Grain size dependence of secondary creep rate in pure metals and single-phase alloys and the mechanisms/models to interpret the dependence have been reviewed. Two types of the grain size dependence were reported, both of which show a negative dependence approximately below 100 μm. The model proposed by Garofalo in 1960s assumes that the density of dislocations generated at grain boundaries and sub-boundaries determines the secondary creep rate, which is not experimentally supported. Mclean’s model considers preferential subgrain growth near grain boundaries, which might be important in practical steels and alloys with sub-grains such as high Cr ferritic heat resistant steels. Grain boundary sliding (GBS) and its accommodation process in grain is considered as a source of the negative grain size dependence. A finite element modeling performed by Crossman and Ashby, which simulated a deformation process where GBS is accommodated by the power law creep process in the grain interior, indicates that the negative grain size dependence cannot be interpreted by the accommodation process. Based on substructure observation and internal stress measurements, Terada established a “core and mantle” model where the mantle region near grain boundary has no internal stress. This model reasonably interprets the negative grain size dependence.

To clarify the effect of alloying elements (M) on the grain boundary segregation behavior of boron (B) and carbon (C) in α-iron, the grain boundary segregation of B, C and alloying elements was evaluated thermodynamically for the Fe–B–1.0 at.%M and the Fe–C–1.0 at.%M ternary systems (M: Al, Ti, V, Cr, Mn, Nb, Mo) using the parallel tangent law proposed by Hillert. In this calculation, the Gibbs energies of the liquid phase in the Fe–B–M and Fe–C–M ternary systems were applied to those of the grain boundaries. According to the calculated results, in the Fe–B–M ternary systems, co-segregation of Ti, V, Mn or Nb with B was predicted, while no co-segregation behavior was confirmed in the case of Al, Cr or Mo addition; in the Fe–C–M ternary systems, co-segregation of Ti, V, Nb or Mo with C was predicted, while no co-segregation behavior was confirmed in the case of Al, Cr or Mn addition. These co-segregation tendencies correspond well with the formation tendencies of metal borides or metal carbides. Although the present calculated results were based on the assumption that substitutional elements can diffuse sufficiently in addition to interstitial elements B and C, we proposed an equation for the parallel tangent law under paraequilibrium condition in which no partitioning of substitutional elements occurs.

The creep deformation behavior and microstructure of a N-containing steel expected to exhibit high creep strength and excellent oxidation resistance were investigated. Even for steel with a high W content, it was possible to form a martensitic microstructure by adding a sufficient amount of N. Comparison of the microstructures of the N-containing steel and a C-containing steel confirmed that the two steels have the same crystal orientation relationship. The N-containing steel precipitated with the Laves phase as a strengthening phase displayed a higher creep strength than conventional steel under relatively high stress. However, the superiority of the creep strength of the N-containing steel relative to the conventional steel decreased under low stress. The stress exponent of the N-containing steel was different from those of the C-containing steel and the conventional steel. This deference considered to be ascribed to the difference of variation behavior of dislocation density during creep deformation.

In order to understand microstructural changes in 9Cr-1Mo-V-Nb weld metal after long term use, microstructure and precipitates distribution before and after aging at 1013 K were investigated. In the weld metal, regions with coarse or fine prior austenite grains were observed due to thermal cycle during welding. In the coarse grain region, precipitate particles inferred to M₂₃C₆ were densely located on grain boundaries, however, in the fine grain regions, they were sparsely observed not only on grain boundaries but also inside grains. Post weld heat treatment (1013 K/7.7 h) followed by aging (1013 K/100 h) led to ferrite grains formation in the fine grain region. EBSD analysis implied that dislocation density in ferrite grains was low. After the aging, mean diameter of particles became coarser and interparticle spacing became sparser in the fine grain region than in the coarse grain region. On the other hand, dislocation density calculated by hardness in martensite structure was almost no deference between these regions before and after the aging. Therefore, it was suggested that ferrite grains were formed because pinning energy by precipitate particles locally reduced in the fine grain region.

Oxide Dispersion Strengthened (ODS) ferritic steel, a candidate material for fast reactor fuel cladding, has low thermal expansion, good thermal conductivity, and excellent resistance to irradiation damage and high temperature strength. The origin of the excellent high-temperature strength lies in the dispersion of fine oxides. In this study, creep tests at 700°C or 750°C, which are close to the operating temperatures of fast reactors, and high-temperature tensile tests at 900°C to 1350°C, which simulate accident conditions, were conducted on 9Cr ODS ferritic steels, M11 and MP23, and 12Cr ODS ferritic steel, F14, to confirm the growth behavior of oxides. In the M11 and F14 creep test samples, there was little oxide growth or decrease in number density from the initial state, indicating that dispersion strengthening by oxides was effective during deformation. After creep deformation of F14, the development of dislocation substructures such as dislocation walls and subgrain boundaries was hardly observed, and mobile dislocations were homogeneously distributed in the grains. The dislocation density increased with increasing stress during the creep test. In the high-temperature ring tensile tests of MP23 and F14, the strength of both steels decreased at higher temperatures. In MP23, elongation decreased with increasing test temperature from 900°C to 1100°C, but increased at 1200°C, decreased drastically at 1250°C, and increased again at 1300°C. In F14, elongation decreased with increasing temperature. It was inferred that the formation of the δ-ferrite phase was responsible for this complex change in mechanical properties of MP23 from 1200 to 1300°C.

This study was set to fundamentally investigate the characteristics of austenite reversion occurring in maraging steels additive-manufactured by laser powder bed fusion (L-PBF). The maraging steel samples manufactured under different L-PBF process conditions (laser power P and scan speed v) were subjected to heat treatments at 550°C for various durations, compared with the results of the austenitized and water-quenched sample with fully martensite structure. The L-PBF manufactured samples exhibited the martensite structure (including localized austenite (γ) phases) containing submicron-sized cellular structures. Enriched alloy elements were detected along the cell boundaries, whereas such cellar structure was not found in the water-quenched sample. The localized alloy elements can be rationalized by the continuous variations in the γ-phase composition in solidification during the L-PBF process. The precipitation of nanoscale intermetallic phases and the following austenitic reversion occurred in all of the experimental samples. The L-PBF manufactured samples exhibited faster kinetics of the precipitation and austenite reversion than the water-quenched sample at elevated temperatures. The kinetics changed depending on the L-PBF process condition. The enriched Ni element (for stabilizing γ phase) localized at cell boundaries would play a role in the nucleation site for the formation of γ phase at 550°C, resulting in the enhanced austenite reversion in the L-PBF manufactured samples. The variation in the reaction kinetics depending on the L-PBF condition would be due to the varied thermal profiles of the manufactured samples by consecutive scanning laser irradiation operated under different P and v values.

Metallurgical factor causing the heat-to-heat variation in creep rupture strength have been investigated for KA-SUS304J1HTB. In the long-term, there was a maximum difference of 3.5 times in creep rupture time between the heat with low creep strength and the heat with high creep strength. In the heat with low creep rupture strength, most of the creep voids occurred at the matrix/σ phase interface. Moreover, in the heat with low creep rupture strength, the area fraction of σ phase was larger than in the heat with high creep rupture strength. Considering that the difference in the area fraction of σ phase in each heat is related to the difference in phase stability of the austenite phase, the Md value in each heat was evaluated. The Md value is the parameter correlated with phase stability. The creep rupture time of each heat was correlated with the Md value. The smaller the Md value, the longer the creep rupture time. Therefore, the heat-to-heat variation in creep rupture strength is caused by the difference in the phase stability of each heat. In other words, in the heats with low phase stability, creep rupture strength is low because a large amount of σ phase precipitates during the creep test.

The microstructure of the B1-type TiC formed during solidification and its mechanical properties were investigated using arc-melted Fe-Ti-C ternary alloys. The TiC formed at relatively high temperatures in the liquid as the primary phase exhibited a dendritic shape. With decreasing temperature and/or decreasing Ti and C content in the liquid, its morphology changed to a cubic shape with the {001}_TiC habit plane, a plate shape with the {011}_TiC habit plane, and a needle shape with the <001>_TiC preferential growth direction. The morphology of the TiC was characterized by the anisotropy of its surface energy and its growth rate. The cubic shape with the {001}_TiC habit plane was formed as a result of the minimum surface energy of the {001}_TiC. However, the plate shape with the {011}_TiC habit plane and the needle shape with the <001>_TiC preferential growth direction should be formed due to the slowest growth rate of <011>_TiC and the fastest growth rate of <001>_TiC, respectively. At room temperature, alloy with the dendritic TiC was fractured in the elastic deformation region because the TiC exhibited no plastic deformation. On the other hand, at 800ºC, it was suggested that the TiC has plastic deformability and alloy with the dendritic TiC was also plastically deformed.