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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.

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.

TiC-added Mo-Si-B alloys are stably composed of Mo ss , Mo 5 SiB 2 , (Ti, Mo)C, and (Mo, Ti) 2 C phases. In the alloys, interfacial cracking mainly occurred at the interface between Mo ss and Mo 5 SiB 2 (T 2 ). Thus, other interfaces such as between Mo ss , (Ti, Mo)C and (Mo, Ti) 2 C were investigated for understanding the mechanical properties. The above three kinds of the interfaces have special atomistic structures probably favorable to the mechanical properties. At the interface of Mo ss /(Mo, Ti) 2 C, (Mo, Ti) 2 C decomposes into Mo ss and (Ti, Mo)C, and then serrated, blade-shaped interfaces are formed. The Mo ss /(Ti, Mo)C interface has an orientation relationship of (110) Mo //(111) (Ti, Mo)C . The (Ti, Mo)C/(Mo, Ti) 2 C interface forms a ledge and terrace semi-coherent structure, and its orientation relationship is (111) (Ti, Mo)C //(100) (Mo, Ti) 2 C and [0 1 ¯ 1] (Ti, Mo)C //[010] (Mo, Ti) 2 C . Therefore, Mo ss , (Ti, Mo)C and (Mo, Ti) 2 C have good coherency in the MoSiBTiC alloy. (119 words)

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.

The electromagnetic acoustic resonance method was applied to measure the elastic constants ( c 11 , c 12 , and c 44 ) of Fe–40, –43 and –46Al (at.%) single crystals from room temperature to 1173 K. The temperature dependence of the energy factor, K , and the elastic anisotropy factor, A , was obtained. A increased monotonously with increasing temperature and reached ∼7 at ∼1200 K. Therefore, it was confirmed that B2-type FeAl shows an anomalous temperature dependence of elasticity.

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.

It has been known that Cu- or Ag-addition Al-1.0mass%Mg2Si (balanced) alloys shows higher hardness and elongation than Cu-free or Ag-free balance alloy. In this study, the alloys with Cu or Ag addition and the alloys with Si / Mg in excess have been investigated by hardness and tensile tests and HRTEM observation. Cu addition is effective for higher hardness, and Ag-addition is useful for improvement of elongation for peak-aged samples. Precipitates in peak aged these alloys have been confirmed by HRTEM. Cu-addition alloy almost includes Q’-phase, and Ag-addition alloy includes b’-phase. The precipitation sequence of Ag- or Cu addition Al-Mg-Si alloy was investigated using HRTEM, SAED, and EDS. The precipitates obtained in the two alloys were classified into several kinds by HRTEM images and SAED patterns. The relative frequencies of precipitates were also investigated and compared with that in the alloy.

In the present work, b’ phase in alloys Al -1.0 mass% Mg2Si -0.5 mass% Ag (Ag-addition) and Al -1.0 mass% Mg2Si (base) was investigated by high resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) to understand the effect of Ag addition. The b’ phase is rod-shape with the longest direction parallel to Al. HRTEM images and SAED patterns obtained along the direction were similar for the b’ phase in both alloys. The unit cell of b’-phase in Ag-addition alloy is hexagonal with the same c-axis dimension as the Ag-free b’, but shorter a-axis. Ag was found in the composition of the rod-shaped precipitates in Ag-addition alloy by energy dispersive X-ray spectroscopy (EDS). In addition, the distribution of Ag was investigated by energy filtered mapping and high annular angular dark field scanning transmission electron microscopy (HAADF-STEM). The Ag-containing atomic column was observed in every b’ unit cell, and the unit cell symmetry is slightly changed as compared with the Ag-free b’. The Ag-containing b’ rods have complicated domain structures. The interfaces of these particles are enriched with Ag atoms that occupy the lattice positions in the Al matrix. The occupancy of the Ag-containing atomic columns seem to vary both inside particles, as well as at the interfaces.

The effect of Al-addition on the mechanism of phase decomposition from β1to α-phase in Cu-40at.%Zn alloy (base alloy) has been investigated using hardness measurement and HRTEM observation. The peak hardness during annealing and the number density of α-phase in Al-bearing alloy were the same as those in the base alloy. Many striations which have been considered as the periodicity of 9R structure were observed in the α-phase of the Al-bearing alloy. There were a lot of striations in the α-phase at the early annealing time, and the number of these striations decreased with increasing annealing time. The rate of decreasing of this striation in the α-phase of the Al-bearing alloy was the slower than that of the base alloy. According to the analysis of the selected area electron diffraction (SAED) patterns obtained for these α-phase in those alloys, transition time from the 9R structure to fcc structure t (9R-fcc) of has been detected by the changing of angles in SAED patterns. The t (9R-fcc) of α-phase in the Al-bearing alloy showed the slower than the base alloy.

Severe plastic deformation (SPD) techniques, such as equal channel angular pressing (ECAP), accumulative roll bonding (ARB) and high pressure torsion (HPT), have been extensively researched to achieve. SPD process makes use of the plastic deformation where no change in the cross-sectional dimension of a work piece occurs during straining. In this work, the effect of HPTed or rolled on aging behavior and microstructure in Al-Mg-Si alloys were investigated. The results show that processing by HPTed leads to significant grain refinement with a grain size of ~400nm. Age-hardening phenomenon was confirmed at 343K and 373K but not at 423K for the HPTed Al-Mg-Si alloy. For the rolled alloys, there is little difference between the hardness of the alloys with different rolling rate. On the other hand, crystal grain with high dislocation density and grain boundary precipitates are observed with TEM analysis for HPTed samples aged at 473 K.

The purpose of study is to investigate aging behavior, crystal structures of metastable phase and relative frequency of metastable phases in aged Al-Mg-Ge and Al-Mg-Ge-Si alloys using high resolution transmission electron microscope (HRTEM), energy dispersive X-ray spectroscopy (EDS) and electron energy-loss spectroscopy (EELS). Every alloy included rod-shaped precipitate which is the same as the typical metastable pahse, ’, in Al-Mg-Si alloy. Except to Mg-rich alloys, the Type-A precipitate, which is a typical metastable phase in the excess Si type Al-Mg-Si alloys and popular at over aged condition, was confirmed as a large rod-shaped precipitates in those alloys. This behavior is probably depends on the ratio of Mg/Ge. Also, Mg, Si and Ge were detected from the ’-phase in Al-Mg-Ge-Si alloy by EELS. This means that the ’-phase in Al-Mg-Ge-Si alloy consists of these 3 elements including Si, not just Ge to form metastable Mg2Ge.

The purpose of this study is to identify the crystal structure of metastable phase in Ag added Al-Mg-Si alloy to compare the formation of β’-phases in Al-Mg-Si alloys without Ag, using images of high resolution transmission electron microscope (HRTEM), selected area electron diffraction (SAED) patterns and an energy dispersive X-ray spectroscopy (EDS). The result of SAED patterns and HRTEM images have been simulated and compared with images then SAED patterns obtained from actual precipitates. SAED patterns and HRTEM images obtained from metastable phase in the Ag added Al-Mg-Si alloy showed similar to those of β’-phase in Al-Mg-Si alloy without Ag and the lattice spacings changed because of the effect of Ag.

It is well known that Ag additional Al-1.0mass%Mg2Si-excess0.4mass%Si alloy (ex. Si-Ag alloy) has higher hardness and elongation than those of Al-1.0mass%Mg2Si-excess 0.4mass%Si alloy (ex. Si alloy). However, precipitation sequence of ex. Si-Ag alloy is not clear yet. In this work, precipitation sequence of ex. Si-Ag alloy has been investigated using high resolution transmission electron microscopy and X-ray energy dispersive spectroscopy. Precipitates were classified into several kinds by HRTEM images and SAED patterns, and relative frequencies of precipitates were also investigated. Its precipitation sequence was compared with that of ex. Si alloy. Type-A, Type-B and Type-C precipitates as special metastable phase in excess Si type Al-Mg-Si alloy, has been observed in ex. Si-Ag alloy, but β’ phase increased and Type-A and Type-B precipitate decreased in this study. Type-A precipitate was found at only grain boundary.

The crystal structure of metastable phase in Ag added Al-Mg-Si alloy was investigated by comparing the β’-phases in Al-Mg-Si alloy without Ag, using images of high resolution transmission electron microscope (HRTEM), selected area electron diffraction patterns (SADPs) and an energy dispersive X-ray spectroscopy (EDS). SADPs and HRTEM images obtained from metastable phase in the Ag added Al-Mg-Si alloy showed similar to those of β’-phase in Al-Mg-Si alloy without Ag and had different lattice spacings because of the effect of Ag. According to our careful analysis on obtained HRTEM images and SADPs, it includes more complicated crystal lattice of distorted hexagons.

The purpose of this study is identity the crystal structure of metastable phase in Ag added Al-Mg-Si alloy by comparing the β’-phases in Al-Mg-Si alloy without Ag, using images of high resolution transmission electron microscope (HRTEM), selected area electron diffraction patterns (SADPs) and an energy dispersive X-ray spectroscopy (EDS). The result of SADPs and HRTEM images have been simulated and compared with images and SADPs obtained from actual precipitates. SADPs and HRTEM images obtained from metastable phase in the Ag added Al-Mg-Si alloy showed similar to those of β’-phase in Al-Mg-Si alloy without Ag and had different lattice spacings because of the effect of Ag.

It has been known that Cu-and Ag-addition Al-1.0mass%Mg2Si alloys (Al-Mg-Si-Cu alloy and Al-Mg-Si-Ag alloy) have higher hardness and elongation than those of Al-1.0mass%Mg2Si alloy. In this study, the aging behaviour of Al-Mg-Si-Cu alloy, Al-Mg-Si-Ag alloy and (Cu+Ag)-addition Al-1.0 mass% Mg2Si alloy (Al –Mg –Si-Cu-Ag alloy) has been investigated by hardness test and TEM observation. The Al-Mg-Si-Cu-Ag alloy has the fastest age-hardening rate in the early aging period and the finest microstructure at the peak hardness among three alloys. Therefore the microstructure of the precipitate in Al–Mg–Si-Cu-Ag alloy has been investigated by HRTEM observation to understand the effect of Cu and Ag addition on aging precipitation.

Conventional time-temperature-parameter (TTP) methods often overestimate long-term rupture life of creep strength enhanced ferritic steels. Decrease in activation energy Q for rupture life in long-term creep is the cause of the overestimation, since the TTP methods cannot deal with the change in Q. Creep rupture data of a heat of Gr.122 steel (up to 26,200 h) were divided into several data sets so that Q was unique in each divided data set. Then a TTP method was applied to each divided data set for rupture life prediction. This is the procedure of multiregion analysis of creep rupture data. The predicted rupture lives have been reported in literature. Long-term rupture lives (up to 51,400 h) of the same heat of the steel have been published in 2013. The multiregion analysis of creep rupture life can predict properly the long-term lives reported. Stress and temperature dependences of rupture life show similar behavior among different heats. Therefore, database on results of the multiregion analyses of various heats of the steel is helpful for rupture life estimation of another heat.

In this study, the early stage of aging in Mg-Gd-Y alloys has been observed by transmission electron microscopy (TEM), high angle annular dark field – scanning transmission electron microscopy (HAADF-STEM) and calculations of images and electron density and bond overlap population (BOP) by first principal to understand the origin of precipitation in this alloy. The small hexagon of 0.37 nm is the first precipitate in this alloy, and this is the evidence of short range ordering of D019 structure. This is referred as the pre β”-phase. In the peak aged condition, β’ phase with bco structure was mainly observed.

The influence of addition of the small amount of transition metals to Al-Mg-Si alloy had reported by many researchers. In the previous our work, β′ phase in alloys Al — 1.0 mass% Mg2Si -0.5 mass% Ag (Ag-addition) and Al -1.0 mass% Mg2Si (base) were investigated by high resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED), in order to understand the effect of Ag. In addition, the distribution of Ag was investigated by energy filtered mapping and high annular angular dark field scanning transmission electron microscopy (HAADF-STEM). One Ag-containing atomic column was observed per β’ unit cell, and the unit cell symmetry is slightly changed as compared with the Ag-free β’. In this work, the microstructure of G.P. zone and β’’ phase was investigated by TEM observation, which were formed before β’ phase. The deformed sample by high pressure torsion (HPT) technique before aging was also investigated to understand its effect for aging in this alloy.

In this study, the aging behaviour of Al-1.0mass%Mg2Si alloys containing 0.35at.%Cu, 0.35at.%Ag, 0.2at.%Cu-0.1at.%Ag and 0.1at.%Cu-0.2at.%Ag has been investigated by hardness measurement and HRTEM observation. 0.35Cu alloy has the highest peak hardness. 0.1Cu-0.2Ag alloy has the shortest aging time to the peak hardness. 0.1Cu-0.2Ag alloy has the fastest age-hardening rate in the early period of aging. 0.35Ag alloy has the finest microstructure at the peak hardness. The precipitates were classified into random-type, parallelogram-type, β’-phase and Q’-phase using HRTEM image with different aging time. Relative frequency of all types of precipitates changed by aging times.

It has been known that Cu- and Ag-added Al-1.0mass%Mg2Si alloys (Al-Mg-Si-Cu alloy and Al-Mg-Si-Ag alloy) have higher hardness and elongation than those of Al-1.0mass%Mg2Si alloy. In this study, the aging behaviour of Al-Mg-Si-Cu alloy, Al-Mg-Si-Ag alloy and (Cu+Ag)-addition Al -1.0 mass% Mg2Si alloy has been investigated by hardness test and TEM observation. The Al-Mg-Si-Cu-Ag alloy has the fastest age-hardening rate in the early aging period and the finest microstructure at the peak hardness among three alloys. Therefore the microstructure of the precipitate in Al-Mg-Si-Cu-Ag alloy has been investigated by HRTEM observation to understand the effect of Cu and Ag addition on aging precipitation.

The present study reports on the correlation between the local atomic structure and cathodoluminescent properties of Al1−xGdxN thin films grown by reactive rf magnetron sputtering at ultra-high vacuum conditions. Those thin films were characterised using X-ray absorption fine structure (XAFS) and cathodoluminescence (CL). From the CL measurements, we have observed a narrow intense ultraviolet emission at 318 nm which is originated from the intra-orbital f-f transition in Gd3+ ions. In order to understand the local atomic structure around the Al1−xGdxN (x=0.1 to 6.0 mol%) thin film, XAFS measurements have been carried out. Analysis of the local atomic structural results showed that both the large distance among Gd atoms and nitrogen vacancies in Al1−xGdxN lattice significantly contribute to the richness in the ultraviolet emission intensity.

We investigated the narrowband ultraviolet emission properties of Al0.94Gd0.06N phosphor thin films pumped by an electron beam. An extremely narrow luminescence line, which was less than 1 nm from the intra-orbital f-f transition in Gd3+ ions, was confirmed at 318 nm. The corresponding emission efficiency was improved by decreasing the growth temperature. The extended X-ray absorption fine structure analysis of the local atomic structure revealed that a low-temperature growth led to the formation of a uniform atomic configuration around Gd, which was found to play a key role in improving the luminescence intensity of the films.

The resolution redistribution (RR) technique can increase the horizontal viewing-zone angle and screen size of electronic holographic display. The present study developed a zone plate method that would reduce hologram calculation time for the RR technique. This method enables calculation of an image displayed on a spatial light modulator by performing additions of the zone plates, while the previous calculation method required performing the Fourier transform twice. The derivation and modeling of the zone plate are shown. In addition, the look-up table approach was introduced for further reduction in computation time. Experimental verification using a holographic display module based on the RR technique is presented.

Mg-Gd-Y alloys are expected that improves mechanical properties at the room temperature. In this study, Mg-Gd-Y alloys have been investigated to clarify the effect of the Gd/Y ratio and total amount of solute atoms on age-hardening and precipitation using HRTEM and calculations of images and electron density by first principal. In the specimen at as-quenched condition, the row contrasts consisted of bright and dark dots lying on the \({\{ 1\bar 100\} _{Mg}}\) planes, which were considered to monolayer were observed.