Prior to performance of fundamental experiments on the chemical process utilized the fluidized bed of the dry decopperization of pyrite cinders, it was seemed necessary to determine the fluidization characteristics.
The reaction tube was made of 2.4cm diameter silica tube and a perforated porcelain plate was used as supporting disc for the sample and silica particles packed with object of preventing falling of the sample.
Pressure drops across the bed of pyrite-cinder particles in the silica tube were measured at decreasing air velocities.
Fluidization tests were run on the 100-150 mesh seven pyrite cinders first and then the works were extended to 200-250 mesh cinder and 65-250 mesh mixture of sized containing four components.
The trends of pressure drops vs. air velocities for 100-150 mesh samples were retraced with good reproducibility below the temperatures in the range from 950 to 1050°C, while at this temperature range the abrupt changes of pressure drops occurred with decreasing velocities, the fluidization ceased and the air-blew through the center of the bed. Other sizes gave similar results.

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The equilibrium of carbon monoxide-carbon dioxide mixtures with carbon dissolved in liquid iron alloys has been studied at 1560°C and represented by the equation:
C (in liquid iron or alloy)+CO2(gas)=2CO(gas)
The results show fairly good agreement with the previous works on the iron-cabon system by Richardson and Dennis (F.D. Richardson, W.E. Dennis, Trans. Faraday Soc., 1953 Vol. 49, pp. 171-180) and by Rist and Chipman (A. Rist, J. Chipman, Rev. de Metall. 1926, 53, pp. 1-12)
Experimental difficulties arising from the deposition of carbon by decomposition of carbon monoxide have required restriction of the esperiments to low-carbon alloys (0.1-0.3%C).
The relationship between the activity of carbon and the concentration of carbon and alloy elements, j, are expressed by the parameters equivalent to the effect of one percent of the alloy on the logarithm of the activity coefficient of carbon as follows:
C+0.200 Ni+0.012 Co+0.012 W-0.003 Mo-0.009

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The total amounts of carbides in the several heat-treated specimens of the high-W high speed steels and Mo high speed steels were determined by electrolytic isolation method in this 2nd report like thd 1st (Tetsu-to-Hagane, 1958, Vol. 44, p. 1186).
Chemical compositions and crystalline structures of the isolated carbides were determined by chemical analysis and X-ray diffraction, and the shapes of the carbides were observed by an electron microscope, and the following results were obtained:
(1) On the as-annealed states, the amounts of carbides in high-W high speed steel were about 30%, and in Mo high speed steel were about 23%, and about 95% of the W, 80% of the V in the steels, and 90% of the Mo in the Mo high speed steel were concentrated in the carbides, while the most portions of the Si and Mn were dissolved in the matrix.
The carbides precipitated in these annealed steels were M₆C, MC and M₂₃C₆, and the most portions of the carbides were M₆C.
(2) In quenched states, the amounts of insoluble carbides in these high speed steels were 10-14%, which was the same about 40-45% of the carbides in as-annealed state.
The insoluble carbides in quenched high-W high speed steels were only M₆C, but the carbide in quenched Mo high speed steel consisted of M₆C and MC.
(3) The higher tempering temperature, the more precipitated the carbides from the austenite. And, especially, the concentration percentage of the special alloy elements in the carbides increased suddenly by the tempering to 575°C.
The M₂₃C₆ carbide, precipitated in the high-W high speed steel by the tempering to 750°C, and the amount of MC carbides in the Mo high speed steel increased by the tempering, and M₂₃C₆ carbide appeared by the tempering to above 575°C.

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High Cr-Fe alloys absorb nitrogen readily when they are heated in the gas at high temperatures. Although such absorbed nitrogen has very interesting effects on the alloys, few studies on this subject have been reported. This paper dealt with the influence of carbon content of 20% Cr-Fe alloys containing 0.04% C, 0.14% C or 0.22% C on the austenite formation due to the nitrogen absorption of these alloys, and clarified the behavior of the formed austenite due to various heat treatments.
The results obtained were as follows:
(1) The absorption of nitrogen during the heating for 4-6 hours in the extremely pure nitrogen, which was predeoxidized and predehydrated by metallic sodium chips, at 1250°C changed the surface zone of 20% Cr-Fe alloys into single austenite phase at the temperatures, and the depth of the austenite zone was enhanced as carbon content of the alloy increased from 0.04% to 0.22%. In inner region of the alloys, a duplex structure of austenite and ferrite was produced by the less content of nitrogen. The higher the carbon content of the alloy was, the more the amount of austenite became.
(2) 20% Cr-Fe alloys heated for 6 hours in the nitrogen atmosphere at 1250°C was found to contain nitrogen from about 0.3 to 0.45% in the single austenite zone at the surface, and the nitrogen content decreased at the inner zone from the surface. At the higher carbon content of the alloys, however, the nitrogen content of these zones did not greatly concern the depth from the surface. As regards the carbon content, surface zone of the single austenite was likely to have fairly higher carbon level than that in the inner zone consisting of a mixture of both austenite and ferrite, which was presumed to be caused by the diffusion of carbon from the inner zone to the outer surface during the nitrogen absorption.
(3) The austenite formed by the nitrogen absorption decomposed at about 850°C during the furnace cooling from 1050°C. Two stages of transformation were observed to occur during the air cooling of the nitrogen-absorbed alloys from 1050°C, of which one took place at about 650°C and the other at about 150°C. As regards the former, the lower the carbon content of the alloy was, the larger the change became. While, the latter stage became larger as the carbon content of the alloy increased. It was also found that the temperature range in which the change from ferrite to austenite occurred on heating was about 900°C-1050°C.
(4) The austenite retained by quenching in the single austenite zone was fairly stable in the heating at temperatures below about 600°C, but it tended to decompose by the temperatures above this, and broke down completely by the tempering for 30 minutes at 650°C or thereabouts.
(5) Such retained austenite was sensible to the subzoro-treatment considerably, but higher carbon alloys were less sensible. In the same alloy, the sensibility was less at the outer surface. While, the holding at room temperature after quenching was found to stabilize the retained austenite. This effect was stronger in the outer zone of the same alloy and with the higher-carbon alloys.

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Effects of practical heat treatment shown by Table 1 on mechanical properties of Inconel X-550, Inco 73°, Nimonic 90 and Inco 700 were investigated.
The heat treatment B which contained water cooling after solution-treatment gave higher short-time tensile strength. But in the case of Inco 739 and Inco 700 which contained higher Al than other alloys, the heat treatment F gave the same short time tensile strength at 750°C as the heat treatment B.
The heat treatment of D and F, which contained solution treatment at 1180°C, gave the highest stress rupture strength at 750 and 816°C to Inco 739, Inco 700 and Inconel X-550.
The reason of the above fact was due to softening during long time test at elevated temperature in the case of heat treatment B.
In the same way as Inco 700 heat-treated by D did not show softening during stress rupture test at 816°C, so heat treatment D showed the highest stress rupture strength at this temperature. But as in the case of other alloys, hardness of specimens decreased remarkably by overaging during stress rupture test at 816°C, so effect of heat treatment on stress rupture strength at this temperature were not clearly.
The heat treatment C and E, which contained coarse aging at 980°C, gave lower shorttime tensile strength and stress rupture strength. It was due to the fact that aging at 980°C cause coarse precipitates, and was softened easily at elevated temperature by overaging.

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