The composition ranges of Zr–Cu–B and Zr–Cu–Ga amorphous alloys and their supercooled liquid region (ΔT_x) and reduced glass transition temperature (T_g⁄T_m) were examined in comparison with those for Zr–Cu–Al amorphous alloys. The glass formation range, ΔT_x and T_g⁄T_m decrease in the order Al>Ga>B. However, the addition of a small amount (5 at%) of B or Ga for the Zr–TM (TM=Co, Ni, Cu) base alloys is effective for the increase in ΔT_x and the largest ΔT_x value is 65 K for Zr₆₅Ni₁₀Cu₂₀B₅ and 73 K for Zr₆₅Co₅Cu₂₅Ga₅ which are larger than that (60 K) for Zr₆₅Ni₁₀Cu₂₀Al₅. The crystallization of the B- or Ga-containing alloy takes place through a single stage due to the polymorphous precipitation of bct Zr₂(Ni, Cu, B) or bct Zr₂(Co, Cu, Ga) while the Al-containing alloy shows a two-stage process of Am→Am′+metastable cubic Zr₂(Ni, Al)→bct Zr₂(Ni, Al)+bct Zr₂(Cu, Al). The addition of B, Ga or Al to the Zr₆₅(Co, Ni, Cu)₃₀Al₅ alloy causes the further increase in ΔT_x to 84, 90 and 104 K respectively, and the crystallization process consists of the polymorphous-type single stage. The reason why the Zr–Cu–B ternary amorphous alloys have the lower glass-forming ability and smaller ΔT_x value among the three Zr–Cu–M (M=B, Ga or Al) ternary systems is presumably because Cu–B pair has a positive heat of mixing and the atomic size ratios between B and Zr or Cu are too large to construct a higher degree of dense random packed structure. This is consistent with the previous result that the dissolution of Al with large negative heats of mixing and different atomic size ratios causes the formation of a higher degree of dense random packed structure and the ΔT_x value increases further for the Zr–Cu–Al amorphous alloys containing a larger amount of Al. These results support the previous concept that the simultaneous satisfaction of the significantly different atomic size ratios and large negative heats of mixing is necessary for the formation of an amorphous alloy with a wide supercooled liquid region before crystallization.

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The glass transition temperature (T_g) and crystallization temperature (T_x) of the Zr₆₅Al₁₀Cu₁₅Ni₁₀ base glassy alloys containing additional M (M=Ti, Hf, V, Nb, Cr, Mo, Fe, Co, Pd or Ag) elements were examined as a function of M elements, with the aim of finding an effective element for the increase in ΔT_x (=T_x−T_g) and of confirming the appropriateness of the previous empirical rules for the appearance of large ΔT_x. As the additional amount of the M elements except Hf increases, T_g increases gradually, whereas T_x decreases significantly and leads to the decrease in ΔT_x. No effective M element leading to the increase in ΔT_x is found. The ineffectiveness is attributed to the partial generation of repulsive bonding nature of Cu–M (M=V, Nb, Cr, Mo, Fe, Co, Pd or Ag) and Zr–M (M=Ti of Hf) pairs which does not satisfy the empirical rules. The area ratio of the glassy region in the arc-melted ingots with a maximum thickness of about 8 mm was found to increase from 67% for the Zr–Al–Cu–Ni alloy to about 90% for the Zr-based alloys containing 5%Ti, 2.5%Nb or 5%Pd, though ΔT_x decreases significantly by the addition of these elements. The great effectiveness of the three elements on the glass-formation tendency in the arc-melted ingots is interpreted to originate from the suppression of the growth reaction of crystalline nulcei which pre-exist in the arc-melted alloy. Furthermore, the disagreement between the glass-formation tendency evaluated by ΔT_x and the area ratio of the glassy region is thought to result from the difference in the dominant factors which are the crystalline nucleation and growth reactions for the ΔT_x of the melt-spun glassy alloys and the growth reaction for the area ratio of the glassy region in the arc-melted ingot. The finding of the additional elements leading to the increase in the glass-formation tendency of the arc-melted alloys, regardless of the magnitude of ΔT_x, seems to be a very encouraging event for future development of bulk glassy alloys.

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Bulk glassy Fe₇₃Al₅Ga₂P₁₁C₅B₄ alloys in cylindrical form with diameters of 0.5 and 1.0 mm were found to form by a copper mold casting method. The further increase in diameter causes the formation of coexistent glassy, Fe₃(B, C), Fe₂B and Fe₃P phases for the 1.5 mm^φ sample and coexistent Fe₃(B, C), Fe₂B and Fe₃P phases for the 2.0 mm^φ sample. It is to be noticed that the maximum thickness for glass formation is about 10 times larger than the largest thickness for Fe-based glassy alloys reported up to date. The glass transition temperature (T_g), crystallization temperature (T_x) and heat of crystallization of the 1.0 mm^φ glassy alloy are 732 K, 785 K and 3.76 kJ/mol, respectively. No appreciable difference in the thermal stability and magnetic properties is seen between the bulk glassy alloys and the melt-spun ribbon. The 1.0 mm^φ glassy alloy has ferromagnetism with a Curie temperature of 606 K and exhibits 1.26 T for saturation magnetization (B_s), 82 A/m for coercivity (H_c) and 0.38 for the ratio of residual magnetization to B_s at room temperature. The large ΔT_x(=T_x−T_g) and large glass-forming ability can be obtained for the Fe-based alloy containing simultaneously the five solute elements. The effectiveness of the multiplication is presumably due to the combination of the following three effects; (1) the suppression of crystalline nuclei due to the increase in dense random packing density for the glassy structure containing P, C and B with significantly different atomic sizes, (2) the difficulty of atomic rearrangements for the precipitation of the Fe-metalloid compounds caused by the generation of Al-metalloid pairs with strongly attractive bonding nature, and (3) the decrease in the preferential precipitation tendency of Fe–B and Fe–C compounds by the dissolution of Ga which is immiscible to B and C and soluble to Fe.

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The anomalous X-ray scattering (AXS) method using Cu and Mo K absorption edges has been employed for obtaining the local structural information of superionic conducting glass having the composition (CuI)_0.3(Cu₂O)_0.35(MoO₃)_0.35. The possible atomic arrangements in near-neighbor region of this glass were estimated by coupling the results with the least-squares analysis so as to reproduce two differential intensity profiles for Cu and Mo as well as the ordinary scattering profile. The coordination number of oxygen around Mo is found to be 6.1 at the distance of 0.187 nm. This implies that the MoO₆ octahedral unit is a more probable structural entity in the glass rather than MoO₄ tetrahedra which has been proposed based on infrared spectroscopy. The pre-peak shoulder observed at about 10 nm⁻¹ may be attributed to density fluctuation originating from the MoO₆ octahedral units connected with the corner sharing linkage, in which the correlation length is about 0.8 nm. The value of the coordination number of I⁻ around Cu⁺ is estimated as 4.3 at 0.261 nm, suggesting an arrangement similar to that in molten CuI.

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The long-period ordered phase (LPOP) formed in Fe-1.83 mass%C martensite during aging at room temperature has been investigated by the electron diffraction method. It was found that the LPOP formed at 298 K possessed an incommensurate superperiod of 10.7c (c is the length of the c axis of martensite), different from the previously reported 12c at higher aging temperatures. This indicated a temperature dependence of superperiod of the LPOP, which could not be explained by previous structure models of the LPOP. The above experimental finding show that the LPOP is consistent with an off-stoichiometric phase γ′-Fe_xC(II) (x=4∼10), the superperiod of which can adopt 10c, 12c, and 14c at different temperatures. The coexistence of two of these structures at intermediate temperatures results in the nominal incommensurate superperiod. Computer simulated electron diffraction (ED) patterns based on this structure model remarkably resembled the observed ED patterns in every aspect.

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Precipitation in Al–Li–Cu ternary alloys has been investigated mainly by a resistivity study. The precipitation sequence in this alloy is as follows: supersaturated solid solution→δ′→T₁. The resistivity decreases at first and then decreases again passing through a peak. This anomalous behaviour has been interpreted using transmission electron microscopy and it is suggested that the resistivity decreases accompanying the precipitation of δ′ (Al₃Li), while it increases with precipitation of the T₁ (Al₂CuLi) phase and then decreases through a maximum. The time taken to arrive at the maximum in resistivity is comparable with the time to peak hardness, which suggests that the increase in resistivity may be attributed to the strain field around T₁ precipitates. Contrast experiments by transmission electron microscopy show the strain field due to the T₁, which disappears on prolonged aging.

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In order to investigate dynamic events concerning atomic behavior of phase transformation and precipitation, in-situ high resolution electron microscope (HREM) observation methods are useful. In-situ HREM observation and observation by dark-field imaging have been carried out on the formation and growth of aged ω-phase in a Ti15Mo5Zr alloy single crystal. The effect of stresses on the formation of the ω-phase has also been investigated by HREM observations. The main results obtained are as follows; (1) ω-phase crystals in thin foil specimens, which were aged for 100 ks at 623 K, did not appear. However, ω-phase crystals appeared in a specimen, which was aged in the bulk state, thinned and then aged. Also, when a specimen was deformed by rolling and then aged for 3.6 ks, ω-phase crystals appeared in the whole of a deformation band. ω-phase appeared in a specimen deformed by tension. These results suggest that some kinds of stresses are necessary to form the aged ω-phase. (2) The results of in-situ HREM observation on the formation of aged ω-phase crystals suggest that the ω-phase grows by the formation of atom-pairs with 0.28 nm distance, which are formed by movement of two neighboring atom lines in opposite ⟨111⟩ directions as required by the model.

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Reaction diffusion in the Mg–Cu system has been investigated by an electron-probe microanalyzer. Two kinds of intermetallic compounds Mg₂Cu and MgCu₂ in the Mg–Cu diffusion couples were observed in the temperature range from 683 to 748 K. The Mg₂Cu phase layer grows faster than MgCu₂ phase layer. The growth rate of the Mg₂Cu layer obeys a parabolic law as a function of annealing time and thus the growth turned out to be diffusion-controlled. The temperature dependence of the growth constant k of the Mg₂Cu layer can be described with the Arrhenius equation:
k_Mg2Cu=1.5×10⁻³exp[−(156±10)kJmol⁻¹⁄RT]m²s⁻¹.
Kirkendall marker shift toward the Cu-rich side was found, which suggests that diffusion of Cu is faster than that of Mg.

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The deformation mechanisms of a fine grained Al-14 mass%Ni-14 mass%Mm [Mm=misch metal] crystalline alloy consolidated from its amorphous powders have been discussed from the mechanical data examined previously at strain-rates between 10⁻³ and 100 s⁻¹ at temperatures from 773 to 898 K. This alloy exhibits superplasticity at unusually high strain rates of nearly 1 s⁻¹ in the temperature range from 848 to 885 K, which is close to the measured melting point. By incorporation of the temperature dependence of grain size, threshold stress and shear modulus into the constitutive equation, the activation energy is found to be 158 kJ mol⁻¹ at temperatures below the melting point. In the temperature range above the melting point, however, the activation energy is very large at more than 350 kJ mol⁻¹. This activation energy in the temperature range below the melting point is similar to that for lattice self-diffusion of aluminum, and all mechanical data in these temperatures ranges can be represented by a single equation. It is postulated that superplastic flow in the Al–Ni–Mm crystalline alloy is controlled by a grain boundary sliding mechanism accommodated by dislocation climb controlled by lattice self-diffusion at the solid state, and in the temperature range higher than the melting point is accommodated by the liquid phase at grain boundaries and/or interfaces. The further deformation mechanisms in high strain rate superplasticity should take the phase state of liquid and solid into account.

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Internal friction in aluminum thin films 0.2 to 2.0 μm thick on silicon substrates has been investigated between 180 and 360 K as a function of strain amplitude by means of a free-decay method of flexural vibration. According to the constitutive equation, the internal friction in the film alone can be evaluated separately from the data on the film/substrate composite. The amplitude-dependent part of internal friction in aluminum films is found in the strain range approximately two orders of magnitude higher than that for bulk aluminum. On the basis of the microplasticity theory, the amplitude-dependent internal friction can be converted into the plastic strain as a function of the effective stress on dislocation motion. The mechanical responses thus obtained for aluminum films show that the plastic strain of the order of 10⁻⁹ increases nonlinearly with increasing stress. These curves tend to shift to a higher stress with decreasing film thickness and also with decreasing temperature, both indicating a suppression of the microplastic deformation. At all temperatures examined, the microflow stress at a constant level of the plastic strain varies inversely with the film thickness, which qualitatively agrees with the variation in macroscopic yield stress.

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Differential pulse voltammetry of Cr⁶⁺/Cr³⁺ in alkali silicates, which has been worked in the previous study (Ref. (11)) was extended to silicate melts containing alkalin earth oxides. The basicity scale of new systems finds more relevant applications to practical slags. As compared at the disilicate composition, the basicity increase in the order of MgO, CaO and Li₂O, BaO, Na₂O, K₂O, Rb₂O, and Cs₂O.

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A model to predict the solidification of spheroidal graphite (SG) cast iron during cooling process was developed. The model treated the problem of the ternary Fe–C–Si system and aimed at simulating the evolution of the solidifying microstructure. When the eutectic temperature is reached, eutectic grains are formed. The rate of nucleation depends on the melt characteristics and supercooling. The nucleation rate was assumed to be presented by power law function of undercooling. The growth of the austenite grain was controlled by carbon diffusion through austenite. Carbon and silicon concentrations at different interfaces are calculated from the Fe–C–Si equilibrium ternary phase diagram and Sheil equation was applied to calculate the silicon content in the bulk liquid and the concentration of carbon in the liquid was determined to keep the eutectic reaction.
The model was applied to a four-step cylindrical sand casting which has a different diameter. The simulated results were compared with the experimental ones using a post processing program which was also developed to simulate the microstructural evolution graphically. The results showed relatively good agreement and the usefulness of the proposed model was confirmed.

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