The tensile and recrystallization properties of electron beam melted molybdenum wires were studied comparing with those of arc melted and sintered molybdenum. The recrystallization temperature range of electron beam melted molybdenum wires was 930°∼1090°C, while those of arc melted and sintered were 820°∼1070°C and 740°∼1260°C respectively. A linear relationship was observed between yield strength and 1⁄T for three kinds of molybdenum wires. The temperature range of the linear relationship was 140°∼−70°C for electron beam melted molybdenum wires, 160°∼80°C for arc melted molybdenum wires and 180°∼below −110°C for sintered molybdenum wires. The parameters for the yield strength vs 1⁄T diagrams were 23.5×10³ kg/mm⁻²T°K, 27.0×10³ kg/mm⁻²T°K and 30.0×10³ kg/mm⁻²T°K for the electron beam melted, arc melted and sintered molybdenum wires, respectively.
The ductile-to-brittle transition temperatures in the tensile test (with strain rate of 2.1×10⁻³ s⁻¹) for recrystallized electron beam melted and arc melted molybdenum wires were −70°C and −80°C, but the transition temperatures varied very sensitively with changes in the annealing temperature.
Sharp peaks were observed at about 200°C in strain hardening exponent vs the reverse of test temperature diagrams for 3 molybdenum wires, and the strain hardening exponents of electron beam melted molybdenum wires were the highest at all the test temperatures and those of arc melted were the lowest.

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The age-hardening of the Cu-7%Al-1.5%Co alloy containing 0.01%P has been experimented by a variety of methods. The main results are as follows:
(1) The age-hardening of the alloy, solution-treated by heating at 970°C for 2 hr and quenching in water, takes place when it reheated above 350°C. The behavior of hardness in the course of the age-hardening is such that is observed with usual precipitation hardening alloys.
(2) The activation energy for the the hardening is 37.3 kcal/mol, being almost equivalent to that for the diffusion of Al in Cu.
(3) The process of the age-hardening was examined by electron micrography, X-ray diffraction, electron probe micro-analysis. The results imply that the hardening is ascribed to the formation of G·P zones, where Al and Co are concentrated. At later stages of the heat treatment the zones grow into precipitates, whose electron microdiffraction indicates that they have a cubic structure.
(4) Between 500°C and 600°C a heat evolution due to the formation of G·P zones is observed by the thermal analysis of the solution-treated specimens, and the evolved heat is determined to be 5.5 cal/g. Below 300°C another heat evolution is observed in the water quenched specimens, but the latter heat evolution seems to do nothing with the age-hardening of this alloy.

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The influence of the addition of small amounts of B, Zr and Cr on grain size, age-hardenability, grain boundary reaction, mechanical properties and electric conductivity of Cu-4 wt%Ti alloys has been investigated. The alloy shows the minimum hardness at the solution annealing temperatures of 800°∼850°C and its hardness is increased somewhat by annealing at 900°C. The solution annealing for 5 min is sufficient for the softening. The addition of B, Zr and Crhas the capacity of grain refinement. It seems that the effect of B is especially due to its presence as an inclusion in the alloy. Zr and Cr are found to relax the hardness drop during long aging above 500°C. The effect of Zr is excellent, and almost inhibits the hardness drop during aging at higher temperatures such as 550°C. The grain boundary reaction has disappeared by the addition of 0.18 wt%Zr. The addition of Cr is also capable of supressing the grain boundary reaction, but its effect is less than that of Zr. It seems that B as an inclusion forms nuclei for the grain boundary reaction and promotes the reaction. Any addition of B, Zr or Cr improves mechanical properties of Cu-4 wt%Ti alloy but somewhat reduces the electric conductivity. The Kb value and Young’s modulus have maximum values by the addition of B 0.02 wt%, Zr 0.18 wt% and Cr 0.76 wt%, respectively.

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The effects of the addition of small amounts of Fe and Sn on the grain size, precipitation at grain boundaries, electric conductivity and mechanical properties of age-hardenable Cu-4 wt%Ti alloy were investigated.
This alloy showed a minimum hardness at the solution-treatment temperature above 750°C when 0.95 wt%Fe was added and also at 800°∼850°C when 0.2∼1.0 wt%Sn was added. The duration of solution annealing required for sufficient softening was 5 min as indicated in a previous work. By X-ray analysis, it was confirmed that the spherical particle observed in the Cu-Ti alloy containing Sn corresponded to a compound of Sn and Ti.
It was assumed that the addition of Fe to the binary alloy would give rise to the formation of a compound with Ti. Fe or Sn is capable of refinement of the alloy, the effect of Fe being more prominent. The addition of 0.95 wt%Fe was fonnd to suppress the hardness drop during a long time aging above 500°C. The rate of the precipitation on grain boundary was increased by the addition of a small amount of Sn, but decreased above 0.5 wt%Sn. When 0.95 wt%Fe or 0.5 wt%Sn was added to the binary alloy, the Kb value (Federbiegegrenze: DIN), Young’s modulus and tensile strength attained maximum values, respectively.
From the result of both the previous and present studies, it has been confirmed that the addition of \simeq0.5 wt%Fe and 0.1∼1.2 wt%Zr is most effective fort he refinement of grains and the suppression of precipitation at grain boundaries, and that the Cu-0.69 wt%Fe-0.10 wt%Zr-4.22 wt%Ti alloy has improved properties as expected.

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Properties of two phase WC-Co alloys are not identical but vary remarkably. This is due to the fact that the tungsten content of the binder phase varies extremely from 2∼3% min to 9∼10% max in the two phase range with a very small variation of the carbon content in the alloys. Moreover, properties of the two phase alloys, when their carbon contents are lower, vary abrubtly by low temperature annealing. For example, the transverse-rupture strength, magnetic saturation, electrical resistivity, and lattice parameter of the binder phase decrease, while the hardness of the binder phase increases. Such aging phenomena in cemented carbides are caused by the precipitation of tungsten atoms from the binder phase.
The results mentioned above were fully described in the previous papers^(1)∼(5), following which the present study was carried out on the aging mechanism of cemented carbides, i. e., the mechanism of precipitation of tungsten atoms when low carbon alloys were annealed. The specimens used were two phase WC-Co alloys vacuum-sintered at 1400°C for 1 hr and subsequently vacuum-annealed up to 30 hr at temperature below 950°C. The results are summarized as follows:
(1) Tungsten atoms supersaturated in the binder phase precipitate in the form of CO₃W according to the reaction γ (f.c.c.)→γ (f.c.c)+Co₃W. (2) The precipitate of Co₃W is formed on (001) cubic planes of the face-centered binder phase. (3) The hardness increase of the binder phase is considered due to a precipitation hardening by Co₃W. (4) The decrease in magnetic saturation is also attributed to the precipitation of Co₃W presumed to be nonmagnetic. (5) The solid solubility of tungsten in the cobalt phase is determined. The strength of low carbon alloys seems to be considerably improved by solution-treatment. But, for the purpose of obtaining high strength alloys, the preparation of high carbon alloys are far more desirable than the solution-treatment of low carbon alloys.

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The present authors have studied the structure of thin oxide films formed on the Zr surface under various conditions. The following results were obtained by electron microscopy and electron diffraction. (1) The oxide films formed by dry oxidation, e.g., heating in air at 300°∼500°C or dipping in a fused salt bath (NaNO₃ 50 g, KNO₃ 50 g, MnO₂ 1 g: 400°C), consisted of only the monoclinic ZrO₂. (2) A number of regularly aligned lines of about 2500 Å in length was observed in some oxide films formed on the surfaces smoothed by cloth-polishing. (3) The tetragonal ZrO₂ as well as the monoclinic ZrO₂ was detected in the Zr oxide films produced by wet oxidation such as immersing in boiling conc. HNO₃. (4) On the hydrogen-absorbed surfaces by either cathodizing in 5%H₂SO₄ or immersing in 35%HCl, the oxide films composed of the tetragonal and the monoclinic ZrO₂ crystals were produced by dry oxidation. (5) When the Zr specimen was quenched in water after heating in air at 400°C, the surface was covered with both the tetragonal and the monoclinic ZrO₂ crystals. (6) The tetragonal ZrO₂ in the oxide films produced by wet oxidation transformed to the monoclinic ZrO₂ by heating in air at above 1150°C. (7) These results indicate that the meta-stable tetragonal ZrO₂ crystals are produced in the oxide films partly due to the existence of water or hydrogen in oxidizing atmosphere.

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Arc fusion-condensation vaporization method in an inert gas stream was applied to the determination of oxygen in iron and steel.
The flow of argon as a carrier gas was controlled to be 500 mL/min.
The degassing action in the furnace was made by arc discharge (current 30 A), the blank value of oxygen being reduced to 15 μg by 2 min discharge. The arc discharge for the extraction of gas from a sample was sufficient with the consumption of the current of 20 A for 1 min.
The precision of oxygen analysis by this method is very high and its reliability is almost equiavalent to that of the vacuum fusion method.
The time of the present analysis from the charge of the sample (approx. 1 g) to its completion is less than 6 min.

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In order to carry out the quantative analysis yttrium in steel, a spectrographic method was applied. The solution technique with a rotating disk electrode was used, because solid standard samples for the spectrographic analysis of yttrium were not prepared and the use chipped samples was intended.
Using the Carl-Zeiss QU 24 spectrograph, effects of discharge conditions, dissolving acids (HCl, HNO₃, H₂SO₄) and electrodes (graphite, Cu) on the sensitivity and precision of the yttrium spectral line were studied.
The result showed the optimum analytical condition to be as follows:
(This article is not displayable. Please see full text pdf.)
The reproducibility of the solution method was 3.9% in terms of the coefficient of variance in the range from 0.01%Y to 0.7%Y.
The Ti content is necessary to be less than 0.01%, because Y 3242.3 Å is affected by Ti 3242.0 Å.
If the Ti content is more than 0.01%, another line pair Fe 3659.5/Y 3664.6 can be applied. In this case, however, the reproducibility is worse in some measure than that of Fe 3205.4/Y 3242.3.
Then, as solid steel samples which contain yttrium were obtained, the content of yttrium was determined by the solution spectrographic method mentioned above and then direct excitation method of solid steel was studied by using those samples for standardization.
It was found that the same excitation conditions with those of the solution spectrographic method was also available favorably in this case, and the sensitivity and precision were almost same to those of the solution spectrographic method.

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Cu-Ni-Zn alloy shows an anomalous increase in hardness by low temperature annealing: In this investigation, the structure of the anomaly was discussed in detail using a 18 wt%Ni alloy. Changes in hardness, elastic modulus and internal friction were measured at room temperature after water-quenching from several temperatures during isochronal and isothermal annealing. (1) The hardness of a specimen furnace-cooled from 700°C was greater than that of a specimen water-quenched from the same temperature. Both of the specimens showed a maximum hardness at about 300°C and it decreased gradualy with further increase of temperature during isochronal annealing. (2) During isothermal annealing at 300°C, the hardness of the water-quenched and cold-worked specimens increased with the annealing time, and became greater with the degree of cold-work, the overall change in hardness. Result (1) sugests the hardening by low temperature annealing to be due to short range ordering; result (2), however, could not be intepreted only by a simple sum of the short range order-hardening and the work-hardening, but it meant that another type hardening mechanism promoted by increasing of the dislocation density existed in this alloy. (3) The change in elastic modulus was nearly the same as that in hardness shown in (2). Activation energies calculated from the change of elastic modulus in the cold-worked and water-quenched specimens were 22∼26 kcal/mol and 17 kcal/mol, respectively. The results of the quenched specimen showed that the short range ordering increased the elastic modulus and was improved with quenched-in excess vacancies. The results of the cold-worked specimen, however, could not be interpreted only by the short range ordering, but it might be reasonably conceived that the strengthening of interaction between high density dislocations and solute atoms would increase the elastic modulus and hardness. (4) The strain-amplitude independent internal friction of the quenched specimen showed a minimum value at about 350°C during isochronal annealing indicating the strongest pinning effect of dislocations at this temperature. The interaction between dislocations and solute atoms was known to serve as the pinnig effect.
Therefore, the authers conclud that the mechanism of hardening in the cold-rolled nickel-silver containing about 18 wt%Ni by low temperature annealing depends not only on the short range ordering but also on the interaction between dislocations and solute atoms, i.e., the segregation of solute atoms into dislocation atmosphere.

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An age-hardenable Cu-Ti alloy (Cu Side) has been shown to precipitate pearlitic or widmanstätten β′-Cu₃Ti (or Cu₇Ti₂) under different conditions of aging. Using a Cu-3 wt%Ti alloy, this paper deals with the changes in the shape of precipitates by aging in two different temperature ranges for the pearlitic precipitation and the widmanstätten precipitation after water-quenching from 900°C. (1) Isochronal aging for 13.5 min with a temperatures interval of 50°C in the temperature range between room-temperature and 800°C: The pearlite became visible above 400°C at grain boundaries where two neighbouring grains have a particular orientation relation with each other. Furthermore, the pearlite seemed to grow into one grain with the orientation affected by that of the other grain. In the temperature range between 650° and 750°C, the pearlite still remained, but different type precipitates (see Photo. 7(b), 8(b)(c)) were observed frequently at grain boundaries. These new type of precipitates could be interpreted as follows: Some of the pearlitic precipitates, whose habit plane coincided with {111} planes of surrounding super-saturated α-phase, grew into the α-phase. By the aging above 700°C, widmanstätten precipitates appeared in the super-saturated α-phase coexsistent with the pearlite. (2) Aging at 700°C after preaging at 550°C: When residual super-saturated α-phase remained after the pre-aging, widmanstätten precipitates appeared in the residual α-phase. On the other hand, pearlitic precipitates changed into straight rod-type precipitates (see Photo. 14) when the entire alloy had been changed into the pearlite by the pre-aging. However, the rod-type precipitates, retained traces of pearlitic precipitates and were different from widmanstätten precipitates.

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The effect of the Ms temperature on the lattice defects in martensite was studied by X-ray diffraction. Fe-Ni alloys with 10, 15 and 20 wt%Ni were chosen as the specimens having different Ms temperatures. The profiles of the Debye-Scherrer lines of martensites in three alloys, were taken with an X-ray diffractometer. Correction was made for the instrumental broadening, and the Fourier analysis was performed.
It is concluded from the results that: (1) The integral width of the martensite line decreases with increasing the Ms temperature. (2) The index-dependences of the broadening are the same for all the specimens examined. (3) The internal strain of the 110 direction in the martensite decreases with increase of the Ms temperature, whereas the effective domain size of that direction does not change so much.
The above results show that, even at a Ms temperature much higher than room temperature, considerable amounts of the lattice defect, including the stacking mistakes, are introduced during the martensitic transformation.

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The reaction between aluminum and sodium fluoride was studied by means of differential thermal analysis and X-ray analysis.
It was noticed that the reaction proceeded in considerable degree through the formation of sodium vapor and cryolite by the reaction of molten aluminum and solid sodium fluoride at a reduced pressure above 1020° to 1150°K.
The reaction proceeded as follows: 2NaF_(s)+1/3Al_(l)=Na_(g)+1/3Na₃AlF_6(s,β).
The reaction pressure of sodium versus temperature was determined by the following equation in the temperature range of 1063° to 1153°K. logP_Na(mmHg)=-9,995/T(°K)+10.752.
Thermodynamic consideration was also given to the above reaction.

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The strength of the high temperature boiler alloy AN 31 was investigated by means of a fatigue test and some other mechanical tests at temperatures ranging from 600°C to 750°C. The structural changes associated with the fatigue were also examined by an optical microscope.
The fatigue strength of the AN 31 alloy was much higher than that of the 18-8 austenitic steel and also than the proof stress of this alloy. The fatigue strength was improned by understressing followed by a process of gradually increasing the amplitude of the alternating stress in small increament.
Considerable carbide precipitations along the slip bands and grain boundaries were observed in the fatigued specimens.
A possible explanation for the observed high fatigue strength of this alloy and for the coaxing effect in fatigue may be given as follows: The niobium carbide can precipitate rapidly from the matrix under a plastic deformation under a repeated stress, and then the fine precipitates harden the matrix, causing the grains to resist fatigue damage.
While a sharp peak was observed in the hot hardness vs temperature curve at about 700°C accompanying a marked drop of the impact value, no peak was observed for the specimen which had been fatigued before the hardness test. This implies that most of the solute niobium in the matrix precipitated out during the fatigue test.

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The effect of solution-treatment prior to cold-rolling on the textures of Al-Cu alloys was studied by the X-ray pole figure method. The following results were obtained:
(1) No marked effect of the solution-treatment on the cold-rolling textures of Al-Cu alloys was observed.
(2) The recrystallization textures varied remarkably by the solution-treatment. The recrystallization textures of Al-Cu alloys which had not been solution-treated were found to be composed of two preferred orientations of (100)[001] and (632)[5\bar5\bar7] or (113)[33\bar2]. Whereas the recrystallization textures of Al-Cu alloys which had been solution-treated before rolling consisted of two preferred orientations of (100)[001] and (110)[1\bar10]. The amount of the (100)[001] type component in the recrystallization textures was increased by the solution-treatment before rolling.
(3) The orientation relationships between the preferred orientations in the cold-rolling texture and those in the recrystallization texture can be expressed as ⟨111⟩ rotations of 20° to 40°.
(4) The variation of the recrystallization texture by the solution-treatment prior to cold-rolling may be attributed partly to the change in the distribution of precipitates during the incubation period of the primary recrystallization, which has influence on the selective nucleation or growth of the nuclei, and partly to the suppression of the random nucleation around the precipitated particles.

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In order to investigate the mechanisms of fatigue at elevated temperatures, the processes of fatigue in annealed specimens of 99.99%Al and Pb were observed microscopically at testing temperatures from 0.5 to 0.85 Tm°K under reversed torsional or bending stresses.
The results observed at the early stage of the high temperature fatigue are summarized as follows:
(1) The grain boundaries move in parallel to the maximum shear stress direction of the specimen axis.
(2) Accordingly, grain boundaries of polycrystals are gradually migrated orthogonaly in the maximum shear stress directions under a reversed stress.
(3) The above characteristics of grain boundary migrations are observed not only at the surface, but also at the interior of the specimen.
The reasons for the above phenomena are discussed in connection with the release of the stored energy resulting from the boundary sliding in repetition, slips in the matrix along near the maximum shear stress direction and a small amount of the vacancy flow as compared with that of creep.

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An analytical study has been carried out on the rate-determining steps of the reduction of an iron-oxide pellet by hydrogen gas. Assuming that the diffusion through the gaseous film around the pellet, the intraparticle diffusion and the chemical reaction proceeds steadily in succession during the reduction process, the comprehensive rate equation (6) is derived in consideration of the resistance of each process. On the basis of the available data of various experimental conditions for the reduction of an iron oxide, the values of the fractional reduction calculated from Eq. (6) are compared with the observed values in the experiments mentioned above. The result shows a satisfactory agreement between the two.
The ratios of the resistance of each process to that of the over-all reduction process are evaluated quantitatively and illustrated graphically as functions of the particls size, porosity, fractional reduction, gas velocity and reaction temperature. Thus, the relation between the rate-determining steps and the operating conditions of the reduction process has beem clarified.

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Grain-boundary sliding and zone shear during a constant rate heating under various constant shear stresses of 10,000 g/sq. cm to 40.000 g/sq. cm were measured in high-purity aluminium bicrystals (99.999%) with a boundary misorientation angle of 26 deg. The amount of grain-boundary sliding increased with increase of the shear stress up to 20,000 g/sq. cm, but when the stress increased to 30,000 g/sq. cm or 40,000 g/sq. cm, the amount of grain-boundary sliding became smaller than that for the shear stress of 20,000 g/sq. cm. A plot of the logarithm of the amount of the grain-boundary sliding per deg. cent. versus the reciprocal of the absolute temperature was found to lie on a series of a straight line over the lower temperature range. This relationship was discucced by the relation between the amount of grain-boundary sliding per deg. cent. and the rate of grain-boundary sliding, derived by the present authers. The activation energy thus obtained was 10,510 cal/mol for the test conducted in vacuum. According to the microscopic observation, grain-boundary sliding was found together with zone shear, but not vice versa.
In a microscopic observation on the specimen heated up to 300°C under the constant shear stress of 30,000 g/sq. cm, a polygonized structure was found within a band region adjacent to the grain-boundary. Voids formation due to grain-boundary siliding was observed in a common grain-boundary, and their number were dependent upon the amount of sliding. These voids obstructed the grain-boundary migration in the grain-boundary sliding process.

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Ten test pieces of different copper contents were first produced in the range of 66.5∼55.0%, which were then subdivided into five kinds of specimens with different lead contents in the range of 1∼4%. Fifty kinds of lead brass specimens thus prepared were cast to examine their drill-machinability under a constant load drilling. The experiment was carried out to investigate the effects of copper contents and lead contents in the free cutting brass whoes standard differs by countries and the relationship between the drill-machinability of the lead brass and its turn-machinability previously experimented. The drill-machinabilty was mainly investigated in the light of the drill resistance and the drilling time. The results obtained were as follows:
(1) In the drilling resistance, the thrust by drilling under a constant load was measured nearly constant regardless of the materials.
(2) When the copper content exceeded 60%, the torque increased distinctly with increase of the lead content, and below 60% the torque increased slightly. The torque increased with the copper content, and the increase in torque was remarkable in the vicinity of the α- or β-single phase in the (α+β) phase.
(3) The time needed to drill the unit thickness of a plate, that is, the drilling time was in close inverse-correalation with the torque.
(4) The tangential force in turning was in linear inverse-correlation with the torque, and the tangential force was in linear normal-correlation with the drilling time. Therefore, it may be said that there were some relations between the turn-machinability and the drill-machinability.
(5) The difference between the standard for the free cutting brass of different countries and the JIS standard was discussed in the light of the drilling.

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The sintering and some properties of iron-bonded titanium carbides were studied. Mixtures of titanium carbide and iron powders of various compositions (10, 20, 30, 40, 50, 60, 70, 80, and 90% iron by weight) were prepared by ball-milling. Linear shrinkages of compacts were observed during sintering, and the optimum sintering temperatures for each composition were determined. Then specimens were prepared by sintering at those temperatures. Some physical and mechanical properties were measured, and the transverse rupture strength and hardness were compared of the specimens in three different conditions; as sintered, quenched and tempered. Although some properties of iron-bonded titanium carbides may be inferior to those of nickel-bonded titanium carbides, the most promising caracteristic of iron-bonded titanium carbides of a higher binder content is that they are heat treatable. The effect of sulphur on the rupture strength of cermets were examined, but no satisfactory result was obtained.

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The effect of nitrogen on the stress corrosion cracking susceptibility of 18-8 austenitic stainless steels was investigated by transmission electron microscopy.
In the specimen to which added no nitrogen (specimen A), there was a tendency towards a cellular distribution of dislocations, but the specimen to which added nitrogen (specimen B) showed a planar arrangement of dislocations. It was confirmed from these observations that the stress corrosion cracking susceptibility of specimen B was augmented.
These thin-foil specimens were further examined after exposure to solutions which cause stress corrosion cracking, and the mechanism of the stress corrosion cracking was discussed.
When thin-foil specimen B was deformed in a boiling 42 percent magnesium chloride solution, the pitting initiated at active slip planes and advanced along these slip steps, but the chemical attack ceased if the stress was removed. Wherese, in the specimen A showing cross slip, no pitting was observed.
The mechanism of the stress corrosion cracking may be considered as follow:
When the restricted slip planes are continuously exposed to a corrosive environment during plastic deformation, the pitting is initiated at these fresh surfaces. The solute atoms are thus segregated to slip planes with the moving dislocations inside the specimen. There occures a high concentration of solute atoms on the slip planes and in their vicinity. The stress corrosion cracking, once initiated at the fresh surface, would proceed in the specimen with an electrochemical process.

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Please see pdf. Wrong:C_p=20.4+0.0173T, C_p=29.8+0.0232T, −52.88 Right:C_p=19.64+0.0208T, C_p=34.62+0.0112T, −52.58

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