Some properties of WC–NbC–10%Co cemented carbides have been studied in relation to carbon and niobium carbide contents. The measured values of their properties were examined in comparison with the results so far obtained by the present authors for various cemented carbides. The experimental results are as follows.
(1) Niobium carbide is insoluble in the binder phase, as in the case of tantalum or titanium carbide in WC–TaC–Co, WC–TiC–Co or WC–Ti(Ta)C–Co alloys. (2) The amount of tungsten carbide dissolved in niobium cardide is about 18%. (3) The composition of the binder phase and the carbon content of the niobium carbide (NbC–WC double carbide) phase vary regularly in the three-phase region due to the slight change in carbon content. This resulted in regular changes in other various properties, as has been confirmed in the above various alloys. (4) The three-phase region extends with increasing niobium carbide content owing to the changes in the carbon content of the niobium carbide phase. (5) The phase diagram of the WC–NbC–10%Co alloy system has been established. (6) The WC–NbC–Co material appears to occupy a position intermediate between WC–TaC–Co and WC–TiC–Co materials.

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The change in the electronic specific heat of dilute alloys, especially iron with a small addition of solute elements, was investigated experimentally in order to understand the electronic band structure of transition elements in relation to the origin of ferromagnetism. Since we have to measure a small change (order of one or two percent) in the electronic specific heat with a small addition of the solute element, a special calorimeter, which is used to measure simultaneously the specific heat of three samples, one pure metal and two alloys, has been developed for this purpose. Using this calorimetry, the change of γ can be measured with higher accuracy in one order of magnitude. The solute elements for iron-based alloys are: Ti, V, Cr, Mn, Co, Ni, Nb, Mo, W, Si and Al. Silver-based alloys with gold are also measured to check the accuracy of this calorimeter. The variation of γ for Ag(Au) alloys is less than experimental uncertainty, while the variation of the Debye temperature shows a linear decrease with gold concentration. For iron-based alloys, out of the eleven solutes, six decrease the γ values linearly with solute concentration, c. Their values of (1⁄γ)Δγ⁄Δc are: Fe(Ti), −1.0; Fe(V), −2.2; Fe(Cr), −2.0; Fe(Co), −0.6; Fe(Mo), −0.8; and Fe(W), −2.4. Four solutes increase the γ-values nonlinearly with respect to the solute concentration. The average values of (1⁄γ)Δγ⁄Δc near c=0 are: Fe(Al), +1.2; Fe(Si), +0.6; Fe(Mn), +2.0 and Fe(Ni), +2.6. These data are utilized to discuss the applicability of the rigid band model as well as to correlate the changes with those of magnetic properties.

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Optimum conditions of the annealing temperature, atmosphere, applied field, etc. for the magnetic field annealing of transformer steel sheets are presented from the practical point of view. The specimens used were randomly oriented hot-rolled commercial sheets of T-class containing about 4.5% Si. Excellent magnetic properties were obtained, when high grade sheets including T-90 were preliminarily hydrogen-annealed at 800°C and then annealed in a magnetic field of several Oersteds at 550°∼650°C and when low grade sheets including T-135 were hydrogen-annealed at 1250°C and then magnetically annealed under the same condition. The maximum permeability thus obtained is about 30000 for the high grade sheets, and the uniaxial anisotropy induced has been found to be 980 erg/cc from the calculation based on magnetization curves. The temperature dependence of the isothermal magnetic field annealing effect expressed as the percentage increase of the maximum permeability is linear when plotted against the inverse of the annealing temperature and can be interpreted as a process with the activation energy of 51 kcal/mol, which is in good agreement with the activation energy for diffusion of Si atoms in ferrite.

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This study has been carried out in order to clarify the interactions between metals and molten salts. Optical absorption spectra from 440 to 1000 mμ were measured for dilute solutions of molten Pb–BiCl₃ mixtures up to 0.2 mol% in solute lead metal at the path length as short as 1 mm and at temperatures of 250°, 300°, 350° and 400°C. It was found that the visible spectrum of these solutions consisted of a large and broad band with a maximum near 570 mμ at every temperature. The formal extinction coefficient, which was defined by Boston and Smith, varied with metal concentration. Clear deviations from Beer’s law were found, which shows that these spectra have the characteristics expected for solutions with more than two light absorbing solute species. These results were compared with those of the Bi–BiCl₃ system and of electrical conductivities of the solution of the Pb–BiCl₃ system.

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As the metal-to-ceramic seal assemblies are usually used as part of a vacuum envelope, so the vacuum tightness seems to be important rather than the seal strength in the metal-to-ceramic seals. The brazing with Cu–Ag eutectic alloy was investigated from a viewpoint of vacuum tightness by means of the heat cycle test.
The results obtained are summarized as follows:
(1) Although the seal strength as brazed is scarcely influenced by brazing temperature, holding time and plating material, it is obvious from the results of the heat cycle test that the higher brazing temperature and the longer holding time give a better seal and Ni plating is superior to Cu plating.
(2) The vacuum tightness should be considered to be distinguished from the seal strength as they are independent essentially.
(3) Leak paths are usually formed not in the metallized region but in the brazed layer, especially along the plate-braze interface on either side of the Kovar or the metallized region.

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A study was carried out to obtain a detailed knowledge about the change in microstructures in connection with the secondary hardening on tempering of a series of vacuum-melted 0.2% carbon steels containing vanadium up to about 0.5%. Specimens were austenitized for 2 hr at 1200°C, quenched into 10% iced-brine, and tempered for 1∼100 hr at various temperatures ranging from 150° to 700°C, and examined by a transmission electron microscope.
The main results are as follows: (1) High resistance for tempering in vanadium steels can be explained in terms of the suppression of dislocation climb and of the reduction of the growth rate of ferrite grains by vanadium atoms in solution as well as finely dispersed V₄C₃. (2) Considerable secondary hardening occurs above 550°C, and the hardness reaches a peak by tempering for 1 hr at 625°C in steels containing more than 0.1% vanadium. At the initial stage of secondary hardening, extremely fine V₄C₃ particles less than 20 Å in diameter are preferentially formed on dislocations. These particles are coherent with the ferrite matrix, and give rise to remarkable strengthening. (3) Vanadium carbide V₄C₃ grows into platelets parallel to {100}_α−Fe planes by tempering for 1∼10 hr at 700°C, and then the platelets spheroidize gradually at subboundaries. (4) The orientation relationship between V₄C₃ and the ferrite matrix is similar to that suggested by Baker and Nutting. Excellent lattice coherency is expected along {100}_α−Fe planes from the orientation relationship.
To summarize, the high strength of vanadium steels on tempering is attributed to the finely dispersed precipitates of coherent V₄C₃, a comparatively high density of retained dislocations, and the smallness of grain size. The fine dispersion of V₄C₃ particles can be attributed largely to the existence of high density of dislocations acting as the preferential nucleation sites.

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The dynamic deformation behaviour of zinc single crystals (99.99% purity) with various orientations has been investigated by compression tests using mainly a bar-bar type impulsive loading apparatus and compared with the static deformation behaviour. The strain rates are 10²∼10³ sec⁻¹ and the testing temperature is room temperature. Especially, for those specimens which are deformed by basal slip the strain rates are varied from 10⁻⁴ to 10³ sec⁻¹ and the temperature from −196° to 100°C. The specimens whose axes are nearly in parallel with the c-axis are deformed by twinning and the density of twins is much larger, their size being much smaller for the dynamically deformed specimens than for the statically deformed ones. The specimens whose axes are perpendicular to the c-axis are deformed by {11\bar22} pyramidal slip. The stress-strain curves for the dynamically deformed specimens show yield drops and their initial work hardening rates are lower than those for the statically deformed ones. The specimens whose axes are inclined at 30°∼80° against the c-axis are deformed by basal slip. The flow stress for basal slip scarcely varies for the strain rate ranging from 10⁻⁴ to 10⁻¹ sec⁻¹, but it increases rapidly with increasing strain rate above 10 sec⁻¹. From this result it is concluded that for the basal slip deformation in zinc single crystals at high strain rates the influence of the frictional resistance against moving dislocations becomes too large to be negligible and controls the relation between the strain rate and the flow stress.

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A Fe-0.75% Mn alloy was creep tested at 500°C to examine the following three points: (1) whether or not the steady state creep of the alloy was in a thermodynamical steady state, (2) which configuration of the dislocation network was contributing to the creep strength, and (3) which change in the dislocation network was the cause of the creep recovery. Specimens were furnace-cooled under stress and then observed by transmission electron microscopy. The creep stress was reduced during the test and the recovery of the creep rate was found to consist of two stages. The decrease in dislocation density was not connected with the rapid and extensive first stage recovery, but with the second stage recovery. The subgrain size changed during steady state creep, indicating that the subgrain boundary was not controlling the creep strength. Though not by a single stage, the recovery was completed; i.e. the history effect disappeared in the true steady state. The structure factor, therefore, did not exist as an independent variable during the true steady state creep of the present simple and stable alloy. This suggests that the thermodynamical steady state was realized at least for the factors controlling the creep strength.

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Magnetic properties of Pt–Co alloys near the equiatomic composition have been investigated in detail. Double aging, consisting of the first aging at an intermediate temperature (680°∼720°C) and the second aging at a lower temperature (600°C), was found to very effectively increase the coerctive force, and it yielded the maximum energy product (B·H)_max of 12.5×10⁶ G·Oe.

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The corrosion behaviors of AISI 304 and 304 L stainless steels were studied to clarify the effects of temperature, surface finish and dissolved oxygen in high-temperature water and superheated steam. In deaerated water, the steels indicated a maximum corrosion rate between 250° and 350°C, and the corrosion rate was affected largely by the surface finish. In aerated water, the corrosion rate increased slightly with increasing temperature and was not so much affected by surface finish. In superheated steam, the corrosion rate increased largely with increasing temperature and was affected markedly by the surface finish at higher temperatures. The corrosion films produced in deaerated water and superheated steam consisted of magnetite and chromite, and the chromite content increased with increasing temperature, whereas those produced in aerated water consisted of γ Fe₂O₃. The α phase transformed from the austenite phase by cold-working or surface abrading decreased the corrosion resistance in deaerated water, but it greatly increased the corrosion resistance in superheated steam.

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It has been shown experimentally that the dislocation recovery in pure iron deformed either at liquid nitrogen temperature or at room temperature for about 5% takes place above 200°C, and that the activation energy of the process depends on the internal stress. But the dislocation mechanism of this process has not so far been discussed. In this paper two dislocation models are proposed to explain the change in activation energy of the recovery process mentioned above; climb of the dislocation at the tip of the piled-up group and climb of the dipole dislocations, both of which were induced by internal stress. In case of the pile-up model, the activation energy, U₀, is expressed as U₀=U−0.71 n σ_a b³, where U is the activation energy of self-diffusion, n is the number of loops in the pile-up, σ_a is the average internal stress, and b is Burgers vector. In case of the dipole model, U₀=U−0.32 μ b⁴⁄r, where μ is shear modulus and r is the distance between the dislocations in a dipole. Of these two the dipole model is tacitly considered to be applicable to the recovery process after low temperature deformation as well as after room temperature deformation. The pile-up of dislocations is considered to play only an auxiliary role in th the recovery after low temperature deformation.

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