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Microstructural evolution in the TiC-reinforced Mo–Si–B-based alloy during tensile creep deformation at 1,500°C and 137 MPa was investigated via scanning electron microscope-backscattered electron diffraction (SEM-EBSD) observations. The creep curve of this alloy displayed no clear steady state but was dominated by the tertiary creep regime. The grain size of the Moss phase increased in the primary creep regime. However, the grain size of the Moss phase was found to remarkably decrease to ss phase occurred by continuous dynamic recrystallization including the transformation of low-angle grain boundaries to high-angle grain boundaries. Accordingly, the deformation of this alloy is most likely to be governed by the grain boundary sliding and the rearrangement of Moss grains such as superplasticity in the tertiary creep regime. In addition, the refinement of the Moss grains surrounding large plate-like T2 grains caused the rotation of their surfaces parallel to the loading axis and consequently the cavitation preferentially occurred at the interphases between the end of the rotated T2 grains and the Moss grains.

A novel procedure of ODS alloy fabrication was examined using a model alloy in the Fe Al O ternary system. Nano-sized Al oxides were precipitated within the α-Fe matrix through a reaction between Fe Al and Fe2O3 powders. Two types of oxides, FeAl2O4 and γ-Al2O3, were found to precipitate during sintering at 1000 °C and the subsequent annealing at 800 °C. The oxides showed sluggish coarsening at 800 °C, which then provided high resistance to coarsening for the α-Fe grains. Microstructure observation and Vickers hardness measurement indicated that the homogeneous dispersion of oxides was achieved through the present procedure without ball milling.

Boron segregation behavior during tempering at 600 and 790 °C was examined in a boron-doped 9Cr-steel by using high-resolution SIMS imaging (nano-SIMS). Boron segregation that developed at prior austenite grain boundaries (PAGBs) during austenitizing was found to decrease reduced during tempering because the equilibrium concentration of B segregation for austenite is larger than that for ferrite so that the segregated boron atoms diffused into block and packet boundaries. The size and the volume fraction of the carbide precipitates at PAGBs were smaller than those at packet boundaries, suggesting that boron segregation at PAGBs suppresses the carbide precipitation upon tempering.

A Ti5Si3-containing multiphase MoSiBTiC alloy with a composition of 38Mo–30Ti–17Si–10C–5B (at.%) was designed and produced by arc-melting. The alloy was composed of five phases—Mo solid solution (Moss), Mo3Si, Mo5SiB2 (T2), Ti5Si3 and TiC—and consistently has good thermal stability at least up to 1700 °C. The density of the alloy was ∼7.0 g/cm3, which is considerably smaller than that of Ni-base superalloys. Microstructure was carefully examined and microstructural anisotropy was confirmed. The anisotropy was considered to be generated by thermal gradient during the solidification process. Microcracking was remarkable across the primary Ti5Si3 phase, which was caused by thermal expansion anisotropy of the Ti5Si3 phase. High-temperature deformability was examined by high-temperature compression tests at 1500 °C. Two kinds of loading axes were chosen for the compression tests with respect to the microstructural anisotropy. The alloy exhibited a peak stress of 450–550 MPa, followed by good deformability at the testing temperature. Microstructure refinement and reduction in microcrack density were observed after hot working. Oxidation tests were conducted on the alloy at 1100 °C and 1300 °C for 24 h. The oxidation curves demonstrated that rapid mass loss finished within several minutes. After that, the mass loss began to slow down and then the specimens' mass decreased almost linearly with increasing testing time. Cross-section observation indicated that oxygen propagated through Moss, whereas T2 and Ti5Si3 phases acted as barriers against oxygen attack during the tests. In addition, it was found that the alloy gained better oxidation resistance after high-temperature deformation, suggesting a positive effect of phase refinement on its high-temperature oxidation resistance.

Heat-to-heat variation of creep strength is significant in grade 91 (Gr.91) steel, but the causes of the variation have not been well understood yet. In the present paper, creep rupture data of 14 heats of Gr.91 steel were analyzed paying attention to their chemical compositions and microstructures. The longest creep rupture lives analyzed are 2 × 105 h at 500 and 550 °C and 105 h at 600 °C. The causes of the heat-to-heat variation are different, depending on creep test conditions. At low temperature and high stress (creep rupture life of 104 h at 500 and 550 °C), creep rupture strength increases with increase of hardness after tempering. This suggests strengthening by a fine subgrain microstructure developed during normalizing and subsequent tempering. At higher temperature and intermediate time range (104 h at 600 °C), creep rupture strength depends on Cr concentration of the heats in addition to the hardness. This finding suggests an important contribution of recovery process of the subgrain microstructures to creep strength of the steel. In long-term creep (2×105 h at 550 °C and 105 h at 600 °C) creep rupture strength primarily increases with increasing grain size of the heats. This suggests that grain boundary sliding is an important deformation mode at low strain rate because of fine grain size usual with Gr.91 steel. Specifications on Ni concentration and N%/Al% ratio are newly introduced in the type II version of Gr.91 steel. They are not effective to eliminate a heat with low creep strength.

Room-temperature fracture toughnesses of TiC-added Mo-Si-B alloys were evaluated for samples of three different compositions prepared using a conventional Ar arc-melting technique. The first alloy (TiCp) had a primary phase during solidification of NaCl-type TiC including an amount of Mo, with a Mo solid solution (Mo ss ) volume fraction of approximately 49% and a TiC volume fraction of approximately 19%, while the volume fraction of Mo 5 SiB 2 (T 2 ) was approximately 31% and the remaining 1% was Mo 2 C including an amount of Ti. The second alloy (T2p) had a primary phase of T 2 , with volume fractions of Mo ss , TiC, Mo 5 SiB 2 (T 2 ), and Mo 2 C of approximately 38%, 4%, 45%, and 13%, respectively. The third alloy (Mop) had a primary phase of Mo ss , with volume fractions of Mo ss , TiC, Mo 5 SiB 2 (T 2 ), and Mo 2 C of approximately 55%, 8%, 32%, and 6%, respectively. Room-temperature fracture toughness was evaluated by three different bending tests using Chevron-notched specimens. Fracture toughness values obtained by the three methods were relatively close with good reproducibility. Consequently, the fracture toughness values of TiCp, T2p, and Mop were evaluated to be ∼15.2 MPa(m) 1/2 , ∼10.5 MPa(m) 1/2 , and ∼13.6 MPa(m) 1/2 , respectively. Fracture surface observations indicated that the Mo ss phase is subject to severe plastic deformation during the fracture process. The TiC phase was also noted to leave river patterns behind through crack propagation. These fractographic results suggest that not only the ductile-phase toughening by the Mo ss phase but also an extra-toughening mechanism by the TiC phase are responsible for the goodness of the room-temperature fracture toughness of the MoSiBTiC alloys.

The microstructures and mechanical properties of as-cast Ti-Mo-Al alloys with various compositions were systematically investigated in this study, aiming to explore the potentials of B2-ordered phase (β 2 ) in high temperature applications. According to the experimental results, these as-cast Ti-Mo-Al alloys were mainly composed of β 2 phase and contained a large number of dislocation walls and sub-grain boundaries. Overall they exhibited a high level of compressive strength at 1073 K, with the maximum value approaching ∼879 MPa. The impressive mechanical properties of β 2 -dominating Ti-Mo-Al alloys are related to the excellent stability of B2 structure and extensive dislocation entanglements at elevated temperature.

Advanced engineering materials 23(11), 2100111 (2021). doi:10.1002/adem.202100111, Herein, synchrotron-generated high-energy X-ray is used to study growthbehavior at the liquid���solid transition of multicomponent alloys during in situdirectional solidification experiments at Deutsche Elektronen-Synchrotron(DESY), Hamburg, Germany. The unique ���FlexiDS��� sample environment is usedto directly investigate crystal growth of a ternary eutectic Mo���17.5Si���8B alloy.During the directional solidification process, high-energy X-rays with a photonenergy of 100 keV (�� �� 0.124 ��) are used in transmission to obtain Debye���Scherrer diffraction rings. The diffraction rings were obtained within the liquidphase, the liquid���solid interphase and the solidified eutectic crystals by scanningthrough the respective observation volume of interest. The results provide strongevidences for a coupled ternary eutectic growth of the phases MoSS, Mo3Si, andMo5SiB2, which can be directly observed during in situ experiments for the firsttime., Published by Deutsche Gesellschaft fu��r Materialkunde, Frankfurt, M.

In this study, α(hcp)/β(bcc) phase boundaries of the Mg-Sc binary phase system were determined by Mg/Sc diffusion couple and conventional equilibrated alloy methods. It was confirmed that α/β phase boundaries exist in regions with a higher Mg content than those in reported diagram. The Mg-Sc alloy with a β single-phase showed tensile strength of 254 MPa and elongation of 25.4%. The tensile strength of the Mg-Sc alloy increased with increasing volume fraction of the α phase ( f α ) and the alloy with a f α of 34% showed high UTS of 310 MPa and elongation of 28.8%, properties which are better balanced than those of conventional Mg alloys. This improvement was suggested to be due to isotropic deformation of the β phase, small grain size and small c / a of the α phase.

Predictions as to 105 h creep rupture strength of grade 91 steel have been made recently. The predicted values are examined with long-term creep rupture data of the steel. Three creep rupture databases were used in the predictions: data of tube products of grade 91 steel reported in National Institute for Materials Science (NIMS) Creep Data Sheet (NIMS T91 database), data of T91 steel collected in Japan, and data of grade 91 steel collected by an American Society of Mechanical Engineers (ASME) code committee. Short-term creep rupture data points were discarded by the following criteria for minimizing overestimation of the strength: selecting long-term data points with low activation energy (multiregion analysis), selecting data points crept at stresses lower than a half of proof stress (σ0.2/2 criterion), and selecting data points longer than 1000 h (cutoff time of 1000 h). In the case of NIMS T91 database, a time–temperature parameter (TTP) analysis of a dataset selected by multiregion analysis can properly describe the long-term data points and gives the creep rupture strength of 68 MPa at 600 °C. However, TTP analyses of datasets selected by σ0.2/2 criterion and cutoff time of 1000 h from the same database overestimate the data points and predict the strength over 80 MPa. Datasets selected by the same criterion from the three databases provide similar values of the strength. The different criteria for data selection have more substantial effects on predicted values of the strength of the steel than difference of the databases.

The site-occupation behavior of Re in Mo5SiB2 (T2) was studied both theoretically and experimentally, and the effect of Re on the solid-solution hardening of T2 was investigated by taking into account the off-stoichiometry of the T2 phase. Mo–Si–B quaternary alloys containing 1.4 at.% Re were produced using a conventional Ar arc-melting technique, and the cast samples were homogenized at 1800 °C for 24 h in an Ar atmosphere. High-resolution high-angle annular dark-field scanning transmission electron microscopy observations strongly suggest that Re preferentially occupies the Mo sites in the Mo–B layers of the T2 unit cell, which was confirmed by the site-occupation behavior predicted by first-principles calculations. Nanoindentation measurements indicate that the hardness of the T2 phase is affected by both the off-stoichiometry and Re addition.

High-temperature compressive properties of two TiC-added Mo-Si-B alloys with nominal compositions of Mo-5Si-10B-7.5TiC (70Mo alloy) and Mo-6.7Si-13.3B-7.5TiC (65Mo alloy) (at.%) were investigated. The alloys were composed of four constituent phases: Mo solid solution (Moss), Mo5SiB2, (Mo,Ti)C, and (Mo,Ti)2C. The primary phases of the 70Mo and 65Mo alloys were Moss and T2, respectively. The compressive deformability of the 65Mo alloy was significantly limited even at 1600°C because of the elongated, coarse primary T2 phase, whereas the 70Mo alloy had good compressive deformability and a high strength in the test-temperature range of 1000–1600°C; the peak stresses were 1800 MPa at 1000°C, 1230 MPa at 1200°C, and 350 MPa at 1600°C. At and above 1200°C, the peak stress values were more than double those of Mo-6.7Si-7.9B, Ti-Zr-Mo, and Mo-Hf-C alloys. The plastic strain in the 70Mo alloy at temperatures lower than the ductile–brittle transition temperature of T2 was generated by plastic deformation of not only Moss but also of (Mo,Ti)C and (Mo,Ti)2C. This work indicates that (Mo,Ti)C and (Mo,Ti)2C play an important role in determining the high-temperature strength and deformation properties of TiC-added Mo-Si-B alloys.

Compressive creep behavior of hot-rolled (40%) Mg-Y binary and Mg-Y-Zn ternary dilute solid solution alloys are investigated in this study. Creep strength is substantially improved by the addition of zinc. Activation Energy for creep in Mg-Y and Mg-Y-Zn alloys are around 200 kJ/mol at the temperature range from 480 to 570 K. These values are higher than the activation energy for self-diffusion coefficient in magnesium (135 kJ/mol). Many stacking faults, which are planar type defects are observed on the basal planes of the magnesium matrix in Mg-Y-Zn ternary alloys. TEM observation has been revealed that the non-basala-dislocation slip is significantly activated by these alloys. The rate controlling mechanism of Mg-Y and Mg-Y-Zn dilute alloys are considered to the cross-slip or prismatic-slip controlled dislocation creep with high activation energy for creep, more than 1.5 times higher than the activation energy for creep controlled dislocation climb.

Recycle-type Fe3Al (hereinafter designated as Re-Fe3Al) based alloys reinforced by the carbides of TiC or ZrC were processed by the high frequency induction melting method using a high-carbon Cr steel sludge, Al can scraps and the transition metals of Ti or Zr. The carbides were synthesized by in-situ reaction between the transition metal and carbon in the molten iron aluminum alloy. Vickers hardness values are 309HV0.5 for Re-Fe3Al/TiC alloy, and 473HV0.5 for Re-Fe3Al/ZrC alloy, which are higher than that of P-Fe3Al (preprared from pure-Fe and-Al). The cutting performance of the Re-Fe3Al baed alloys was compared with a High-Speed-Steel (HSS) by cutting tests for pure-Cu extruded bar (C1020) using a lathe under a dry condition. Tool life limit was estimated from frank wear length after the cutting tests of C1020 by finish-machining. Tool life limit of Re-Fe3Al/TiC alloy is more than16 min; P-Fe3Al was 12 min; HSS was 8 min, Re-Fe3Al/ZrC alloy was 7 min at the cutting speed of 100m/min. Also, tool life limit of the Re-Fe3Al/TiC alloy was more than twice times as long as that of the HSS at the cutting speed of 300/min. The relationship between cutting speed and tool life limit clearly indicated that the Re-Fe3Al/TiC alloy was better than the HSS at a higher cutting speed. Therefore, it was concluded that Re-Fe3Al/TiC alloy has excellent cutting tool performance.

The phase formation behavior during solidification of Mo–Si–B ternary alloys in the Mo-rich compositional portion was experimentally examined with special attention to the Moss–T2 and Mo–Mo3Si–T2 eutectics. Primary phases observed in this study were in very good agreement with those shown in a liquidus projection thermodynamically calculated by Yang and Chang. A Mo–9.5Si–14.2B (at.%) alloy, which has a composition close to the Moss–T2 eutectic point, had two specific microstructural areas consisting of a superfine Moss–T2 eutectic and a fine Moss–Mo3Si–T2 three-phase eutectic. A Mo–19.1Si–6.5B alloy aimed to the Moss–Mo3Si–T2 eutectic point also had a fine Mo3Si–T2 eutectic and a Moss–Mo3Si–T2 eutectic. The compositions of the eutectic phases were determined by electron probe micro-analysis calibrated using a multipoint measurement method and details of the solidification pathways in the Mo-rich compositional portion are discussed based on the experimentally obtained results compared with the liquidus projection.

TiC was added to Mo-Si-B alloys using a conventional Ar arc-melting technique, and the phase equilibria, microstructure evolution, and high-temperature strength at 1673 K (1400 °C) were investigated. The primary phase changed to Mo solid solution (Moss), Mo5SiB2 (T2), or TiC depending on the composition. Following the primary phase solidification, a Moss + TiC, Moss + T2, or Moss + T2 + TiC + Mo2C eutectic reaction took place as the secondary solidification step. In some alloys, Moss + T2 + TiC and Moss + T2 + Mo2C eutectic reactions were present as higher-order solidification steps. After annealing at 2073 K (1800 °C) for 24 hours, Moss, T2, TiC, and Mo2C coexisted stably with microstructural coarsening. The coarsening rate was much faster in an alloy with no TiC dispersion, suggesting that TiC has a strong pinning effect on the grain boundary and interface migration. Compression tests conducted at 1673 K (1400 °C) revealed strength properties of almost all the alloys that were better than those of the Mo-Hf-C alloy (MHC). Alloy densities were 9 g/cm3 or less, which is lighter than pure Mo and MHC (≥10 g/cm3) and competitive with Ni-base superalloys. TiC-added Mo-Si-B alloys are promising candidates for ultrahigh-temperature materials beyond Ni-base superalloys.

Creep rupture strength of creep strength enhanced ferritic steels is often overestimated, and its evaluated value has been reduced repeatedly. In this paper, the cause of the overestimation is discussed, and the creep rupture strength of T91 steel is assessed with its updated creep rupture data. Effects of residual Ni concentration on the creep rupture strength and necessity of F factor in T91 steel are also discussed. Decrease in activation energy Q for rupture life in long-term creep is the cause of the overestimation, since conventional time–temperature parameter (TTP) methods cannot deal with the change in Q. Due to the decrease in Q, long-term creep rupture strength evaluated decreases as longer-term data points are added or shorter-term data points are discarded in the conventional TTP analysis. The long-term region with small values of activation energy and stress exponent is named region L2 in this paper. Region L2 appears in all the heats of T91 steel and plate products of Gr.91 steel. Since service conditions of the T91 steel are usually in region L2, the creep rupture strength under the service conditions should be evaluated from the rupture data in region L2 only. The 5 × 105 hrs rupture strength at 550 °C decreases from 129 MPa (evaluated from the whole data of T91 steel) to 79 MPa (evaluated from the data in region L2 only) with increasing cut-off time for data selection. The 105 hrs rupture strength at 600 °C also decreases from 87 MPa (whole data) to 70 MPa (region L2 only) despite sufficient number of long-term data points at 600 °C. Careful consideration on the data selection is necessary in evaluation of creep rupture strength of the T91 steel. A multiregion rupture data analysis (MRA) is helpful to select data points belonging to region L2.

Single-phase regions in Mo-rich Mo–Si–B alloys at 1800 °C were experimentally determined using a field-emission (FE) gun-type electron probe micro-analyzer (EPMA) in this study. A quantitative analysis by FE-EPMA was conducted with the calibration method we applied in our previous study to improve the accuracy in the B measurements. The compositional ranges of the determined Mo solid solution, silicide and boride phases were in good agreement with those of Mo–Si and Mo–B binary phase diagrams. On the other hand, the determined solubility of B in Mo solid solution, Mo 3 Si and T 1 (Mo 5 Si 3 ) were quite different from that indicated in previously reported ternary phase diagrams. The compositional region of the T 2 (Mo 5 SiB 2 ) single-phase ranges from 9.7 to 13.3 at% for Si and from 23.5 to 26.8 at% for B. The following two points were made clear by the ternary phase diagram as determined in this study. One is that no stoichiometric T 2 composition is in the T 2 single-phase region, confirming no T 2 single-phase material can be obtained at the stoichiometric composition at 1800 °C. The another is that the T 2 region of this study expands to a Si-rich area from the stoichiometric composition.

First of all, the as-cast microstructures of Mo-rich Mo-Si-B ternary alloys were investigated around the triple junction point of the primary Mo solid solution, Mo5SiB2 and Mo2B in this work, based on the liquidus projections of the Mo-Si-B system which have been reported in earlier studies. Subsequently, their microstructural evolution through heat treatment was investigated. Since Mo2B crystallizes out during solidification into a primary or secondary phase even though the alloy composition lies in the triangle of Mo-Mo5SiB2-Mo3Si in the Mo-Si-B equilibrium phase diagram, the as-cast microstructures include the non-equilibrated Mo2B in wide compositional ranges. However, Mo2B was completely decomposed during heat treatment at 1800 °C for 24 h and this contributed to the development of homogeneous, fine microstructures. On the other hand, since Mo2B was not decomposed perfectly during 24 h of 1600 °C heat treatment, as-cast microstructures largely remained. Therefore, it is realized that the heat treatment at 1800 °C is necessary to obtain well-developed microstructures of Mo-Si-B alloys.

Microstructures and wear properties of a recycled Fe3Al-based alloy made from high-carbon Cr steel sludge and Al can scraps were first investigated in this study. The alloy was composed of D03-type Fe3Al matrix and Fe3AlC0.5 dispersions (κ phase). Carbon in the steel sludge and Al can scraps contributes to the formation of the κ phase. Next, the recycled Fe3Al-based composites reinforced by transition metal carbides (TMC = TiC, NbC, V4C3, and ZrC) were fabricated by in-situ reactions between transition metals (TM = Ti, Nb, V, and Zr) and carbon when these transition metals are added to the molten Fe3Al-based alloy. The (Fe,Al)2Nb intermetallic compound and a Fe–Al–Zr–C carbide are also formed in the Nb- and Zr-added Fe3Al-based composites with NbC and ZrC, respectively. The Vickers hardness values of the recycled Fe3Al-based composites reinforced by TMC were higher than that reinforced only by the κ phase. Abrasive wear behavior is observed for these composites. The carbides of the κ phase, TiC, NbC, V4C3 and ZrC are still embedded in the matrix after abrasive wear without delamination or destruction. This result indicates that the carbides significantly affect the wear resistance of these composites. The wear resistance of the composites reinforced by TMCs is higher than that reinforced by the κ phase. Moreover, the wear resistance of the recycled Fe3Al-based composites reinforced by TiC is much higher than that of gray cast iron; the former value is comparable to that of high-carbon Cr steels. Thus, carbon in the steel sludge and Al can scraps is useful to synthesize hard carbides with transition metals. We have succeeded in preparing the valuable recycled Fe3Al-based composites with excellent wear resistance from the industrial steel and Al scraps.

Recycle-typed Fe3Al composites reinforced by Fe3AlC0.5 or TiC phases with several volume fractions were prepared from high-carbon Cr steel sludge, Al can scraps and pure-Ti metal. The feasibility of recycle-typed Fe3Al composites for practical use as cutting tools was investigated by conducting cutting tests under dry conditions using, as work materials, an Al alloy (A5056) for rough machining and pure-Cu (C1020) for finish machining. In the cutting tests of A5056 by rough machining, adhesive wear was the dominant wear mechanism. The ranking of tool life was inconsistent with that of tool hardness, and the tool life of recycle-typed Fe3Al (3.3Ti) was longer than that of high-speed steel (HSS). In the cutting tests of C1020 by finish machining, the tool lives of recycle-typed Fe3Al composites were longer than that of HSS. After a 9 min cutting test with a recycle-typed Fe3Al (1.9Ti) tool, the surface roughness of C1020 was 0.8 μm, about half of that cut with the HSS tool. These results suggested that recycle-typed Fe3Al composites have potential as cutting tools, and showed promise as viable alternatives to HSS.

Creep rupture data and microstructural degradation during aging of high Cr ferritic boiler steels with enhanced creep strength have been studied with special attention to prediction of long term creep rupture life. Tempered lath martensite structure in the high Cr ferritic steels remains unchanged during short term aging, whereas static recovery of the lath martensite structure proceeds when diffusion distance during aging becomes sufficiently long as is the case in long term creep. The static recovery brings about premature failure in long term creep and decreases in apparent activation energy for creep life. The decrease in activation energy is responsible for overestimation of rupture life reported in strength enhanced high Cr ferritic steels. The boundary from a short term region with high activation energy QH to a long term region with low activation energy QL moves towards longer time with decreasing Cr concentration. The difference in activation energy (QH − QL) primarily determines the extent of overestimation of rupture life predicted from short term data. In general, the extent of overestimation is less serious at 9%Cr as compared to 12%Cr.

The influence of annealing atmosphere, temperature and time on the surface oxidation and vacancy condensation behavior of {111}-oriented single crystals of B2-type FeAl was investigated. AFM observation showed that as-annealed surfaces under a high vacuum were rugged and covered with a thin oxide film. The results obtained by TEM indicated that the thin oxide film was {001}-oriented κ-Al2O3 epitaxially grown on the {111}-oriented FeAl surface. AES measurements showed that the thickness of the oxide film was almost twice as thick as that of the passive Al2O3 film formed on the FeAl surface in an ambient atmosphere. It was found that the growth of surface mesopores is attributable to both the condensation of supersaturated vacancies in FeAl substrate and the Kirkendall effect by the surface oxidation during the vacuum annealing.

The phase formation behavior during solidification in the Mo-Si-B system was experimentally examined around the triple junction point (T) of the Mo solid solution (Mo SS ), Mo 5 SiB 2 (T 2 ) and Mo 2 B primary phases in the Mo-Si-B liquidus projection suggested by Yang and Chang through thermodynamic calculations, and the solidification routes of the Mo-Si-B alloys were reconsidered based on the obtained results. Alloys of four different compositions around the T point were produced by arc-melting. The primary phases observed in these alloys were in excellent agreement with the liquidus projection proposed by Yang and Chang. The solidification routes were basically Mo SS (primary) → Mo SS + Mo 2 B → Moss + T 2 eutectic → Moss + T 2 + Mo 3 Si (A15), T 2 (primary) → Mo SS + T 2 eutectic → Mo SS + T 2 + A15, or Mo 2 B (primary) → Moss + T 2 eutectic → Moss + T 2 + A15, depending on the composition. However, a quantitative EPMA analysis indicated the compositions of the Mo SS -T 2 eutectic and Moss + T 2 + A15 phases differed from the reported liquidus projection. This is due to the difficulties interpreting the unusual solidification routes in the Mo-Si-B system and their as-cast microstructures.

Creep rupture life in the time range of 105 h and longer has been evaluated on the basis of creep rupture data available on modified 9Cr–1Mo (Gr.91) steel. The steel is strengthened by a tempered martensite lath structure, which degrades by means of strain induced recovery in short-term creep (Region H). The degradation leads to creep rupture. In long-term creep (Region L), static recovery of the martensite lath structure takes place in addition to the strain induced recovery, resulting in premature creep rupture and the consequent decrease in stress exponent n for rupture life. The premature creep rupture always occurs in Gr.91 steel, but the activation energy QL for rupture life in region L is usually close to the value QH in Region H in the steel. In such a case, long-term rupture life can be predicted correctly by simple extrapolation of short-term data. In some heats of Gr.91 steel, however, another region with low values of n and Q, namely Region L2, appears in the time range longer than 104 h. Simple extrapolation of short-term data results in overestimation of rupture life in Region L2. The static recovery and the consequent premature creep rupture are controlled by migration of sub-boundaries stabilized by precipitates. M23C6 as well as MX precipitates in Gr.91 steel have higher thermal stability than those in other high Cr ferritic steels, resulting in the later appearance or absence of Regions L and L2 in the steel.

Overestimation of creep life is a serious problem that happens in long term creep of welded joints made of 9–12%Cr ferritic steels. Creep rupture data of welded joints made of grades 91, 92 and 122 steels were subjected to multiregion analyses paying attention to change in activation energy Q for creep rupture life. A region with a low Q value was found in long term creep of the welded joints. Therefore, long term creep rupture life of the welded joints can be overestimated when it is evaluated by temperature acceleration tests based on a high Q value of short term data. The decrease in activation energy was closely related to the change in fracture location from base metal to heat affected zone. Activation energy and stress exponent for creep rupture life take similar values both in base metal and in the welded joint, though the transition to the low Q region starts at shorter time in cross-welded joints. This fact suggests that creep rupture life can be formulated in a unified manner irrespectiv...

Samarium (Sm) additions to as-cast Mg-5Al-3Ca-based alloys result in changes, such as equiaxed grains and a refined grain size. The microstructure of as-cast Mg-5Al-3Ca- x Sm alloys consists of an α-Mg matrix, a (Mg, Al) 2 Ca eutectic phase, and an Al 2 Sm intermetallic compound. In as-cast alloys, the (Mg, Al) 2 Ca eutectic phase was located at grain boundaries with a chain structure, and the Al 2 Sm intermetallic compounds were homogeneously distributed at the α-Mg matrix and grain boundaries. The eutectic phase of the extruded alloys was elongated in the extrusion direction and crushed into fine particles because of severe deformation during hot extrusion, and the grain size was refined with an increased amount of Sm addition. The maximum values of the yield strength and tensile strength were 313 MPa and 330 MPa at 2 wt%Sm alloy content, respectively.

The as-cast microstructure of Mg-5Al-3Ca-xY alloy consists of dendritic � -Mg matrix, (Mg, Al)2Ca eutectic phase and Al2Y intermetallic compounds. These two kind of (Mg, Al)2Ca compounds were observed: coarse irregular-shape structure at grain boundary and fine needle-shape structure in the � -Mg matrix grain. This (Mg, Al)2Ca phase of the extruded alloys was elongated to extrusion direction and size of this phase was refined comparing with that of as-cast alloys because of severe deformation during hot extrusion. Maximum yield and ultimate strength value of extruded alloys was 326 and 331 MPa at Mg-5Al-3Ca-3Y alloy, respectively. From these results, it is conclusively demonstrated that Y additions on Mg-5Al-3Ca alloy have more effect to improve mechanical properties. [doi:10.2320/matertrans.MC200777]

The as-cast microstructure of Mg-5Al-3Ca-xNd alloys consists of equiaxed α-Mg matrix, (Mg, Al) 2 Ca eutectic phase and Al-Nd rich intermetallic compounds. With the increase of Nd addition, α-Mg matrix morphology was changed from dendritic to equiaxed grains due to suppression of grain growth by formation of homogeneous intermetallic compounds containing Nd dispersed at grain boundary and α-Mg matrix. And the grain size of as-cast alloys was decreased as addition of Nd was increased. This eutectic phase of the extruded alloys was elongated to extrusion direction and size of this phase was redined comparing to that of as-cast alloys because of severe deformation during hot extrusion. Maximum yield and ultimate tensile strength value of the as-extruded alloys was 322 and 335 MPa at Mg-5Al-3Ca-2Nd alloy, respectively.

The as-cast microstructure of Mg-5Al-3Ca-2Sm alloy consists of equiaxed α-Mg matrix, (Mg, Al)2Ca eutectic phase and Al-Sm rich intermetallic compounds. This eutectic phase of the extruded alloys was elongated to extrusion direction and size of this phase was finered compare to that of as-cast alloys because of severe deformation during hot extrusion. After hot extrusion, the average grain size of Mg-5Al-3Ca and Mg-5Al-3Ca-2Sm alloys was 4.8 *m and 3.8 *m, respectively. In load-unload hardness test, penetration depth was decreased with added Sm and after extrusion procedure because of grain size refining by addition Sm and large deformation. Hardness value of the alloys containing Sm was higher than that of Mg-5Al-3Ca alloy due to grain refining and formation Al-Sm rich intermetallic compound at gain boundary and α-Mg matrix. Maximum hardness value was obtained at the extruded Mg-5Al-3Ca-2Sm alloy at elevated temperatures.

In this study, the high temperature oxidation behavior of polycrystalline MoSi2 in a low-pressure atmosphere was investigated. Polycrystalline MoSi2 was produced by the spark plasma sintering process. Oxidation tests were carried out at 1500°C at either 10Torr or 760Torr in an Ar-20%O2 atmosphere. For both conditions, the weight change peaked at the initial oxidation stage, and then their weights gradually increased with increasing oxidation time. The sample weight became heavier in the ambient pressure than in the low-pressure, but the evaporation oxidation was not significant in the low-pressure condition. After the low-pressure oxidation tests, the formation of Mo5Si3 in the MoSi2 substrate was identified. The oxidation resistance of MoSi2 at 1500°C is discussed based on the obtained results.

In this study, rapidly solidified Co-Sn-Ti and Co-Sn-Si alloys were investigated as candidate cathodic materials for lithium-ion batteries. (Co0.5Sn0.5)1-xTix (x=0.1, 0.2, 0.3, 0.4) and (Co0.5Sn0.5)0.9Si0.1 alloys were produced by Ar arc-melting. These alloys were further rapidly solidified into ribbons by the single-roll melt-spinning method. Constituent phases in the ribbons were identified by XRD, indicating that all of the alloys were composed of crystalline multi-phases with no amorphous phase. The microstructure of the ribbons was observed by SEM and TEM, and the chemical composition of constituent phases was analyzed by TEM-EDS. It is found from the SEM and TEM observations that the grain size of the constituent phases ranged from tens to hundreds nanometers due to the rapid solidification effect.

The apparent activation energy for rupture life sometimes changes from a high value of short term creep to a low value of long term creep. This change results in overestimation of rupture life recognized recently in advanced high Cr ferritic steels. The present study examined how to detect the decrease of activation energy in 9-12 %Cr steels with tempered martensitic lath microstructure. During aging without stress hardness of the tempered martensite microstructures remains almost constant in short term, whereas it decreases with increasing time after long term exposure. The onset of hardness drop can be a good measure of the decrease of activation energy. Causes of the hardness drop and the decrease of activation energy are discussed.

The aims of this research are to investigate the effect of Sm addition in Mg-Al-Ca alloys on microstructure and mechanical properties. Sm addition to Mg-5Al-3Ca based alloys results in the change from dendritic to equiaxed grain morphorlogy and formation of Al-Sm rich itermetallic compounds at grain boundary and α-Mg matrix. And these Al-Sm rich intermetallic compounds were dispersed homogeously and stabilized at high temperature. And maximum yield and ultimate strength value was obtained at Mg-5Al-3Ca-2Sm alloys at elevated temperature because of homogeneous dispersion of stable Al-Sm rich intermetallic compound at high temperature.

In this study, the synthesis of Mo-Si-B multi-phase alloys, so-called in-situ composites, was attempted with the combination of mechanical alloying (MA) and spark plasma sintering (SPS) processes. MA was conducted with mixed powders of Mo, Si and B using a planetary ball mill under various milling conditions. MAed powders were characterized by X-ray diffractometry (XRD) and scanning electron microscopy (SEM). The results obtained by XRD indicated that Mo-Si-B alloyed powders were successfully produced when elemental powders were milled at a higher milling energy. Vacuum heat treatments after the MA process promoted the formation of Mo-Si-B intermetallic phase in MAed powders. On the other hand, the MAed powders were successfully consolidated by a SPS technique, and as a result, sound compacts of Mo5SiB2-based composites were synthesized.

As samarium addition was increased, α- Mg matrix morphology was changed from dendritic to equiaxed grains and average value of grain size was decreased from 101.6㎛ to 39.3㎛. Samarium addition to Mg-5Al-3Ca based alloys resulted in the formation of Mg-Al-Sm thernary intermetallic compounds at grain boundarys and α-Mg matrix grains. In these alloys, two kinds of eutectic structure were observed; coarse irregular-shape structure at grain boundary and fine needle-shape structure in the α-Mg matrix grain. It is found that the yield strength and ultimate strength showed the maximum value of 109.1MPa and 139.3 at Mg-5Al-3Ca-2Sm alloy, respectively.

The formation of vacancy complexes through aging heat treatments between 800 and 1200 K was studied for rapidly solidified NiAl and CoAl ribbons using TEM. In the case of the NiAl ribbons, the density of dislocations decreased with increasing annealing temperature, while voids with tens to a hundred nanometers in size were formed in a low density. In the CoAl ribbons, noteworthy microstructures of dislocations and voids appeared after annealing. A zone consisting of dislocation loops, in which prismatic loops 20–50 nm in diameter had formed at high density, was found in the CoAl ribbons annealed at 800 K for 360 ks. On the other hand, voids with a size of several nanometers and a remarkably high density were formed by annealing not only at 800 K but also at 1200 K, suggesting that the growth rate of voids is extremely low because of low diffusivity in CoAl.

Nanoscale dynamic transformations during tensile and compressive deformation of Zr65Al7:5Ni10Cu17:5 and Zr65Al7:5Ni10Pd17:5 bulk metallic glasses have been investigated. Although no apparent differences are observed in the stress-strain curves in the tensile deformation between the two alloys, fine striations and depression zones with viscous flow appear at the fracture surface near the edge in the Zr65Al7:5Ni10Pd17:5 alloy. Unlike the Zr65Al7:5Ni10Cu17:5 alloy and other bulk metallic glasses, the Zr65Al7:5Ni10Pd17:5 bulk metallic glass exhibits a large plastic strain of approximately 7% during compressive deformation. By detailed examination of the microstructure, we provide direct evidence for nanoscale multistep shear band formation in the Zr65Al7:5Ni10Pd17:5 metallic glass. A novel nanoscale structure where fcc Zr2Ni nanocrystalline particles are arranged in ‘‘bandlike’’ areas in the glassy matrix is observed near the compressive fracture tip. The suppression of the propagation of the shear bands due to dynamic nanocrystallization causes this structure. Furthermore, the results are recognized as a novel phenomenon, a nanoscale dynamic structural change by shear band propagation, and provide a new method for improving the mechanical properties of bulk metallic glasses. [doi:10.2320/matertrans.MF200615]

In this work, the characteristics of shear bands and fracture surfaces of Zr 65 Al 7.5 Ni 10 Pd 17.5 bulk metallic glass fractured by a tensile test was investigated. Zr 65 Al 7.5 Ni 10 Pd 17.5 bulk metallic glass shows a yield stress of approximately 1.3 GPa, a fracture stress of approximately 1.5 GPa and a tensile plastic strain of 0.1-0.2% irrespective of the applied strain. Wavy, meandering shear bands were observed in the relatively wide area of specimen surfaces around the point of failure, and typical vein patterns were observed on the fracture surfaces. Shear bands and fracture surfaces were further examined by confocal microscopy to obtain more precise information on their roughness. On the other hand, evidence of viscous flow due to crack propagation was also obtained around the edge at the point of failure on specimen surfaces by confocal microscopy. The deformability of Zr 65 Al 7.5 Ni 10 Pd 17.5 bulk metallic glass is discussed on the basis of the obtained results.