The brittle fractures of low carbon steels containing notches or microcracks occur easily under high strain rates. In this report, test pieces in which Al2O3 was dispersed in iron were examined by a tensile tester under a strain rate range of 10−4∼10−2 l/sec. The dependence of the strain rate on the yield stress and conditions of the stress concentration owing to inclusions under the strain rates range was investigated by means of transmission electron-microscopy. Ashby’s dislocation model with voids formation around the particles was also discussed in terms of dislocations and voids around the inclusions. The results obtained are as follows: (1) Al2O3 particles were uniformly dispersed in iron. The average diameter of inclusions is about 5 μ. The distances between the inclusions for the test pieces containing 4 wt%Al2O3 and 2 wt%Al2O3 were about 22 μ and 35 μ respectively. (2) The brittle fracture of the specimens containing 4 wt%Al2O3 occurred at a strain rate of 10−2 l/sec. The cause may be attributed to the formation of micro-cracks around the inclusions (Al2O3 type) due to the stress concentration. (3) The micro-cracks around the inclusions in the tensile-fractured specimens under the strain rate of 10−2 l/sec and the voids in the tensile-fractured specimens under the strain rate range of 10−4∼10−3 l/sec were observed by means of transmission electron microscopy. (4) The size of the voids in the tensile-fractured specimens under the strain rate of 10−4 l/sec was bigger than that of the specimens under the strain rate of 10−3 l/sec. (5) The dependence of the strain rates on the yield stress (τy) is expressed as (This article is not displayable. Please see full text pdf.) where \dotε is strain rate and A, B and C are constants.
The aim of the present work is to obtain more detailed information on the phase transformation of metastable beta-type titanium alloys, Ti-(12∼20)Mo. Preliminary investigations revealed that an automatic transformation apparatus “Formastor-F” was applicable to titanium alloys. The phase changes of these alloys during isothermal transformation were mainly detected by the changes in dimension, hardness and microstructure using this apparatus, and then TTT diagrams were obtained. The start point of the contraction coincides with the point where quenched omega transforms to aged omega in the TTT diagram, and the starting point of expansion or the contraction end point also coincides with the point where softening due to the omega→alpha transformation occurs. The contraction and hardening of these three alloys were maximum at 400°∼450°C. In the Ti-15Mo specimens isothermally transformed at 650° and 700°C, the alpha phase precipitated along the grain boundary, while the alpha phase precipitated in the grain at 500°∼600°C. This indicates that the beta→alpha reaction is not so easy as the beta→omega→alpha reaction and occurs only at grain boundary at high temperature.
The effect of Zr addition (0∼40 wt%) on the stability of a metastable beta titanium alloy, Ti-15 wt%Mo, was investigated. The measurement of the mechanical properties in various annealing conditions revealed that the stability of Ti-15Mo-Zr alloy changed abruptly at about 5%Zr. When the alloys containing less than 5%Zr were slow-cooled from the beta region, hardening and embrittlement due to the decomposition of beta was remarkable, and when quenched, a low yield/tensile strength ratio was shown because of martensitic strain transformation. On the other hand, the mechanical properties of the alloys containing more than 5%Zr did not depend on the rate of cooling from the beta region, and the alloys were stable mechanically and thermally. The TTT diagram obtained by an automatic transformation apparatus and the changes in the mechanical properties during aging showed that Zr in a meta-stable beta alloy, Ti-15Mo, suppresses and retards the omega precipitation. The beta stabilizing power of Zr which was calculated from the slope of β⁄α+β solvus in the Ti-Zr equiliblrium diagram was weak, but the above-mentioned beta stabilizing effect of Zr in the Ti-15Mo alloy was as strong as that of Mo.
Internal friction in hydrogenated Ti-Mn alloys have been measured. The frequency of vibration was 70∼80 c/s and the hydrogen concentrations were varied from 1 to 20 at%. The peak observed in the temperature range −160° to −190°C is suggested to be due to stress-induced diffusion of hydrogen atoms dissolved in the beta-phase (bcc structure) of the alloys. The height of an internal friction peak is not proportional to the hydrogen concentration and the peak temperature shifts towards higher temperature as the hydrogen concentration increases. This is suggestive of some interactions between hydrogen atoms. For equivalent hydrogen concentrations, the peak height decreases as the manganese concentration increases. This gives rise to suggestion that the unsymmetrical distortion introduced in the lattice when the solute atoms occupy the interstitial positions decreases with increasing manganese concentration. Each internal friction curve observed is much broader than the theoretical one, implying either the distribution of relaxation times or more than two relaxation phenomena. This deviation from a single relaxation and the lack of knowledge as to the location of the hydrogen atoms in the lattice would complicate the interpretation of the observed relaxation phenomenon.
The structure of columnar dendrite, especially the dendrite element size and the dendrite arm spacing, have been investigated in Cu-Sn and Cu-Zn alloys. In Cu-Zn alloys the dendrite element size is uniform everywhere on a given transverse section of columnar crystals, while in Cu-Sn alloys fine dendrite elements appear in group together with coarse dendrite elements on the transverse section. The fine elements are continuous as a bundle along the growth direction of columnar crystals and the average concentration of tin in the region of fine elements is similar to that of the other region. The dendrite element size is inversely proportional to 0.3∼0.5 power of the cooling rate in the solid-liquid coexisting region, Vs, and the dendrite arm spacing is inversely proportional to 0.1∼0.3 power of Vs. At a given cooling rate, the dendrite element size is roughly proportional to the square root of the solute content for Cu-Zn alloys and is proportional to the solute content for Cu-Sn alloys. The dendrite arm spacing in Cu-Sn alloys is constant at about 10 wt%Sn or less and decreases with increasing solute content at about 10 wt%Sn or more. From these results, it is concluded that the factors governing the columnar dendrite structure in Cu-base binary alloys are similar to those in Al-base alloys.
For the purpose of obtaining high-strength copper alloys having high corrosion resistance, the aging characteristics and the spring property of Cu-30%Ni containing 0.47 wt%Be and 0.5∼1.8 wt%Ti have mainly been investigated by means of the hardness test, X-ray diffraction, optical and electron microscopy (reprica method) and the spring test. The results obtained were as follows: (1) These alloys showed a maximum hardness value at 500°∼600°C when aged at 300°∼800°C for 1 hr after water quenching from 1050°C, and the maximum hardness increased with the addition of titanium. (2) It was concluded that age hardening occurred by heating at 600°C after water quenching from 1050°C was due to the precipitation process of NiBe (β) and Ni3Ti (η), and that the precipitation process of NiBe contributed to the hardening at an earlier stage of aging and the precipitation process of Ni3Ti at a later stage of aging. (3) The added titanium suppressed the grain boundary reaction developed by the precipitation of NiBe and restrained the over-aging of these alloys, but Ni3Ti precipitates grew newly from the grain-boundary in the high titanium alloys aged for a long time. (4) The alloy containing both beryllium and titanium showed a Kb value higher than that of the alloy containing beryllium.
The relationship between compositions and magnetic properties of Al-Ni-Co type permanent magnet alloys has been discussed. All components, excluding Cu, were divided into two groups, i.e. α1 former (Fe and Co) and α2 former (Al, Ti and Ni), based on the Fe-Ni-Al phase diagram and Curie points of several alloys. In order to make a simplified treatment of the compositions, four ratios in atomic % of each element were adopted as follows: (This article is not displayable. Please see full text pdf.) In the K1-K2-K3 solid diagram obtained, the compositions of commercial permanent magnet alloys, viz. Alnico 5, Alnico 8 and Ticonal 2000 were observed to fall on a straight line. Therefore, the various compositions suitable for permanent magnets were easily obtained from the line in the diagram.
Effects of compositions on the magnetic properties of Al-Ni-Co type permanent magnet alloys have been investigated over a wider range of composition than the commercial one. The conditions for solution treatments were determined metallographically. After solution treatments, in order to determine the optimum condition for thermomagnetic treatment (TMT), specimens were held at different temperatures between 700°C and 850°C for 5∼30 min and aged at 600°C for 5 hr, and their magnetic properties were measured. The optimum condition for TMT was the combination of time and temperature at which the highest coercive force was obtained. Moreover, after the optimum TMT, the magnetic properties at different ageing temperatures were measured. The variations in magnetic properties with composition were given in contour diagrams. The diagrams were explained by the variations in Curie points, packing factors and shapes of the α1 and α2 phases with composition.
The effect of the replacement of Al↔Ti on the magnetic properties of Al-Ni-Co type permanent magnet alloys has been investigated. The optimum heat treatments were determined by microstructure observations and magnetic measurements. The variations in the magnetic properties with the Al↔Ti replacement were given in contour diagrams. Four ratios in atomic % of each element were adopted as follows: (This article is not displayable. Please see full text pdf.) Coercive forces more than 1.3 kOe were obtained for the alloys in the range of K1: 1.4<2.3, K2: 0.5∼2.0, K3: 0.4∼0.8 and K4: 0.6∼3.1, and coercive forces more than 1.5 kOe were obtained in the range of K1: 1.5∼2.2, K2: 0.6∼1.1, K3: 0.5∼0.7 and K4: 0.9∼2.5. Curie points and packing factors of α1 (θ1 and p) increased gradually with increasing K4 to broad peaks. The variation in 4πIs with the replacement corresponded to those of θ1 and p. Intrinsic coercive force iHc, however, decreased drastically in the range of higher K4.
Many high-strength alloys, such as Cu-Be alloy, Al-Cu-Mg-Zn alloy (duralumin), 17-7 PH stainless steel and maraging steel are strengthened by dispersing fine precipitates into a matrix. However, the increase of fatigue strength by precipitation hardening is usually smaller than the values expected from other mechanical properties. For example, carbon steel has a value σw⁄σB=0.4∼0.5 (σw: fatigue limit, σB: tensile strength), but as for duralumin, σw⁄σB=0.2∼0.3. These properties paticular to fatigue strength of precipitation-hardening alloys are assumed to be mainly due to the generation of lattice vacancies or interstitials during fatigue process. So, the effect of excess vacancies on ageing process during fatigue is very interesting. Then the effects of point defects generated by fatigue during the ageing process and the influence of fine precipitates on fatigue strength were investigated for Cu-Be alloy, Fe-Ni-Ti alloy, and Fe-Cr-Mo alloy which have relatively high recrystallization temperatures. As a result of experiments, the fatigue strength of these alloys was found to increase by precipitating fine particles in the matrix. The precipitating process is accelerated by the generation of lattice vacancies and interstitials during fatigue. When precipitation-hardening alloys are aged isothermally, the increase rate of fatigue strength is larger than that of other mechanical properties. However, the absolute value for the increase of fatigue strength is small as compared with other mechanical properties. The influence of the size and distribution of fine precipitate on fatigue strength was investigated using a transmission electron microscope.
Crystal structures and magnetic properties of Au-Mn-Sb alloys have been investigated by means of X-ray and magnetic analysis. It has been found that a new Clb type intermetallic compound AuMnSb exists in the Au-Mn-Sb system and forms a solid solution in a narrow compositional range which has lattice parameters of about 6.368∼6.3775 Å at room temperature. The compound AuMnSb has a lattice parameter of 6.373 Å at room temperature, and its relationship between reciprocal susceptibility and temperature satisfies the Curie-Weiss law, From this curve, the paramagnetic effective Bohr magneton number per Mn atom was evaluated to be 5.98 and the paramagnetic Curie temperature to be about 98°K.
The phenomena relating to the two-phase region in TiC-Mo-30%Ni alloys have been studied in order to obtain basic information as to the properties of TiC-Mo-Ni alloy systems. TiC-(0∼30)%Mo-30%Ni alloys were vacuum-sintered at 1300°∼1450°C. The average cooling rate from sintering temperatures to 700°C was about 24°C/min. The results obtained were as follows: (1) The width and location of two-phase region in each alloy seemed not to be affected by the sintering temperatures. (2) The two-phase region of the straight alloy was in the range of 18.1 to 19.3%C(calculated in terms of the carbon content in cabide). The two-phase region in the straight alloy was extended with increasing addition of Mo. For example, the phase region of the 30%Mo alloy existed in the range of 12.7∼14.5%C. (3) The added Mo appeared to dissolve in titanium carbide in the form of Mo2C. (4) Compositions of the binder and carbide phases of each alloy showed a definite variation in the two-phase region with carbon content, as has been confirmed in various WC-Co base alloys. The dissolved amount of Ti in the binder phase of the straight alloy was found to change sharply in the two-phase region from a minimum of about 3% (in the highest carbon alloy) to a maximum of about 10% (in the lowest carbon alloy). Even in the alloys containing Mo, the dissolved amount of Ti was nearly the same as in the straight alloy, but in this case Mo also dissolved in the range from about 0 to 6% depending on the Mo and the carbon contents. (5) The dissolved amount of Ni in titanium carbide was about 0.4% in the straight alloy. The amount of Mo dissolved in the ε phase was about 5% in the alloys containing Mo. (6) The corrosion and oxidation resistance of the alloys were confirmed to vary as can be expected from the change in composition of the binder phase.
Growth of a single pit developed on austenitic SUS 27 (type 304) and SUS 28 (type 304 L) stainless steels has been studied potentiostatically and microscopically in 1 N-H2SO4 containing 0.5 N-NaCl at 25°C. The single pit was developed in the specimen surface of 0.95 cm2 by using, for example, a SUS 27 specimen pretreated in 1 N-H2SO4 at 1200 mV (SCE) and treated for repassivating film-breakthrough sites at 250 mV in the solution containing chloride. The critical potential for pitting is 400 mV for SUS 27 and 250 mV for SUS 28, respectively. This is the boundary potential between the repassivation and pit-growth regions. A distribution hystogram of diameters of film-breakthrough sites occurring in the repassivation region is given. A peak lies at 30∼40 μ. Concernning pit growth, the potential dependence of growth rate was determined. The current density per mouth area of the single pit is 0.3 A/cm2 for both steels at the critical potential. The pit grown near the critical potential has a thick lid (Deckel) over its mouth. The critical potential for pitting seems to be a minimum potential which allows the flow of current enough to concentrate chloride in the pit to a critical value. The lid would be convenient to maintain the high chloride concentration.
The relationship between the structure and hot workability of Al-Mg-Si alloy ingots has been investigated. The ingots were cast at various cooling rates and their metallographic structures were examined. Hot workability of the ingots were evaluated by means of the hot torsion test at 450°C and at strain rate 10 sec−1. It is shown that the maximum shear stress of the ingots, τmax, is given by the relation, (This article is not displayable. Please see full text pdf.) where d is the mean dendrite arm spacing, D is the mean grain diameter and τ0, k and K are constants. This means that the rapidly solidified ingots, with smaller values of d and D, show higher deformation resistance under hot working conditions. Grain refining also results in higher τmax. However, when the cooling rate is too slow during solidification, coarse second phase particles and shrinkage cavities are formed in the interdendritic region and poor high temperature ductility is observed. When those ingots are solution and precipitation-treated, τmax values decreased irrespective of the solidification condition. The practical implication of the obtained results is that the most desirable ingots for hot working can be obtained by rapid cooling during solidification and subsequently by heat treatment under a condition to minimize their high temperature deformation resistance. The measured values of τmax and fracture shear strain are also presented for the properly heat treated direct-chill cast ingots when torsion tested in the temperature range 300°∼500°C and at strain rate 1∼100 sec−1.
The recrystallization of vanadium and its alloy containing 20% titanium was studied by hardness measurement, tensile testing, X-ray diffraction and optical and electron microscopy. The materials were cold rolled to 20 and 50% in reduction and subsequently annealed in vacuum for 1 hr in the temperature range from 500° to 1100°C. The strength of the V-20%Ti alloy decreased at temperatures between 500° and 700°C upon heating after the cold rolling, while in the pure vanadium the decrease occurred less sharply in the temperature range from 500° to 1100°C. Microscopic observation and X-ray diffraction analysis revealed that the recrystallization was complete at 900°C in both materials. The recrystallization nuclei were formed through well-defined subgrain growth in the V-20%Ti alloy, whereas in the pure vanadium the nuclei were formed by a process of strain-induced boundary migration. In both materials, there was no definite correspondence between the decrease in strength and the recrystallization process observed by optical and electron microscopy. The dislocations in the pure vanadium remained locally after heating above 900°C, while in the V-20%Ti alloy dislocations disappeared almost completely after heating above 900°C and precipitates were observed.