In the previous report (Tetsu-to-Hagané, 47 (1961) No. 3, p. 390), it was revealed that the cooling of a billet in the mold was largely reduced by formation of the air gap between the billet and the mold in the lower part of the mold.
In this report, the result of experiments which were held in order to search the commencement of formation of the air gap and the thickness of the air gap between the billet and the mold shall be revealed.
The air gap between the billet and the mold occurred when the contraction of the outer shell of the billet by cooling became greater than the creep of the outer shell of the billet by the ferrostatic pressure of the molten steel which was retained yet in the core of the billet.
The contraction of the outer shell of the billet could be calculated by the temperature change of the shell. On the other hand, the creep of the shell could be calculated by the thickness and the temperature change of the shell and the ferrostatic pressure of the molten steel which was retained in the core of the billet. Experiments were carried to search the thickness of the shell by means of the pour-out test and to search the temperature change of the shell by means of inserting the thermocouple into the shell of the billet.
The experimental results were as follows. The commencement of the air gap formation was at 250mm below the meniscus. The taper which occurred between the billet and the mold was found to be 1·66/1000-3·15/1000.
These figures about the air gap between the billet and the mold should have a great meaning n analysing the reasonable cooling of the billet in the mold in the following report.

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Experiments were performed to make clear the true nature of harmful effects of sulfides on mechanical properties and workability etc. of high-sulphur steels such as S-type or S-Pb-type free-cutting steels, and to solve these problems from the results. Various experimental specimens were melted by an atmosphere-controlling high-frequency furnace, and then studies were made on the morphology, deformation of sulfides by working and their effect for the toughness in the length and breadth direction of steels.
The results obtained were as follows; (1) Excess Al or Si in steels elongated sulfides remarkably by working, and this tendency was found also on increase of Mn/S ratio in steel.
(2) There was a large difference between the toughness in the length and breadth direction in such materials, and the toughness in the breadth direction was very inferior to the length direction. (3) On the contrary, decrease of Mn/S ratio or existence of oxygen in steel restrained the elongation of sulfide due to working. (In this case, MnS presumably dissolved FeS or (Fe, Mn) O, or they existed together.) (4) However, the small Mn/S ratio materials was liable to be cracked by hot working. The fact was ascribed to high temperature brittleness. (5) Each problem abovementioned could be solved by a displacement of part of the Mn with the Zr.

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The effect of aluminum nitride (AlN) precipitates on the high-temperature ductility of medium-carbon steel was investigated. Small tensile test pieces (Fig. 1) were made from the steels containing different amounts of aluminum and nitrogen (Table 1 and 2), and were drawn at 500-1000°C.
At first, tensile test pieces were made from the 50kg ingots as cast and were drawn without preheat-treatment.
1) The high-temperature ductility of medium-carbon steels as cast was decreased extremely by the addition of aluminum (Fig. 2 and 3). The steel containing about 0·05% acidsoluble aluminum had the lowest ductility, which was reduced from half to one-fourth of the ductility of the steel containing no aluminum. This phenomenon seemed to be on account that the small AlN particles precipitated dispersedly and pearlite obstruct to deform uniformly, and consepuently the rupture proceeded along the line of AlN particles precipitated at the grain-boundary of austenite.
2) The steels containing more than 0·05% acid-soluble aluminum had higher ductility (Fig. 2 and 3). This phenomenon seemed to be due to the fact that the more the aluminum content, the higher the precipitation temperature of AlN became, and consequently the large AlN particles precipitated at wider intervals.
Next, the effect of AlN. precipitates on the high temperature ductility during cooling after the solution-treatment of AlN precipitates was investigated. The tensile test pieces as cast were heated for 30mn at 1350°C in order to dissolve AlN precipitates. Some of them were cooled slowly (cooling rate 150°C/h) and some rapidly from 1350°C to 1000-500°C, and then drawn.
3) The ductility in austenite (above 800°C) during slow-cooling from 1350°C did not decrease by the addition of aluminum (Fig. 6). This phenomenon seemed to be due to the fact that the AlN particles precipitated during slow-cooling from 1350°C are large and at wide intervals.
4) The ductility in ferrite and pearlite (below 700°C) during slow-cooling from 1350°C decreased by the addition of aluminum (Fig. 6). This phenomenon seemed to be due to the fact that the rupture started from the line of AlN particles precipitated at the center of ferrite-net and proceeds along it.
5) The steel containing about 0·05% acid-soluble aluminum had the lowest ductility below 700°C and the steels containing more than 0·05% acid-soluble aluminum had higher ductility (Fig. 6). This phenomenon seemed to be due to the fact that the more the aluminum content, the higher the precipitation temperature of AlN becomes and the cylindrical AlN particles precipitated at wider intervals, and consequently the brittle rupture became difficult to occur.
6) The ductility during rapid-cooling from 1350°C was decreased by the addition of aluminum. This phenomenon seemed to be dependent on the fact for the line of small AlN particles precipitated along the austenite grain-boundary during cooling.
Next, with the tensile test pieces made from the forged steel, the effect of AlN precipitates on the high-temperature ductility during cooling after solution-treatment was investigated.
7) The ductility during cooling of the forged steels was also influenced by the addition of aluminum like the ductility of the steels as cast.
Finally, the mechanism of panel cracking found in medium-carbon steel ingot was considered.
8) Presumably there was a relation between the phenomenon that the panel cracking of medium-carbon steel ingot was increased by the addition of aluminum and the experiment resulted in that the ductility at high temperature was decreased by the addition of aluminum.

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Influence of melting atmosphere of 18%Cr-8%Ni stainless steels on austenite stability, hardness change due to subzero-treatment, ageing after cold rolling and mechanical properties at room-and high temperatures was studied. Main results obtained were as follows:
(1) The steels NM (0·158%N) and NNM (0·177%N) both melted in nitrogen atmosphere at 600 mm Hg had greater stability of austenite to martensite transformation due to high content of nitrogen compared with the steels melted in vacuum (VM) or in air (AM).
(2) In the steels VM and AM, isothermal martensite transformation was found to occur at temperatures below their Ms point, and in AM even at temperatures above the Ms point.
(3) Remarkable work-hardening caused by rolling at room-and subzero-temperatures, and further hardening caused by subsequent ageing at 300°C were observed with all of the steels except with the steel VM, which consisted largely of martensite in a solution-quenched state.
(4) High-nitrogen steels NM and NNM exhibited lower strength and higher ductility at room temperature than VM and AM in solution-quenched state due to greater austenite stability, while the strength of the former was increased remarkably and the ductility was decreased somewhat by 30%cold-rolling.
(5) Tensile creep-rupture strength of the steels at 700°C was improved considerably with increasing content of nitrogen.

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This report concerned effects of Ni and Mo content on structure of Cr-Ni stainless steels, especially on the amount of δ ferrite and effects of composition and structural conditions on aging hardness, tensile strength and creep-rupture strength.
Specimens used in the experiment were 18Cr-12Ni austenitic stainless steels containing 4%, 6% and 8% Mo (M series) and 18Cr-4Ni, 18Cr-5Ni and 18Cr-6Ni martensitic stainless steels (S series).
With specimens of S series, γ phase precipitated from surrounding portion of δ ferrite and secondnry ferrite precipitated in matrix during aging at 750°C. phase was formed from ferrite after a long-time aging. Specimens were hardened by 5 precipitation of a secondary ferrite and formation of a phase from ferrite.
With specimens of M series, 5 phase was formed quite rapidly from δ ferrite during aging at 750°C.
Tensile strength of specimens of S series depended not on Ni content but on structures.
Tensile strength was increased with raising the amount of δ ferrite and martensite at a test temperature below 500°C.
But strength of δ ferrite was lower than that of austenite above 600°C. Tensile strength of specimens of M series was increased with Mo content.
Creep rupture strength at 650°C was increased with Ni and Mo content but decreased with enhancing the solution-treatment temperature because the amount of δ ferrite was increased with raising the solution-treatment temperature.

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The present investigation was undertaken to develop the low-Mn AISI type 347 stainless steel for nuclear reactor materials. The hot shorthness caused by lowering Mn content in. steel, and favorable effect of Ti or Zr addition on low-Mn alloys.
The experimental results obtained were as follows;
(1) Mn content in steel could be reduced to below 0·10% through careful selection of raw materials.
(2) The hot shortness was apt to occur when Mn content in steel was decreased to 0·60%. Addition of Ti or Zr over 0·10% presumably was effective against the hot shortness.
(3) It can be considered that the drop in hot ductility for low-Mn steels and the solving this problem by the addition of Ti or Zr might not only depend on the type of oxides and oxygen content, but also mainly on the morphology of sulphides.

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A study was made to determine the adequate amount of nickel, chromium and molybdenum enough to obtain the desired hardness in as-cast state in accordance with the wall thickness of the product.
Experimental variables were as follows:
Nickel percent: 0, 1·0, 2·0
Chromium percent: 0·5, 1·0
Molybdenum percent: 0·5, 1·0
Sample size: 30, 45, 60 mm in diam.
Molds: green, dry
To simplify the experiment a “latin-square” design was applied and the results obtained were as follows.
(1) Bainitic structure was obtained when 2% nickel, 1% chromium and 1% molybdenum had been added.
(2) The mixed structure of eutectic cementite and bainite was obtained in the low-silicon range (1·5% Si) when each content of nickel, chromium and molybdenum was 1%, respectively.
(3) From the conclusions mentioned above, the following experimental formula to determine nickel, chromium and molybdenum amount required for 300-600 Brinell hardness number was developed:
Dry mold……HB=114·4X-68·2XY+178·7 (Y+Z) +9·9
Green mold……HB=35·2X+195·OY+58·2XY+195·3Z-10·0,
where X=nickel %,
Y=chromium %, and
Z=molybdenum %.

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