The Knoop hardness measurements of NbB₂ and TaB₂ single crystals were made in the temperature range 20°-1100°C. The long axis orientations of the indentations made were along the ‹1120› and ‹1010› directions on the (0001) plane and along the [0001] and ‹1210› directions on the {1010} plane. Hardness decreased monotonously through all the testing temperatures, maintaining the relationships of H_{‹1010›(0001)}>H_{‹1120›(0001)} and H_{‹1210›{1001}}>H_[0001]{1010}. The hardness values of the (0001) plane were higher than those of {1010} plane in the given temperature. ln(H) vs. T kept linear relationship having steps near 0.15 and 0.3T_m. Predominant prismatic slips were observed at higher temperatures.

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Severe wear of lining refractories in torpedo car occurs when injection of desulphurizing agent containing calcium carbide is done to desulphurize molten pig iron. To make clear the corrosion mechanism of the refractories, reactions between calcium carbide and Al₂O₃-SiO₂ ceramics were studied. Three kinds of mullite and corundum disks were embeded in calcium carbide powder, and heated at 1200°-1400°C in a reducing atmosphere. Mullite ceramics formed a reacted zone, leaving metallic silicon particles in it. The thickness of the reacted zone is proportional to the square root of time. Thin boundary zone, which is mainly consisted of alumina, is formed between the reacted and the original ceramic zones. On the other hand, corundum ceramics did not form any reaction zone.
From these results, SiO₂ component in the mullite ceramics is considered to be reduced by CaC₂, forming metallic silicon, calcium oxide, and carbon monoxide or carbon. The reaction would proceed by the following steps:
(1) Transportation of CaC₂ through the reacted zone and the boundary zone.
(2) Decomposition of mullite: 3Al₂O₃⋅2SiO₂→3Al₂O₃+2SiO₂
(3) Reduction of SiO₂ by CaC₂: 3SiO₂+CaC₂→3SiO(g)+CaO+2CO
(4) Transportation of SiO gas through the boundary zone
(5) Reduction of SiO by CaC₂: 3SiO+CaC₂→3Si+CaO+2CO
The rate of the reaction is considered to be controlled by transportation of CaC₂ through reacted zone and boundary zone.

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Calcium phosphate ceramics with high mechanical strength are of potential use as artificial bones. Fundamental conditions for preparing this kind of ceramics by unidirectional solidification of their melts were investigated. Starting materials with nominal compositions of the 2CaO⋅P₂O₅-3CaO⋅P₂O₅ eutectic and its neighbors (Table 1) were prepared by melting the mixtures of reagent grade chemicals of CaHPO₄⋅2H₂O and CaCO₃ in a platinum crucible at 1400°C for 30min and casting them into another platinum crucible. The starting materials were remelted except for their lowest part about 15mm thick in the platinum crucible placed in a temperature gradient furnace (Fig. 2) and unidirectionally solidified upwards under various conditions (Table 2). The lowest part of the starting material left unmolten acted as seed crystals. When the starting material with the exact eutectic composition (E+1.0P₂O₅ in Table 3) were first remelted below 1350°C and then unidirectionally solidified at rates of 2-20mm/h under a thermal gradient of 20°C/cm (d-d′′′ in Table 2), ingots with regular structure were obtained (Fig. 5, E+1.0P₂O₅), In these ingots, lamellar α- and β-3CaO⋅P₂O₅ crystals were aligned parallel to the solidification direction in a α-2CaO⋅P₂O₅ crystal matrix. The crystallographic direction perpendicular to the (002) plane of the α-2CaO⋅P₂O₅, and those perpendicular to the (113) and (220) planes of the α- and β-3CaO⋅P₂O₅ crystals, respectively, coincided all with the solidification direction (Fig. 6), The lamellar spacing λ decreased from 16.5 to 6.0μm with increasing solidification rates R from 2 to 20mm/h according to the relation λ∝R^-1/2, like for other eutectic systems (Fig. 7). The ingots were pore-free but contained some microcracks, as shown in Fig. 8. These microcracks arose probably from thermal stresses due to a large difference in thermal expansion between the 2CaO⋅P₂O₅ and 3CaO⋅P₂O₅ crystals, which were formed during cooling of the solidified ingots (Fig. 9 and Table 5). When the composition of the starting material deviated from the eutectic even by 0.5%, the resultant unidirectionally solidified ingots showed irregular structure and contained many pores (Table 3 and Fig. 5). The compositional deviation were found to occur when the starting material was prepared by melting above 1400°C, where the P₂O₅ constituent tended to vaporize separately from the melt. The formation of the irregular strlucture was attributed to a constitutional supercooling of the melts at the front of the growing crystals (Figs. 10 and 11).

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(1) TiO₂-rich melts in the BaO-SiO₂-TiO₂ system were formed into plates (1-3mm thick) and fibers (60-300μmφ) by pressing between two steel plates and by drawing up with the tip of a Pt rod, respectively. The compositional region from which clear glass plates and fibers were obtained are shown in Fig. 2. Automatic drawing of fibers through a Pt orifice was difficult for the above compositions, but became feasible by substituting Na₂O, CaO and SrO for BaO and ZrO₂ for TiO₂, partially (Fig. 1).
(2) 2g of grains (297-500μmφ) converted from the plate glasses of the BaO-SiO₂-TiO₂ system was immersed in a 100ml of 2N NaOH aqueous solution at 95°C for 18h. The weight losses due to corrosion of the glasses having TiO₂/SiO₂ mole ratio higher than 0.75 were less than that of G 20 glass (1Li₂O⋅11Na₂O⋅1Al₂O₃⋅71SiO₂⋅16ZrO₂, wt%). On the contrary, however, diameter reduction of the fibers of the same compositions, were much larger than that of G 20 glass, when 5 pieces of them (250-300μmφ×5mm) were immersed in the same solution (Fig. 3).
(3) The diameter reduction of the fibers of the BaO-SiO₂-TiO₂ system was decreased by partial substitutions of CaO or SrO for BaO (Figs. 4 and 5). The diameter reduction of fibers of S4-15Sr glass (15SrO⋅15BaO⋅30SiO₂⋅40TiO₂, mol%) could be further supressed by addition of a small amount of S4-15Sr glass grains or powdered chemicals of Sr(OH)₂⋅8H₂O, Ba(OH)₂⋅8H₂O and/or TiO₂⋅1.5H₂O to the solutlon (Fig. 6, Table 3). Presence of a TiO₂-rich thin layer on the surface of S4-15Sr glass fibers was confirmed by electron probe microanalysis after the fibers were immersed in the NaOH solution (Table 4).
(4) On the basis of the above results, alkali corrosion process of the TiO₂-rich BaO-SiO₂-TiO₂ glasses was suggested as follows: The NaOH solution near the surface of the glasses would firstly be saturated with TiO₂ dissolved from the glasses because of its extremely low solubility, leading to formation of a thin TiO₂-rich layer on the glass surface. Ba²⁺ ions dissolved from the glasses would be adsorbed by the TiO₂-rich layer because of their high affinity to TiO₂. The TiO₂-rich layer thus formed would suppress further dissolution of BaO and SiO₂ from the interior of the glasses.

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The freeze-dried sulfate was prepared from the equi-mole mixed aqueous solution of magnesium and aluminum sulfates and calcined at 900°-1400°C for 1h. These calcined spinel powders were characterized and their sinterability was investigated. The results obtained were as follows:
(1) The freeze-dried sulfate was amorphous and its composition was MgAl₂(SO₄)₄⋅8-9H₂O which corresponded to partially dehydrated MgAl₂(SO₄)₄⋅22H₂O.
(2) The sulfate completely decomposed to spinel above the firing temperature of 1000°C.
(3) Powders obtained from the calcination of sulfate at 1100°-1300°C were nearly perfect spinel crystals, showing that their lattice parameter was 8.082Å, but their apparent crystallite size was small, about 600Å, and their specific surface area was relatively large, 20-30m²/g.
(4) Initial sintering of these specimens was mainly controlled by volume diffusion mechanism and the apparent activation energy of sintering for the specimen calcined at 1300°C was 127kcal/mol.
(5) The bulk density of these speciments fired at 1500°C for 2h was higher than that of spinels prepared from oxide mixture, hydroxide mixture and coprecipitated materials reported previously and its maximum density was 3.47g/cm³, 97% of theoretical density, with specimen consisting of the powder calcined at 1300°C.

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The elastic modulus of grinding wheel is one of the very important mechanical properties on the grinding performance and theory of grinding. To obtain this value, the direct measurements on grinding wheels or the mathematical methods from the theoretical relation among volume fractions of grains, bonds and pores that are constituents of a grinding wheel and from each elastic moduli of the constituents are tried in the past. No simple and effective explanation was, however, established between the elastic modulus and composition of grinding wheel. In this paper, a simplified structure model was contrived and from this model, was induced the theoretical formula of the elastic modulus. The elastic modulus of grinding wheel is expressed as E, elastic moduli of grains and bonds as Eg and E_b, volume fractions of abrasive grains and bonds in the grinding wheel as V_g and V_b. Then, following relation is obtained: E={1/V_g^1/3E_g+3(1-V_g^1/3)²/V_bE_b}^-1
To verify this formula, test specimens of vitrified grinding wheel were trial-manufactured, based on the grain percentage: 43-51%, grade (hardness): H-R, and grain size: #46-#120 of WA, #60 and #80 of GC, and the elastic moduli of specimens were measured by bending method.
The specimen size is 5(thickness)×12.8(width)×120mm(length). The obtained values were compared with the calculated ones from the formula, and the results showed a good agreement in each combination of the volume fractions of grains and bonds, employing elastic modulus of WA grain E_g: 35.7×10⁵kgf/cm² and of vitrified bond E_b: 6.31×10⁵kgf/cm². The elastic moduli of specimens are in the range of 1-8×10⁵kgf/cm². Therefore, the values of elastic moduli from the simplified model of grinding wheel are inferred to be reasonable.

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Ultramarines have the crystal structure made up of close packings of a sodalite-type cage which is also a main constituent in the structure of Linde A molecular sieves (MS). The formation of ultramarines from MS should be easy from this reason and in fact ultramarine blues were able to be derived from MS5A and MS4A via the following processes (1) to (4); (1) impregnation of MS with Na₂S in the saturated aqueous solution of Na₂S and drying the resulting MS-Na₂S (in N₂ gas), (2) absorption of sulfur vapor on the MS-Na₂S at 500°C (in N₂ gas), (3) heating of the MS-Na₂S-S up to 820°C (in N₂ gas, heating rate 17°C/min), (4) cooling of the MS-Na₂S-S from 820°C to 500°C followed by the air oxidation at 500°C.
The oxidized products from MS5A and MS4A were sky-blue and green, respectively and they turned blue by re-heating in air for a few hours at 820°C. The formation of the ultramarine from MS4A was found much easier than that from the sintered product between kaolin and Na₂CO₃ with the composition of NaAlSiO₄ which has no sodalite-like structure. It was presumed from the above result that the formation of ultramarine green was caused by the intrusion of polysulfides into sodalite-type cages in MS4A.
This method for the synthesis of ultramarine blue is not practical for industrial manufacturing at present because of the high cost of molecular sieves and the low concentration of blue chromophor in the produets, but this method has an advantage for the study of the reaction mechanisms involved in the ultramarine synthesis.

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Transparent and homogeneous coating films of the compositions SiO₂, 98SiO₂⋅2TiO₂, 95SiO₂⋅5TiO₂ and 78.3SiO₂⋅1.7TiO₂⋅20CuO in mol% were prepared by immersing a sodalime glass substrate into metal alkoxide solutions, pulling it up, drying and heating. Effect of preparation conditions on the properties of films was examined. The time for hydrolysis of the solution, that is, interval between preparation of the solution and its application, and the drying time were varied from 1h to 2d and from 1min to 1h, respectively. The films of 0.1-0.3μm in thickness were transparent and homogeneous when the time for hydrolysis of the solution and drying time were short. The thickness of the films could be increased up to 0.5μm by adding a viscosity-increaseing agent to the solution. Defects such as crack, opaqueness or peeling off were observed when the film was thicker than 0.5μm. Scratch strength of the films was higher than that of the glass substrate and increased with increasing heating temperature and heating time. Adhesive strength of the SiO₂ film was higher than 100kg/cm².

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