The operation of a slip system whose Schmid factor with respect to the transverse direction is negative will give rise to a lateral spread, while in the positive case a lateral contraction will occur. Consequently, when slip on one or several slip systems causes the lateral spread, systems which serve to deny it must operate with them. In this paper, the behaviors of slip rotations for cubically aligned copper during rolling are discussed in detail. The results obtained are as follows : A polycrystalline aggregate initially having a (001) [100] orientation rotated first toward (011) [100]. Since the slip systems leading to this rotation brought about the lateral spread, the system serving to check it began to operate on the way to (011) [100]. Consequently the resultant slip rotation took place and contributed to the initiation of displacement toward the stable end position.

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A theoretical analysis of rotations during rolling due to slip on systems of the type {111}⟨110⟩ predicts the occurrence of two kinds of slip rotation. When a stress axis is oriented for single slip, it will rotate toward the boundary where two slip systems have the same resolved shear stress. Then, conjugate slip will begin to operate and lead to the development of the rolling texture described as {011}⟨211⟩. On the other hand, when two stress axes lie in the neighborhood of the same {011} great circle, two texture components {112}⟨111⟩ and {011}⟨100⟩ with large spreads will be formed. With increasing rolling reduction, the {011}⟨100⟩ orientation diminishes rapidly due to slip toward {011}⟨211⟩, and the {112}⟨111⟩ orientation tends to rotate toward {011}⟨211⟩ as well. However, the displacement from {112}⟨111⟩ toward {011}⟨211⟩ will be retarted, in case cross-slip should occur extensively. Consequently, for the metals of high stacking fault energy, the pure metal type rolling texture containing the {112}⟨111⟩ minor component will be developed.

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Slip rotations toward the stable end positions are theoretically studied on the standard stereographic projection, assuming that slip systems of the type {011}⟨111⟩, {112}⟨111⟩ and {123}⟨111⟩ operate together. Materials initially oriented so that the direction ⟨112⟩ may lie between the rolling direction and the operative slip direction, will take up primarily the {111}⟨112⟩ orientation. Then, if conjugate-slip on systems of the type {011}⟨111⟩ occurs extensively, the {111}⟨112⟩ texture component will gradually be displaced toward the most stable end orientation {112}⟨110⟩. In case that the direction ⟨110⟩ is situated between the rolling direction and the operative slip direction, the rolling plane normal approaches the great circle {001}-{112}-{111}, while the rolling direction is displaced to ⟨110⟩. Then double-slip on systems of the type {011}⟨111⟩ and {112}⟨111⟩ begins operating, and the two stable end orientations {112}⟨110⟩ and {001}⟨110⟩ will be developed.

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Masumoto et al. discovered in a previous work that Ni–Mn binary alloys containing about 22% Mn developed high permeabilities by a suitable heat treatment. These alloys were named Nimalloy. The discovery was followed by much subsequent work on the effects of additions of various elements on the properties of Ni–Mn alloys, and the highest initial permeability of 76000 and the highest maximum permeability of 441000 were obtained in the case of Cr additions to Ni–Mn–Fe alloys. With increasing additions of Sb and Sn to Ni–Mn alloys, both initial and maximum permeabilities were gradually increased. In the Ni–Mn–Sb system, the highest initial permeability of 10870 and the highest maximum permeability of 80700 were obtained the alloy comprised of 76.43% Ni, 16.27% Mn and 7.30% Sb which was cooled at a rate of 2800°C/hr from 900°C and then reheated at 420°C for 4.5 hr and the alloy of 76.17% Ni, 18.53% Mn and 5.30% Sb which was cooled at a rate of 240°C/hr from 900°C and reheated at 460°C for 2 hr, respectively. In the case of the Ni–Mn–Sn system, the highest initial permeability of 11200 was obtained with the alloy comprised of 77.07% Ni, 17.36% Mn and 5.57% Sn which was cooled at a rate of 240°C/hr from 900°C and then reheated at 440°C for 1 hr, and the highest maximum permeability of 73700 with the alloy of 75.19% Ni, 11.20% Mn and 13.61% Sn which was cooled at a rate of 2800°C/hr from 900°C and then reheated at 380°C for 15 hr, respectively.

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An attempt was made to identify an oxide film formed on the surface of 18 Cr–8 Ni steel, heated at 500°, 800° and 1050°C in air in less than a few minutes. The structures of the oxide film were determined with the aid of an electron microscope equipped with an X-ray microanalyzer, which enables us to obtain electron micrographs, diffraction patterns and X-ray emission spectra from the same area of the film. The oxide film was composed of three oxides; Cr₂O₃, Fe₃O₄ and a manganese oxide (MnO or Mn₃O₄). Ni could not be detected from any oxide films. The mono-metallic oxides of each metal are first formed at the initial stage of oxidation at high temperatures, but not the compound oxides which are formed after a prolonged oxidation.

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An adiabatic calorimeter used for specific heat measurements was improved for the measurements of heats of mixing. Heats of mixing in the binary liquid alloys consisting of silver and aluminum, gallium, indium, thallium, germanium, tin, lead or bismuth were determined at 970°C. The experimental data obtained were represented by the following ξ function.
(Remark: Graphics omitted.)
The main results obtained are as follows.
(1) Some systematic correlations were found between heats of mixing and positions of the considered couples on the periodic table.
(2) Behavior of heats of mixing seems to be explained qualitatively in terms of the electronegativity, size and valence factors. Each factor has an intimate relation with the periodic table.
(3) Heats of mixing have intimate relations with phase diagrams, suggesting that the atomic bonding in liquid silver alloys is similar to that of solid.
(4) The ξ function of silver alloys having secondary solid solutions in a solid phase decreases markedly with increasing silver concentration. Hence it is considered that the interaction energy depends considerably on the silver concentration.

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Chromized steel is produced by the thermal diffusion process of evaporated chromium on steel. Therefore, the chromium contents depend on the depth from the surface of the sample. An attempt has been made to determine the surface chromium contents of the chromized steel by X-ray fluorescent spectrometric analysis. An excitation potential of 6.5 kV was used for all measurements so as to generate CrK_α, but not FeK_α radiation. The sample which has the concentration gradient of chromium, e.g. chromized steel, was not identically treated with a usual uniform sample. In this case, some correction is necessary for the sample density and absorption coefficients, considering the concentration gradient. The following two methods were used for the correction of the surface chromium contents: One was to employ the intensity ratio of CrK_α radiation from chromized steel to that from the uniform sample, and the other was to employ the thickness d_1⁄2 (50% emission) in consideration of the linear gradient of chromium concentration. The chromium contents used for the above corrections were estimated by using the theoretically introduced formula. The concentration gradient of chromium was obtained by measuring the intensities of CrK_α emitted from the electrolytically polished samples. Consequently, it was found that the surface content of chromium should be corrected by about 3% in case of the 0.34% μ chromium gradient.

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To investigate the oxidation behavior of TiC-base alloys containing molybdenum, TiC-10% Ni and TiC-10% Ni-(2.5∼18.57)% Mo alloys were prepared by sintering. These alloys were exposed in still air at 600°∼1300°C. The oxidation process and the effect of molybdenum on oxidation were studied by the weight gain, surface recession, X-ray diffraction and microstructure observations of scales. The following results were obtained.
(1) Addition of molybdenum to TiC-10% Ni alloy increased the weight gain due to oxidation below 1100°C, which was most effective at 700°∼800°C. This effect was, however, decreased at higher temperatures and at 1300°C molybdenum increased the oxidation resistance reversely.
(2) From the X-ray diffraction data, it appeared that at higher temperatures the oxidation of molybdenum was suppressed to cause the appearance of metallic molybdenum in the scales and NiO·TiO₂ was also increased, while at lower temperatures NiO was combined preferentially with NiO·MoO₃ and decreased the NiO·TiO₂ content in the surface layer of scales.
(3) From measurements of the surface recession, it was found that the progress of oxidation of all the alloys at 1000°C was parabolic and the parabolic rate constant increased in proportion to the molybdenum content. But at 1300°C, the surface layer of scales became porous and no parabolic relation was found in all the alloys.

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The yield point and the tensile strength of steels worked at elevated temperatures (about 300°C) are higher than those of steels worked at room temperature or of steels annealed at various temperatures after cold working. It can be considered that strengthening of steel by warm working is attributed to the strengthening of α-iron in steel; the dislocation density of α-iron worked at elevated temperatures is higher than that of α-iron cold worked or annealed after cold working because of the multiplication of dislocations due to dynamic strain aging in warm working. 0.02% C α-iron warm worked, cold worked and annealed after cold working is studied, respectively, by means of transmission microscopy and the hardness test.
The results obtained are as follows:
(1) Both the low temperature annealing effect after cold working and the warm working contribute mostly to the strengthening of 0.02% C α-iron in the same way as for other steels, so that the strengthening mechanism of 0.02% C α-iron is considered to be the same as that of other steels.
(2) The strength of the warm-worked steel is higher than that of the steel annealed at various temperatures for various times after cold working.
(3) The dislocation density of α-iron worked at 300°C is higher than that of α-iron worked at other temperatures and annealed after cold working. The dislocation density of α-iron corresponds to the hardness of α-iron.
(4) The mechanism of strengthening α-iron by warm working is based on the increase of the dislocation density by multiplication of dislocations due to dynamic strain aging.

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The structures of Ti-4∼14% Al-2∼8% Co alloys have been examined by optical microscopy, after annealing for a long period in the range of 600°∼1100°C. A partial titanium-aluminium-vanadium ternary diagram for this range has been established.
To distinguish the α and α₂ phases, use was made of the fact that no precipitation from α₂ occurs with decreasing temperature, but α₂ is precipitated from α in the high aluminium alloys.
Ti-6% Al-4% V alloy is situated on the boundary for the formation of the α₂ phase. The increase of aluminium or vanadium contents in this alloy brings about the formation of the α₂ phase.

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A mass spectrometer was applied for vacuum fusion analysis of micro-amounts of gases in iron and steel produced by vacuum melting and vacuum casting. The gases extracted from the metal specimen in a vacuum fusion furnace were collected directely in the gas reservoir through a mercury diffusion pump, and the individual constituents were determined by the mass spectrometer (Gas Detector, ATLAS–WERKE AG.). The limits of detection for this apparatus were 0.07 μg of oxygen, 0.1 μg of nitrogen and 0.005 μg of hydrogen. The analysis took approximately 30 minutes. From the extraction curve with vacuum fusion, it was proved to be extracted gases from iron and steel within 2 minutes.

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