This paper presents the carbon brick-work and its wear during operation of Nos.1, 2, 3, and 4 blast furnaces at Wakayama Steel Works. Carbon bricks in the following parts are quickly worn out;
(1) the part where is inadequately cooled from outside,
(2) the part where carbon brick may be oxidized by charged oxide materials,
(3) the part where carbon brick receives mechanical abrasion by charged materials.
Therefore, the followings are concluded:
(1) Carbon brick should be used for the parts where cooling effect is sufficient, i. e. the lower part of furnace bottom and the peripheral part.
(2) In the upper part of bosh, schamotte brick is rather suitable, while carbon brick can be used at the lower part of bosh.
On the basis of the above conclusion, the brick-work of No4 blast furnace and No1 blast furnace (reconstructed) was designed. Since No4 blast furnace was blown-in (March, 1967), there have not been observed any troubles in brick-work, and the furnace operation is going well.

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Application of carbon blocks to blast furnace bottom initiated in 1951 at Yawata. Since then, troubles of bottom linings have decreased prominently and total iron production of a furnace has increased.
At Yawata, six carbon bottom furnaces have been blown out up to today and properties of used carbon blocks have been investigated. As the results, it is found that carbon blocks absorb considerably large amounts of foreign components, especially molten iron before wear proceeds, resulting in disintegration of brick texture, strength lowering, pore enlargement, crack formation and shrinkage. These property changes were considered to accelerate wearing rate of carbon blocks. From the investigation of used carbon blocks, the wear mechanism of carbon blocks is discussed in relation with carburization, oxidization, floatation and so on. To reduce the degree of property alteration, the authors have studied on improvement of carbon block quality and found anthracite superior to metallurgical coke in high temperature stability and alkali attack resistance. At Yawata, anthracite base carbon blocks are adopted and trouble-free operation is established since 1966.

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Some times, graphite are used for the rotor of viscometers and for the crusible, in which slag is melted, for the viscosity measurement.
The slip between graphite rotor and molten slag was neglegibly small in this work with the circumferential velocity of the rotor of about 0.1-2cm/sec. However, the slip may not be neglegible, when the rotor velocity increased in highly viscous flow.
The viscosity was decreased by addition of graphite powder and the activation energy for viscous flow was increased at higher temperature range.
Therefore, it should be avoided to use graphite for the detecting parts of viscometers or for the crucible when graphite reacts chemically with molten slag or may be dispersed in slag as suspended particles.

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The radioactive lime (⁴⁵CaO) were added in the reducing slag of 60t electric furnace. Radio activity in the macroscopic inclusions extracted from the killed steel ingot with slime method were measured.
The results are as follows.
1) CaO in macroscopic inclusions in killed steel are originated from the reducing slag in electric furnace process by 50-70%. But the ratio of included slag to total compositions of inclusions are 4.5%.
2) The amount of inclusion from slag in molten steel are reduced by 40% during holding the ladle in 10min after tapping.

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The solidification of rimmed steel and the segregation of the elements in the solidified steel are described and discussed.
Experimental results show that the primary blowholes created in the solidifing ingot are very small compared with normal rimmed steel ingot. It will be due to the fact that molten steel of this work are poured more slowly than in ordinary operation. During the solidification with rimming, carbon and sulphur are most remarkably segregated, and phosphorus, not so much as them, white manganese is hardly segregated.
From the autoradiography of ¹⁹⁸Au on the cross-section of ingots, it is observed that the solid-liquid interface during the rimming period is very smooth and solidification velocity of liquid steel is reduced from the bottom to the top of ingot. And also the solidification velocity is decreased by the existence of blowholes too. The experimental equations of the solidification velocity are as follows:
A_m…x=2.15 (√t-0.37)
B_m…x=2.22 (√t-0.35)
C_m…x=2.28 (√t-0.32)
D_m…x=2.00 (√t-0.25)
Symbols A, B, C and D show the portions corresponding to 1/5, 2/5, 3/5 and 4/5 in the height of 10 t flat steel ingot (1800mm) respectively, and subscripts u, m and l show upper, middle and lower parts in each portion respectively, x is the thickness of solidified steel (cm) and t, time (min). The thickness of the impure layer ahead of the solid-liquid interface, which, in generally, depends on the solidification velocity of liquid steel, is decreased from the bottom to the top of ingots where the rimming action occurs more violently.
Moreover, by the fundamental consideration based on “BURTON'S theory”, it is clarified that the CO evolution from the rim-layer during solidification is more intense in low-manganese content ingot than in high-manganese one, and also that the amount of CO evolved is, in either kinds of ingot, much more at the top of ingot than the bottom. This result is very interesting in relating with the powerfulness of rimming action, the degree of deoxidation and the static pressure of liquid steel.

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A study has been made of sulphur segregation in 22t and 33t ingots. The former has 2800mm in height, and the latter has 1030mm in thickness, both of which are almost the limiting sizes for their blooming mill capacity.
The results reported here are summarized as follows:
i) Position of the maximum segregation is raised up towards top side in a longer ingot, and is lowered down in a thicker ingot.
ii) The maximum segregation of sulphur in an ingot may not be increased even in the 33t ingot, but the analysis from center sheets indicates that the 33 t ingot shows higher segregation than ordinary 17t ingot.
iii) The amount of sulphur segregation is, however, determined predominantly by the ladle analysis. The authors also pointed out that the maximum segregation ratio [(max. S-ladle S)×100/ladle S] is increased with increasing the sulfur content in the ladle.
iv) Those results mentioned above are discussed, referring to the experimental data on solidification in the core zone of ingot.

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Structural changes during austenitizing of a Ni-Cr-Mo-V steel, which has martensitic structure initially, were observed.
Austenite nucleated at A_c1 temperature grows with orientation, therefore, shape of residual ferrite becomes acicular. The final structure includes acicular pattern derived from acicular shaped residual ferrite in austenitic matrix. This is the structure that have previously been named “DEC”. Variations of heat-treating condition induce the changes of distribution of acicular pattern.
Equiaxed grains formed in DEC, which have been named “LEC”, increase with increasing austenitizing temperature, and the increasing rate is smaller in the early stage of its formation and larger in the later stage. Variations of heat-treating condition induce the changes of increasing rate of LEC in the later stage of its formation.
It was shown that, in the case when coarser austenitic grain is formed, relatively rougher distribution of acicular pattern in DEC, larger increasing rate of LEC in later stage of its formation, and larger LEC size are obtained.
It was suggested that examinations on fine austenitic grains which constitute the DEC is important for the explanation of these phenomena.

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Ductility and mass effect in large steel castings were studied on production scale, using experimental castings of Al killed steel and 0.36%C steel of which dimensions are 200×200×650mm, 300∅×650mm, 600×650mm∅, 1000∅×1050mm and 80×140×180mm respectively under both conditions of single normalizing and double normalizing.
The results obtained are summarized as follows:
1) Fine austenitic grains were obtained by double normalizing, compared with single normalizing, in both high Al-contained cast steel and low Al-contained cast steel. Coarsening temperature can considered to be 1050°C for high Al-contained cast steel and 1000°C for low Al-contained cast steel.
2) Mixed grain obtained in the low temperature range in single normalizing is thought to be primary mixed grain influenced by the initial grain at the as-cast state, and that in the higher temperature range is secondary mixed grain caused by mixture of coarse grain and incompletely grown grain. Primary mixed grain in double normalizing is seemed to be caused by great difference between coarse grain of prior mixed grain and new grain formed by subsequent treatment.
3) Ductility of steel castings is generally more significantly improved by double normalizing than by single normalizing, but injured by mixed grains that appear in the begining of coarsening.
4) Concerning to ductility of sound and segregating zone in large steel castings, low temperature of 1000°C-1050°C is effective for diffusion annealing of sound zone, and high temperature of 1100°C forsegregating zone. Consequently, in accordance with ductility level demanded for cast steel, it is industrially necessary to select optimum conditions of diffusion annealing; diffusion annealing should be carried out in the low temperature range when ductility of sound zone is not enough high, but in the high temperature range when ductility of segregating zone is important.
5) Mass effect generated on casting in large steel castings is greatly influenced by C content in this experimental casting alloyed with small amount of Ni and Cr.
6) As for mass effect generated on heat treatment, ductility of 300mm dia casting hardly decreases in even the inside, but it decreases abruptly for inside castings of 600, 1000mm dia.
If factor of mass effect is considered, it is also possible for steel castings to estimate tensile strength nondestructively from chemical compositions and heat treated conditions.

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In our previous investigation of the behavior of oxide inclusions in steels during hot-rolling, the specimens containing uniformly a certain amount of specific oxide inclusions were made using the powder metallurgical technique. These specimens were somewhat more porous than those manufactured by melting. When the specimens with high porosity were heated in the air, complex oxide inclusions were frequently formed in those specimens, because the oxide inclusions inserted were combined with the iron oxides which were formed during heating. This investigation was carried out to know how change would occur on Al₂O₃, TiO₂ and SiO₂ inclusions in the above specimens.
Consequently, Al₂O₃, TiO₂ and SiO₂ changed easily into FeO. Al₂O₃ (Hercynite), 2FeO·TiO₂ (Pseudo-brookite) and 2FeO·SiO₂ (Fayalite) respectively. Particularly, the complex oxide 2FeO·TiO₂was identified by various experiments, since, in our previous experiment, it could not be identified by A. S. T. M. cards. Furthermore, the rates of formation of the complex oxides were measured, by means of the X-ray diffraction technique, on the basis of the diffusion reaction between particles of the powder mixtures: Fe, Fe₂O₃ and simple oxides. The activation energy in the formation of those complex oxides was also calculated and discussed.

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(1) The electrolytic extraction with 5% Na-citrate and 10% NaCl electrolyte and the subsequent treatment for purification were elaborately studied and was established the procedure for the completecollection of large inclusions (over ca.10μ) from the steel specimen.
(2) New methods were established for classifying the extracted inclusions by size, determining the composition of a particle of inclusions and observing the inner structure of inclusions and others.
(3) By utilizing this new method, the distributions of inclusions in the ingot and slab, as well as their size and composition were determined for rimmed, capped and killed steels.

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In the future, with the growing use of laterite are as a raw material for steelmaking, metallic impurities such as Ni and Cr will accumulate increasingly in final products. In more general case the final products may be contaminated with Ni and Cr contained in steel scrap. But there is little information on the effects of small amounts of Ni and Cr on properties of plain carbon steels which are in the greatest demand. From such a viewpoint, the authors investigated the effects of small amounts of Ni and Cr on properties of low, medium and high carbon steels.
This paper is concerned with the effects of Ni and Cr on microstructural factors of low carbon ferrite plus pearlite steels. Studies were made for JIS S15C steel, and microstructural factors, such as pearlite fraction, mean ferrite path, condition for CrN-precipitation and critical cooling rate for bainite formation were investigated. Particularly, the condition for CrN-precipitation was discussed with equilibrium state in mild steel, and the free energy equations for CrN-precipitation were presented for the case that Cr and N coexist with Al and Si.

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The effect of small additions of niobium on the cold work-annealing behavior of low carbon steels was investigated to assess feasibility of application of such process to niobium steels. The base composition of test steels was 0.1%C-(0.15, 0.40)%Si-(0.6, 1.2)%Mn and niobium was added in two levels of 0.03% and 0.06%. After cold reduction of 20-50%, isothermal and isochronal annealing was performed at temperatures between 550 and 700°C.The addition of niobium strongly retarded softening and recrystallization, and resulted in an elongated grain structure after recrystallization. A high manganese content remarkably intensified this retardation effect.
The hardening caused by precipitation of Nb (CN) remained same as in as-hot-rolled condition even after recrystallization. From observations of thin film specimens and analysis of acid-insoluble niobium, a mechanism is proposed to account for the retardation in terms of pre-existing Nb (CN), precipitation of niobium originally in solid solution, and also reversion-re-precipitation of Nb (CN) due to introduction of a high density of preferred sites of precipitation (i. e. dislocations).

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With solution treatment at 800°C, ferritic Fe-5Ni-2Al alloy showed almost the same age hardening property as that of solution treated one in full austenite condition. By utilizing this phenomenon, the precipitation process of ferritic Fe-Ni-Al alloy was investigated.
Two kinds of extra spots besides bcc lattice reflection were observed in the electron diffraction patterns of the aged specimens. The one was the spot due to epitaxial oxide film sometimes produced on the specimen surface, which could be removed by carefull treatment of the specimen. The other kind of spot was the super lattice reflections of ordered precipitates, and using this reflected beam, it wasreaffirmed through dark field image that C_sCl type (B2) precipitates arise in the same orientation as the bcc matrix. In the aged Fe-5Ni-1Al-1.7Ti alloy DO₃ type diffraction pattern was obtained, but not the case in Fe-5Ni-2Al alloy aged at 550°C or above. By measurements of electric resistivity andlattice parameter of a -solid solution parallel with observation of electron diffraction pattern, it was assumed that the precipitation of NiAl occurred from the beginning of aging, e. g., at 550°C.

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An examination of fracture surface of commercial high tensile strength steel and high carbon steel specimens broken in low cycle, high stress, fatigue has revealed the occurrence of fracture patterns similar in appearance to those resulting from high cycle, low stress, fatigue, these patterns seem to reveal the fundamental nature of fatigue in ferritic material itself, but do not coincide with “striation” usually obtained in Aluminum and its alloys. A tentative explanation for the fatigue fracture pattern is offered on the basis of recently reported investigations.

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