The emissivity change caused by oxidation might produce a large error in temperature measurement by a pyrometer in a steel making process, such as continuous annealing line of cold rolled steel. We measured in detail the change of spectral emissivity of metal samples during an oxidation process. It was confirmed that there were two types of emissivity changes caused by the oxidation. One type was the oscillation, and the other was monotonous increase. In cold rolled steel samples, the former appeared when the Si and Mn concentrations were low, and the latter appeared when the Si and Mn concentrations were high. These phenomena can be explained by the fact that surface silicon oxide layer prevented further oxidation. Regarding the relationship between oxide thickness and emissivity change, emissivity oscillated as the result of the interference at the thin oxide layer on the sample surface. Because emissivity changes in a very complex way during an oxidation process, it is difficult to compensate the change of emissivity. In order to solve this problem we employed a multiwavelength pyrometer. This method was effective when emissivity was less than 0.8 and the oscillation did not occur. We manufactured a pyrometer using four wavelengths, and the accuracy of temperature measurement was ± 10°C with the emissivity range from 0.4 to 0.8 in a laboratory experiment. Another method was proposed which could be applied to the oscillation.
Ti-6Al-4V seamless pipe was manufactured by the inclined rolling process to take advantage of its higher productivity than the conventional hot extrusion process. Properties of the rolled Ti-6Al-4V seamless pipe was investigated. Also, the pipes, extruded at alpha+beta phase temperature and beta phase temperature, were compared each other. The inclined rolling process was successfully applied to produce Ti-6Al-4V seamless pipes. The pipes possessed uniform fine accicular microstructure and no surface flaws. The balance between 0.2% proof stress and elongation was optimum in the rolled pipe, followed by the alpha+beta extrude pipe with elongated primary alpha phase, and the beta extrude pipe which consisted of coarse beta grains with accicular structure. Crystallographic texture was weakly formed in the rolled pipe, which was similar to that of rolled plate. Similar, but slightly intense texture was formed in beta extruded pipe. Tensile and fracturing properties were isotropic in the rolled and the beta extruded pipe. On the other hand, intense texture was formed in the alpha + beta extruded pipe. C axis direction of titanium alpha phase was intensely orientated to both the transverse and the longitudinal directions. Tensile and fracturing properties were anisotropic in the alpha+beta extruded pipe, due to the intense texture formation. The degree of anisotropy in the properties was correlated with the ratio of the C axis orientation fraction, regardless their manufacturing processes.
Effect of Si and temperature on an initial stage of oxidation of Fe-Si (01.5 mass%) alloys in air was investigated for up to 150s at temperatures between 1323 and 1473K. Iron and the Fe-0.1Si alloy formed a triplex oxide layer structure of FeO (Fe3O4 included)/Fe3O4/Fe2O3, obeying a parabolic rate law and the temperature dependence of the parabolic rate constants yield activation energies of 101 and 156 kJ/mol for Fe and an Fe-0.1%Si alloy, respectively. The alloys containing 1.0 and 1.5 Si showed very slow oxidation below 1373K due to a formation of duplex Fe2O3 and SiO2 layers. These alloys oxidized faster, obeying a linear rate law owing to liquid Fe2SiO4 formation in the triplexlayer structure of FeO (Fe3O4 included)/Fe3O4/Fe2O3 at 1473K. At temperatures of 1373 and 1423K the alloys containing 0.4, 0.5, and 1.0 Si oxidized slowly during initial periods and then the oxidation rate increased rapidly due to liquid Fe2SiO4 formation. Measurements of sample temperatures showed that this is due to over-heating of the sample by the rapid oxidation. The composite FeO and Fe3O4 layer between the FeO and Fe3O4 layers was suppressed in the Fe and Fe-low Si alloys, when oxidized at temperatures below 1173K, while it appeared even in rapidly quenched samples oxidized above 1273K..
Effects of retained austenite parameters (volume fraction and stability) and second phase morphology ("a network structure" and "an isolated fine and acicular one") on the deep drawability in high-strength TRIP-aided dual-phase (TDP) sheet steels with different silicon and manganese contents were investigated. The deep drawability was evaluated with a limiting drawing ratio (LDR=Do/dp), where the Do and the dp are a maximum blank diameter drawn out and a punch diameter, respectively. The deep drawability was affected by the volume fraction of the retained austenite rather than by its stability (carbon concentration) and morphology. Namely, the higher the volume fraction of the retained austenite, the larger the strength-deep drawability balance, i.e., the product of tensile strength and LDR. Such excellent deep drawability was caused by large local necking resistance at the cup wall just above the punch bottom due to "the transformation hardening" and "the stress relaxation" resulting from the strain-induced martensite transformation, as well as a low drawing resistance of the shrinking flange. Furthermore, the earing behavior of drawn cup was preferably decreased by an acicular type of retained austenite that was mainly isolated in the ferrite matrix, away from bainite phase.
Torsional fatigue strength of induction hardened steels has been affected by fracture origin. The purpose of this study is to clarify a relationship between fracture origin site in torsional fatigue test and hardness distribution of induction hardened steels. In the test-pieces of a shallow case hardeninng depth, fracture tended to occur on internal origin. When case hardening depth increased and exceeded a certain level, the fracture origin changed from internal to surface. A new indicator, "projected core hardness", defined as a core hardness projected from internal fracture origin site to surface along stress distribution, was proposed. The site of fractue origin was dependent on the ratio of projected core hardness and case hardness. That ratio of 1 corresponded to the critical condition which fracture occurred on internal or surface origin. In addition, the site of fractue origin was dependent on stress amplitude, too. The fracture tended to occur on surface origin with the test condition of high stress amplitude. This was because the compressive residual stress at surface decreased with increasing stress amplitude. Torsional fatigue mode maps, shown as a function of the ratio of projected core hardness, case hardness and stress amplitude, were newly proposed. These maps enable to predict the fracture origin site in torsional fatigue test.
The purpose of this study is to clarify an effect of hardness distribution on the torsional fatigue strength of induction hardened steels. The factors of hardness controlling crack initiation life and breakage life were investigated against the each fracture mode of internal and surface fracture origin. Torsional fatigue strength of induction hardened steels was dependent on each of the hardness and depth of the hardened layer and the hardness of core zone on the internal and surface fracture origin, respectively. In addition torsional fatigue strength increased with increasing these factors of hardness distribution. The crack initiation life of the test-pieces fractured with surface origin could be expressed by a new indicater, "equivalent hardness", defined as a mean hardness weighted with the radius squared. The breakage life of the test-pieces fractured with surface origin in low cycles faigue test could be expressed by the equivalent hardness, too. This was because the crack initiation life was dominant to the breakage life in low cycles fatigue test. The crack initiation life of the test-pieces fractured with internal origin could be expressed by a new indicater, "projected core hardness", defined as a core hardness projected from internal fracture origin site to surface along stress distribution. The breakage life of the test-pieces fractured with internal origin could be expressed roughly by the projected core hardness, too.
In 1977, Japan Institute of Art Japanese Sword was reconstructed Tatara furnace for direct steelmaking from iron sand and charcoal, so called Kera-oshi method. This Tatara furnace called Nittouho Tatara was constructed on the base of Yasukuni Tatara furnace which had been operated until the end of the World War II in Yokota city in Shimane prefecture. For this reconstruction of the furnace and the direct steelmaking, operation, the great effort of Mr. Yoshizo Abe as a leader Murage had been paid and his techniques should be made clear. Until the age of Yasukuni Tatara, Kera-oshi method was consisted of 4 stages; Komori, Komoritsugi, Nobori and Kudari. In the 2nd stage of Komoritsugi, Komori iron sand had been used to charge because of easy reduction and production of pig iron. In 1977, Mr. Abe had met difficulty to collect Komori iron sand. Then, he developed the new technique of Tatara operation using only Masa iron sand for the last two stages in spite of Komori iron sand. He controlled the wet of iron sand and made the residual time of iron sand longer in furnace. The reduced iron particles have enough time to absorb carbon for producing pig iron.
The Tatara furnace, so called Nittoho Tatara, was reconstructed in 1977 based on "Yasukuni Tatara" which had been workedbefore the end ofthe World War II. Mr. Yoshizo Abe as a leader "Murage" developed the new techniques of direct steelmaking, "Kera-oshi" for Tatara operation only using "Masa" iron sand for the 3 stages of "Komori", "Nobori" and "Kudari". Until the age of "Yasukuni Tatara", the operation hadbeen carried out in 4 stages with "Komoritsugi" following "Komori", and in the first two stages, the "Komori" and "Komoritsugi" iron sandswith higher (Fe2O3)/(FeO) ratio had been used, respectively. In the presentwork, 7 Tatara operations with 3stages were carried out using"Komori" iron sand or "Masa" iron sand for the "Komori" stage followed by the last two stages using "Masa" iron sand.The operation using"Komori" iron sand produced more amountand higher carbon content of "Tamahagane" than that only using "Masa" iron sand, because "Komori" iron sand was reduced easier than"Masa" iron sand.