oxygen content and superconducting characteristics have been studied by changing Pb substitution content, and annealing treatment. Bulk samples were prepared by the coprecipitation method. Furthermore, in order to elucidate the effect of metal substitution and oxygen content on J_c of the Bi-2223 system, films which have the same metal composition of the bulk, were prepared on MgO (100) single crystal substrate by the spin coating-pyrolysis method. Measurements of J_c and T_c were carried out on highly oriented films of Bi-2223 phase with changing Pb content and oxygen content: T_c (zero) experienced a maximum at Pb content, z=0.50 (nominal composition), and T_c (zero) decreased with increasing z in the region z>0.50. On the other hand, J_c also showed a maximum at z=0.50, but it was constant with increasing z in the region z>0.50. Oxygen underdoped and overdoped regions, and the highest T_c (zero) and J_c were obtained upon adjusting the optimum excess oxygen content, by changing Pb content and annealing treatment.

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Influence of composition deviation from around Ba(Zn_1/3Ta_2/3)O₃ (BZT) on superlattice ordering and microwave Q factor were studied. The superlattice ordering ratio (ordered-disordered structure formation) was found to be dependent on slight composition deviations around BZT. Ordered structure with perovskite single phase was obtained in the vicinity of the stoichiometric BZT. In this region, the significant high Q factor (Q•f=110000 GHz) was found. Disordered perovskite with lower density and secondary phase in the ordered perovskite were thought to be the causes of Q factor decrement. Suppressions of the secondary phases and the disordered phases by strict composition control can provide further Q factor improvements for the BZT system.

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Due to the consumption of huge amount of energy in the ceramic industry, its energy assessment is significant. This paper provides on overview of energy consumption in ceramic fabrication including powder synthesis, powder grinding, forming and sintering. Furthermore, we obtained energy consumption in each fabrication stage of alumina ceramics and its total energy consumption as an example case, and discussed the relationship between fabrication processes for energy-saving and optimum particle size.

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In order to fabricate machinable spinel/mica composite, 40 mass% mica-composition (KMg₃AlSi₃O₁₀F₂) glass powder was mixed with magnesia and alumina powders and the powder mixture was sintered. Furthermore, in order to improve the mechanical properties of the composite, silicon carbide was added to the powder mixture. The densification and the microstructure development of spinel/mica composite progressed drastically at 1150-1200°C. The dense machinable spinel/mica composite of which relative density was about 99% was obtained at 1200°C. The machinability was caused by the interlocking microstructure of mica crystals. The addition of silicon carbide retarded the densification and the grain growth of spinel and mica. However, the sintered composite also showed machinability and the fracture toughness and hardness were enhanced by the silicon carbide. Though the sintered composite was not densified sufficiently by the addition of silicon carbide, the finer microstructure caused a little improvement of the bending strength.

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AlZrC₂ powder prepared by solid-state reaction was sintered using a pulse electric current sintering technique in a temperature range from 1700 to 2000°C under the pressure of 40 MPa in vacuum for 5 min. The density and mechanical properties of the sintered AlZrC₂ monolith were estimated. As heating temperature increased, the density of the AlZrC₂ monolith increased and the mechanical properties were improved. The bulk density of the AlZrC₂ monolith sintered at 2000°C was 5.12 g/cm³. The maximum values of the fracture strength and Vickers' hardness were 380 MPa and 11.1 GPa, respectively.

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Wrong:A vector material is defined as the one to manipulate surrounding inorganic and organic compounds, microorganisms, and tissues using a force irradiated locally and continuously from the material itself. The effect given by a vector material is ca materials are of unparalleled importance under the circumstance of the limited use of external forces such as in vivo. We have recently experimentally demonstrated the effectiveness of electrically polarized hydroxyapatite (HA) electrets as vector ceramics, which can independently irradiate an electrostatic force to the surrounding constituents; crystal growth from a simulated body fluid was accelerated or decelerated and proliferation of a microorganism was controlled on the surfaces of polarized HA, depending upon the electric signs of surface charges. Prior to our study, certain types of vector effects have already been employed in biomedical applications; the so-called 45S5-type bioactive glass (Bioglass^®), developed by Hench et al., has long been recognized to be bioactive because of its solubility in a aqueous medium. Radioactive ceramics irradiate β-ray to and work on surrounding cancer cells. These are also vector materials. In addition to these examples, various kinds of vector effects and ceramics are now under development. This advanced report presents the concept of vector effects and reviews the examples of vector ceramics.
Right:A vector material is defined as the one to manipulate surrounding inorganic and organic compounds, microorganisms, and tissues using a force irradiated locally and continuously from the material itself. The effect given by a vector material is called the vector effect. From the viewpoint of bio-interface engineering for the manipulation of their surrounding media, vector materials are of unparalleled importance under the circumstance of the limited use of external forces such as in vivo. We have recently experimentally demonstrated the effectiveness of electrically polarized hydroxyapatite (HA) electrets as vector ceramics, which can independently irradiate an electrostatic force to the surrounding constituents; crystal growth from a simulated body fluid was accelerated or decelerated and proliferation of a microorganism was controlled on the surfaces of polarized HA, depending upon the electric signs of surface charges. Prior to our study, certain types of vector effects have already been employed in biomedical applications; the so-called 45S5-type bioactive glass (Bioglass^®), developed by Hench et al., has long been recognized to be bioactive because of its solubility in a aqueous medium. Radioactive ceramics irradiate β-ray to and work on surrounding cancer cells. These are also vector materials. In addition to these examples, various kinds of vector effects and ceramics are now under development. This advanced report presents the concept of vector effects and reviews the examples of vector ceramics.

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