In order to improve the wear and fatigue properties, radical nitriding was applied to Ni-base superalloy, Alloy 718, which is difficult to nitride at low temperature because of its stable passive film. Nitriding conditions selected were at 500°C∼570°C and for 10h∼20h in consideration of practical application. Compound layers formed by these nitriding conditions were 4μm∼10μm in thickness. By nitriding at higher temperature and for longer time, base alloy was softened, though the thickness of compound layer was increased. Wear properties were improved by the nitriding at both of room temperature and 500°C, especially at room temperature. The improvement of wear properties was caused by surface hardening. Moreover, fatigue strength at room temperature increased. The increase in fatigue strength was mainly caused by suppression of fatigue crack initiation due to compound layer.
Effect of radical nitriding on fatigue strength of Alloy 718 was investigated at room temperature and 500°C under push-pull loading. Nitrided alloy used was with compound layer of 4μm. Fatigue strength was increased at room temperature by nitriding, though there was no or little difference in fatigue strengths between aged alloy and nitrided one at 500°C. All of fracture occurred from specimen surface in both alloys at room temperature. On the other hand, fracture origins were specimen surface at high stress levels and subsurface at low stress levels in both alloys at 500°C. Surface crack in the aged alloy initiated in ductile manner and the one in nitrided alloy was in brittle manner. The increase in fatigue strength at room temperature was mainly caused by the suppression of crack initiation by compound layer. The reasons for the change in fracture origin from specimen surface to subsurface at 500°C were the suppressions of crack initiation by compound layer in the nitrided alloy and the one of propagation of small crack by oxide induced crack closure effect in the aged alloy. In both of aged and nitrided alloys, fatigue strength at 500°C was larger than that at room temperature in high cycle region where subsurface fracture occurred.
Rotating bending fatigue tests were carried out using a radical nitrided nickel base superalloy, Alloy 718, to investigate the effect of the nitriding on initiation and propagation behavior of a fatigue crack. The nitriding was done for 12h at 500°C and for 20h at 570°C in consideration of practical application. Compound layers were formed on the specimen surface with the thickness of about 4mm and 10mm respectively under these conditions. Hardness of the base alloy was hardly affected by nitriding at 500°C, but it was decreased at 570°C. Fatigue strength increased by nitriding. The increase in fatigue strength was mainly caused by the suppression of a crack initiation through the compound layer. A surface crack in the aged alloy initiated in ductile manner and one in the nitride alloy initiated in brittle manner from the specimen surface. Also, crack growth rate in the alloy was slightly decreased by nitriding at 500°C, but it was accelerated by nitriding at 570°C.
Crack propagation tests of lead-free solder were conducted using center-notched plates under cyclic tension-compression of three waveforms : pp waveform having fast loading and unloading rates, cp-h waveform having a hold time under tension, and cc-h waveform having a hold time under tension and compression. In fatigue loading at fast loading-unloading rates, i.e. pp waveform, the path of crack propagation was macroscopically straight, perpendicular to the maximum principal stress direction. The introduction of creep components by tension and compression holds in cc-h waveform promoted shear-mode crack propagation even under tension-compression loading. For fatigue loading of pp waveform, the crack propagation rate was expressed as a power function of the fatigue J-integral and the relation was identical for load-controlled and displacement-controlled conditions. The creep component due to the hold time greatly accelerates the crack propagation rate when compared at the same values of the fatigue J integral or the total J integral (the sum of fatigue J and creep J integrals). The creep crack propagation rate was expressed as a power function of the creep J integral for each case of cp-h and cc-h waveforms. The crack propagation rate for cp-h waveform is higher than that for cc-h waveform. The predominant feature of fracture surfaces was striations for pp waveform and grain boundary fracture for cp-h waveform. Grain fragmentation was evident on the fracture surface made by cc-h waveform.
Microstructures of Gr.122 and Gr.92 weld joints were investigated mainly by electron backscatter diffraction pattern (EBSP) in order to clarify the difference of creep void initiation site between 9Cr and 12Cr steels. Creep tests of welded specimens were conducted at 898 and 923K for 90MPa. Most of the initial grain boundaries of Gr.122 had straight configurations and their rotation angles were around 60° in both the base metal and HAZ. In the creep ruptured specimens, many fine grains and sub-boundaries were observed in addition to the block boundaries and consequently the decrease of the boundaries with 60° misorientation angle and the increase of boundaries with low misorientation angle were detected. This tendency was relatively remarkable in fine grained HAZ. It seemed that the recovery of martensite structure in fine grained HAZ was faster than other region. In addition, creep voids formed on the prior austenite boundaries in fine grained HAZ of Gr.122 weld. On the other hand, the ratio of random grain boundaries in fine grained HAZ of the as-heat treated conditions was higher in the Gr.92 than in the Gr.122. Many creep voids formed on the random grain boundaries in the Gr.92 weld. It seemed that the formation of creep voids in Gr.122 weld was suppressed because the martensite structure in fine grain HAZ was more stable for the Gr.122 weld than for Gr.92 weld.
Microstructures of KA-SCMV28, KA-STPA29 and KA-SUS410J3 weld joints were investigated mainly by electron backscatter diffraction pattern (EBSP) in order to clarify the differences of fine grained HAZ formation mechanism among these steels. The welded specimens were heat treated at 740-760°C for several hours as the post weld heat treatment (PWHT). In the base metal of KA-SCMV28 and KA-STPA29, most of the observed boundaries were block and packet boundaries and the ratio of the boundaries with rotation angle of 60° was high. In the fine grained HAZ region of these steels, boundaries with rotation angles of 22 to 46°, which were different from K-S relationship, were detected and the ratio of block boundaries with rotation angle of 60° decreased. In order to investigate the formation mechanism of fine grains, as-welded specimen of KA-SCMV28 without PWHT was observed. Boundaries (22-46°) which were different from K-S relationship have already formed in the as-welded condition but the ratio of these boundaries was lower than in the specimen after PWHT. It was considered that fine grains were formed by the transformation of α’ to γ to α’ or α’ to α during welding cycle. On the other hand, in KA-SUS410J3 most of boundaries were block and packet boundaries and distinct fine grain region was not observed. In addition, distribution of the rotation angle was not much different between the base metal and HAZ region.
A new recrystallization phase-field model is proposed, in which the three stages of static recrystallization phenomena, i.e., recovery, nucleation and nucleus growth are sequentially taken into account in a computation. From the information of subgrain patterns and crystal orientations in a polycrystal that are obtained by a dislocation-crystal plasticity FE analysis based on a reaction-diffusion model, subgrain groups surrounded by high angle boundary are found out. Next, subgrains in the group are coalesced into a nucleus by rotation of crystal orientation and migration of subgrain boundaries through a phase-field simulation. Then a computation of nucleus growth is performed also using the phase-field model on account of an autonomic incubation period of nucleation, in which stored dislocation energy assumes a role of driving force. It is shown that the present method can numerically reproduce the three stages of static recrystallization as a sequence of computational procedure.
In this study, an attempt was made to develop tank and modified tank models of surface runoff and groundwater recharge of porous pavements based on the results of model tests conducted to evaluate the stormwater infiltration and storage functions of porous pavements. As a result, it was revealed that the tank and modified tank models were fully applicable to the evaluation of groundwater recharge functions of small-scale artificial structures like porous pavements and that the amounts of surface runoff and groundwater recharge by porous pavements could be accurately reproduced by the tank or modified tank model by using different parameters in areas at rainfall intensities above and below 70.0 mm/hr where stormwater penetration behavior in porous pavements varied. In order to evaluate the stormwater storage capacity of porous pavements, numerical tests were conducted using the modified tank model. As a result, it was verified that porous pavements had little stormwater storage capacity.
The tests for sliding characteristic of sintered aluminum with retained oil for sliding bearing were conducted . The sintered aluminum materials blended of silicon with 40wt% weight in maximum were made for experiments. Results obtained are as follows. There was the relation among the oil content, the radial crushing strength and the hardness. There was the limitation of mean grain size and weight fraction of silicon for the higher quantity of oil and the higher mechanical properties. In that condition, the hardness increased with the quantity of retained oil in subsurface of tested rings.
A clear yield point was not seen in the tensile strength of the aluminum sintering body. As for the molding pressure, 170MPa was the upper limit value. Hardness, the radial crushing strength increased, but the oil content. As for the sintering temperature, 545°C were the upper limit values. Hardness, the radial crushing strength increased, but the oil content increased to 545°C, too. The friction examination of molding pressure 170MPa showed limit PV value 6.46MPa.m/s. The friction examination of sintering temperature 545°C showed aspect pressure 17MPa. The life examination maintained friction temperature to approximately 40°C until 30hours.