Chemical state analyses of nitride layers in polycrystalline Ti and Zr produced by high dose implantation of nitrogen ions were performed by X-ray photoelectron spectroscopy (XPS).
From the results, it was found that the accompanying the nitrogen ion implantation, a shift in the peaks of the Ti 2p_3/2 or the Zr 3d_5/2 spectra occurred, and that the magnitude of the binding energy shift became the largest in the vicinity of the maximum concentration of nitrogen. The magnitude of the shift was nearly equal to that of the MeN_1.0 thin films formed under the equilibrium condition.
The change in the binding energy difference between the Ti 2p_3/2 and the N 1s core peaks shifted to the high-energy side along with the increase of the nitrogen concentration. At a nitrogen concentration of about 30%, a clear inflection point was observed. It was shown that this inflection point corresponds to the stoichiometric change from Ti₂N to TiN_1.0.
The N 1s-Zr 3d_5/2 binding energy difference values showed a nearly constant value of 218.3eV in the low concentration region where the nitrogen atoms are α-phase, and it became clear that they decreased continuously from the region for forming ZrN_x (25 at. %N).
It was understood that the change in the binding energy shows a change in the electron charge transfer from metals to nitrides accompanying an increase in the nitrogen atom concentration.
The possibility that nitride thin layers having a gradient composition of nitrogen atoms in the depth direction can be formed was suggested.

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The effects of ultrasonic waves on electrolysis consist mainly of cavitation action. The collapse of cavitation bubbles in liquid generate collapse pressure, jet flow and impulse pressure due to the jet flow. First, the collapsing motion of a single bubble near a electrode surface was modeled, and equations of motion for a spherical gas bubble were obtained. Next, the properties of deposits were studied on the basis of the equations of motion. The jet flow speed and liquid pressure due to water hammer were calculated using equations (33) and (34) in liquid. For example, when the jet flow speed was 120m·s^-1, the water pressure was about 2000atm. The mass transfer on the electrode was a turbulent diffusion. The jet flow and pressure affected the entire mechanism of electrodeposition: the rate of deposition, current efficiency, hardness, roughness and particle size of deposits were affected.

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The hydrogen absorption characteristics of sputtered LaNi_XM_Y alloy films (M=Al and Cu) were examined using a number of electrochemical methods. The alloy films did not exhibit a clear dissociation pressure plateau in P-C-T curves, probably because of their amorphous structure. The hydrogen dissociation pressure of La-Ni alloy films increased as the amount of nickel component increased. In contrast with LaNi_XM_Y films, it decreased with the addition of aluminum or copper. The apparent plateau pressure of the LaNi_5.1 film was estimated to be on the order of 10⁴ Pa at 0°C, and that of the LaNi_4.5 Al_0.1 film was estimated to be on the order of 10² Pa at 20°C. Addition of aluminum was highly effective in lowering the hydrogen dissociation pressure. Elevation of the electrode operation temperature decreased the amount of hydrogen absorbed in the films. However at high temperature the amount of hydrogen that could be discharged galvanostatically increased. It could be considered that this occurred because the improvements in the reactivity and the hydrogen diffusion of alloy films at high temperature resulted in a fall in the discharge overvoltage, and because the dischargeable region of the hydrogen concentration in the P-C-T curves was extended.

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The basic electrodeposition behavior of Dy-Fe was investigated and a rectangular pulse electrodeposition technique was applied to this system. The current efficiency of DC electrodeposition decreased at values over 50A m^-2, but electrodeposition with rather high current efficiency was possible in the pulse current density range between 100 and 250A m^-2. A predominant pulse effect was observed at a duty cycle [on time/[on time+off time]] below 50%. A smooth, bright and amorphous film was obtained using the pulse technique.

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The hardness of deposited silver from cyanide baths was studied. The hardness of the deposited silver was Hv (80∼120) and became Hv (65∼85) after immersion in boiling water for 1h, Hv (50∼65) after baking at 200°C for 1h, Hv (38∼40) after baking at 600°C for 1h and Hv (72) after melting at 1000°C. The hardness of deposited silver from baths containning polyethilen glicol nonylphenol ether showed the same tendencies as deposited silver from cyanide baths with added Se was Hv (138∼148) and became Hv (70∼73) after immersion in boiling water, Hv (51∼55) after baking at 200°C for 1h and Hv (38∼40) after baking at 600°C for 1h. The effectiveness of hardening for silver deposited from a Se-added bath disappeared after immersion in boiling water and due to softening after 3 months.
The hardness of deposited silver from cyanide baths added Sb was Hv (100∼170) and became Hv (95∼165) after dipping in boiling water for 1h, Hv (48∼120) after baking at 200°C for 1h, Hv (38∼40) after baking at 600°C for 1h, and Hv 72 after melting at 1000°C. The hardness of deposited silver from baths with added Sb depended on the current densities, (1∼6A/dm², Hv100∼170).
X-ray diffraction patterns and crystal sizes of deposited silver did not vary at 20∼200°C but did vary at 600°C. with on increased (111) plane and crystal size.
The orientation of the deposited silver became the basic pattern for pure silver (ASTM-Ag), with a crystal size of 29.9nm after melting at 1000°C. It Was presumed that the hardness of deposited silver was decreased by diffusion of H absorbed in the silver and was changed due to crystal transformation or melting at 600∼1000°C. It is possible to adjust the hardness of deposited silver by heat-treatment and by adding hardeners to tho bath.
Various types hardened silver deposits are utilized in functional materials.

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ADC 12 specimens were anodically oxidized in 13M and 1.5M sulfuric acid solutions at constant current densities (i_a), and the amounts of oxide film (W_oxide), Al³⁺ions dissolved in the solutions (W_d) and the volume of gas evolved from the anode (V_g) were examined as functions of the anodizing conditions. There was more W_d, less V_g, and more W_oxide in the 13M solution than in the 1.5M solution. This suggests that the Al³⁺ions dissolution current (i_d) is larger, O₂ gas evolution current (i_g) is smaller, and oxide formation current (i_ox=i_a-(i_d+i_g)) is larger in the 13M solution than in the 1.5M solution. Since the films formed in the 13M solution had larger values of i_ox and larger amounts of SO₄^2-ions than the films formed in the 1.5M solution, the film weight, coating ratio, Si content, and thickness were larger for the 13M-films than for the 1.5M-films. Because microvoids were present in both the barrier layer and the porous walls, the hardness of the 13M-films was lower than that of the 1.5M-films.

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To clarify the mechanism of thermally accelerated adhesion-loss of plated films of tin on phosphor-bronze and to find methods for providing protection from this phenomenon, the effects of the plating conditions of copper undercoat were examined with plating baths of copper cyanide, copper pyrophosphate, copper sulfate and copper fluoborate.
Main findings are summarized as follows.
1. Tin-coated phosphor-bronze is very insensitive to thermal peeling when copper undercoat is electroplated from copper cyanide baths, and very sensitive when the undercoat is electroplated from copper pyrophosphate baths.
2. In electroplated tin-copper systems, no effect of copper electroplating bath on diffusivity was observed.
It is suggested that the susceptibility of thermal peeling of tin-coated phosphor-bronze to plating conditions of copper undercoat is due not to any acceleration of diffusion by the undercoat but to the interaction of vacancies induced in the undercoat by diffusion with the matrix.

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