We designed the tensile impact testing machine. The maximum test speed is at 25m/sec. We examined the tensile impact behavior of 5 kinds of plastics, polyethylene, polypropylene, etc., and discussed their dependence on the impact velocity, and the notch effect.
This paper describes a non-destructive approach to measure the distribution of brittle transition temperature of welded steel plate by means of hardness measurement at low temperature. The transition temperature in this hardness test is defined as the intersections of two linear stages in the description of the logarithmic Vickers hardness versus the reciprocal of absolute temperature. The two intersections were found in the range between room temperature and -196°C. The higher one is called as the first discontinuity temperature and the lower is the second discontinuity temperature. Both the temperatures are plotted as a function of distance from the center of welding, and compared with the result by usual Charpy impact test. The hardness method has a special advantage for the case to change the properties in various ways in the narrow range as the heat affected zone by welding. The characteristics of the three stages in the temperature dependency of hardness are discussed from the observations of the deformations around the Vickers indentations by means of microscope and roughness meter.
Investigation was made of the structure of crack region observed at the center of the first weld layer of AN 31 alloy steel. The weld crack originated in the boundaries of dendritic grains containing the primary and secondary precipitates. From the results of the X-ray, microphotographic and differential thermal analyses, it is clearly to be seen that the primary precipitates are Nb (C·N) intermetallic compound consisting of lattice constant: 4.404Å which does not dissolve in matrix even by heating at 1300°C, and that the secondary fine precipitates that appear in the grain boundary regions by tempering at about 900°C are the same as Z phase (MNb2) which were reported by Andrews et al. (1957). From these results, it is evident that the weld crack had originated in the incoherent interfaces of the matrix and the primary precipitates at higher temperature than 1000°C in rapid cooling stage after solidification.
A brief report of studies is made hereunder of the structure of first weld layer of AN 31 alloy steel mixed with about 5% ferrite by using the NC-32 welding rod, and of its bending creep strength at 650°C in air. The results obtained are as follows. (1) The boundaries of austenite dendritic grains in the first weld layer are inclosed among the mixed ferrite grains, and no weld crack is observed in this region. (2) The mixed ferrite phase dissolves into austenite +Cr23 C6 carbite by being heated at 650°C for a long period, and these Cr carbides dispersed in the grain boundary regions accelerate preferred oxidation at the grain boundaries during the creep test. (3) During the creep test at 650°C the weld specimens mixed with ferrite in the first weld layer show higher creep rate in a steady state for a certain period than the specimens with no mixture It is considered that the increase of creep rate is due to the fact that the continuity existing among the crystal grains under the maximum bending stress is broken by the preferred oxidation on the boundary.
In order to determine high fatigue strength of 18-8 stainless steel at room temperature (to be hereunder abridged as R. T.), the authors performed the two following series of experiments. Experiment A: (1) Measurement of the amount of ferrite and micro-Vickers hardness at R. T. of the specimens to which various extents of stretching were given at -77°C. (2) Measurements of the same at R. T. of the specimens heat-treated at several elevated temperatures after sub-zero working. Experiments B: Rotary bending fatigue test at R. T. of the four followiug kinds of test specimens: (1) sub-zero worked specimens, 20% stretched at -77°C. (2) cold worked specimens, 20% stretched at R. T. (3) specimens 20% sub-zero stretched at -77°C, and later heat-treated at 550°C for 100 minutes. (4) specimens heat-treated in solution at 1050°C for 60 minutes (S. H. T.), followed by water quenching. The results of experiments can be summarized as follows; Experiment A: An extraordinary amount of ferrite was recognized in the specimens heat-treated at 100°C for 30 minutes. This phenomenon disappeared at 300°C. Experiment B: The combination of sub-zero working and subsequent heat-treatment at elevated temperature was found to increase fatigue strength effectively at R. T. The specimens (3) showed the highest fatigue strength 53.2kg/mm2, i.e. 90% greater than that of the specimen (4). These phenomena may be explained as due to the decomposition of metastable austenite and to the solubility of micro carbides.
The dynamic behaviors of super duralumin (24S) have been studied through experiments in which the material in the form of bars was shot from a gun of compressed nitrogen gas at a steel anvil bar. The impact speeds was varied from tens of meters per sec to 190 meters per sec by which the test was made in the range of about 20% strain. The test of the specimens was performed just after the quenching and after the age-hardening for 80min at 100°C. The main results obtained are as follows; (1) By comparison of the dynamic stress-strain curve of which the strain rate is estimated at about 104 1/sec and the static stress-strain curve of which the strain rate is (1∼4)×10-4 1/sec, the sensitivity to the strain rate is hardly recognized in the age-hardened specimen, but is obviously recognized in the specimen fresh following the quenching. (2) It is assumed that even after the age-hardening of the material there will be variation in the internal structure of the specimen due to the nature of the deformation, whether it is static or dynamic. The effect of the variation on flow stress is not obvious till the strain is increased up to about 10%. (3) All the stress-strain curves obtained statically and dynamically of the specimen after age-hardening and those immediately following the quenching are approximately represented by the following equation, σa=σe+C2(εa-σa/E)1/2, where σe: yield stress, E: Young's modulus and C2: constant. It can be recognized in these curves that the stress value extrapolated from these curves at strain 100% is approximate to the characteristically constant value of the material. (4) The relation between the impact speed V1 and the strain ε1, V1=∫ε10(1/ρdσ/dε)1/2 dε, is applicable to the range of strain up to about 18% or to the range of impact speed up to about 180 meters per sec of both the specimen fresh following the quenching and of the age-hardened specimen.
The packing models such as those devised by Hudson and Horsfield and ordinarily used for mixed powders of different sizes are not applicable to fine powders in which the force of interaction between the powder particles can not be neglected. The change of porosity based upon mixing of fine powders having different particle size, was measured. The porosity increased rapidly by addition of small amount of fine powders. However, the porosity seemed to remain constant within certain range of further addition of fine powders. But it increased again when more amount of fine powders were added beyond this range. The mechanism of this process can be explained by a simple assumption that cohesion of small particles to larger ones increased their points of contact. And the relation between the weight fraction of the small particles added and the increase ratio of porosity of powder bed in every mixed powder was obtained as Ws/f(ε)×(W0-Ws)=α+βWs/W0 where Ws is the weight of added powder, W0 is that of total powder mixed and ε is the porosity. The constant β was related to the equilibrium porosity, but the meaning of the constant α was not yet known.