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The kinetics and mechanism of the full reduction of a thin RuO2film with a stoichiometric (110) surface on Ru(0001) by H2has been studied in the temperature range of 100 to 400 °C at 10–2to 10–4Pa. The reduction kinetics is dominated by the creation of oxygen vacancies and their annihilation upon transformation of RuO2into metallic Ru. The temperature-dependent reduction rate increases linearly with H2pressure. In the temperature range of 100 up to 200 °C, initially hydrogenation of the RuO2(110) surface occurs. Next, oxygen vacancies are created due to desorption of water vapor, which accelerates the reduction by place exchange of oxygen bulk atoms with an activation energy of 0.45 eV. In the temperature range of 200 to 300 °C, slow reduction of RuO2by H2already occurs in the initial period with an activation energy of 0.48 eV and is followed by faster reduction. In the temperature range of 300 to 400 °C, the reduction of RuO2starts immediately when exposed to H2and the activation energy (0.48 eV) is similar to the activation energy in the lower temperature range (100 to 200 °C). Apparently, the annihilation of oxygen vacancies during reduction is more prominent with increasing temperature.

In this review, recent progress in understanding the short- and long-term oxidation behaviour of MCrAlY coatings and the instruments that can be employed to influence the oxidation behaviour of the MCrAlY's are discussed. Particular attention is given on the transient oxidation behaviour, the initial distribution of the Y in the oxide scale and how these initial oxide products influence the long-term oxide growth rate and adhesion of the oxide scale. It is shown that the growth kinetics and adhesion of the oxide scale are to a large extent determined by the amount and distribution of Y-oxide inclusions in the scale. Oxide scales with a regularly distributed and larger amount of small Y-oxide inclusions are faster growing, but perform better in thermal cycle tests. Transient oxide products like metastable aluminas (?- or ?-Al2O3) and spinels do not have a large influence on long-term oxidation behaviour except when the oxide scale is virtually free of Y-oxide inclusions due to Y depletion. Then, the nucleation of metastable aluminas at the onset of oxidation may be beneficial due to a slower-growing alumina scale with a larger grain size after transformation into a-Al2O3. Spinel oxides are considered detrimental for long-term oxidation performance, but their amount is usually limited on most types of MCrAlY coatings due to the rapid supply of Al along the numerous grain and phase boundaries in the fine-grained coatings. The key instruments that can be employed to manipulate the oxidation behaviour of MCrAlY coatings are: (i) the composition and size of the main phases in the coating (?-Ni and ß-NiAl), (ii) the initial distribution, chemical state (metallic or oxidized) and reservoir of the Y in the coating, (iii) the composition and roughness at the coating surface, and (iv) the oxidation conditions during post-deposition heat treatment (temperature, time, partial oxygen pressure).

Gold catalysts supported on TiO2 and TiO2/SiO2 were used in gas-phase propene epoxidation with a hydrogen−oxygen mixture. The catalysts were characterized by ¹⁹⁷Au Mössbauer absorption spectroscopy (MAS), X-ray photoelectron spectroscopy (XPS), and transmission electron microscopy (TEM). Gold particle sizes of 1 wt % Au catalysts calcined at 673 K ranged from 3 to 6 nm. Two Au contributions were found in Mössbauer spectra and assigned to bulk metallic Au atoms in the core of a gold particle and metallic gold on the outer surface of this particle. By MAS, no evidence for charge transfer from support to Au particle could be found. The Auger parameters confirmed that the surface layer of 3−5 nm gold particles is metallic. Deactivation is not due to a change in the active gold species but is related to the TiO2 support. In the preparation via deposition-precipitation, Au(OH)3 species are converted to metallic gold during calcination. Gold particles do not gradually grow during calcination, possibly due to the simultaneous conversion of Au(OH)3 moieties with dehydroxylation of the TiO2 support. Epoxidation activity increases with the amount of surface metallic gold. No evidence for oxidized gold under reaction conditions was found.

A novel technique for the controlled fabrication of π-conjugated oligomers bonded to a substrate is presented. On a gold substrate, monolayers of small monomeric organic units are alternately deposited from solution. They order and connect to the previous layer by self-assembly. In each deposition step, chemisorption takes place on top of the previous monolayer through a condensation reaction. The oligomer chain growth has been analyzed by scanning tunneling microscopy, ellipsometry, and X-ray photoelectron spectroscopy. The azomethine molecules order approximately perpendicular to the gold surface. This growth process allows for flexible preparation of a range of oligomers in a self-assembled monolayer with variable electronic properties.

The incorporation of Kr in sputtered a-Si films has been investigated in a systematic way by varying the Kr to Si flux, yielding Kr concentrations up to 5 at%. Compositions were determined with X-ray microanalysis. A model has been applied to describe the composition of the growing film. The layers were characterized by positron annihilation, Raman spectroscopy and Mössbauer spectroscopy. The present results clearly indicate that ion assisted growth leads to a strong reduction of open volume defects, and that the Kr resides in very small clusters.

TiC layers were chemically vapour deposited at 1273 K mn Fe-C substrates with carbon contents between 0.06 and 1.20w%C. X-ray diffraction stress analyses showed that large compressive stresses are present in the IC coatings and that small tensile stresses any present in the substrates. The stresses developed during cooling from the deposition temperature to room temperature, owing to the difference in thermal shrink between coating and substrate. However, stress relaxation was also evident. This was provoked by the phase transformation processes occuring in the substrate on cooling. Stress relaxation was hindered when grain-boundary cementite formed in the substrates. The stresses present in the TiC coatings on substrates without grain-boundary cementite can be predicted quantitatively.

TiCx coatings were chemically vapour deposited in an industrial reactor on Fe-C substrates with carbon contents between 0.06 and 1.20 wt % C. Electron probe microanalyses showed that significant amounts of chromium and iron were present in the coatings and that chromium was also present in the substrate region adjacent to the coatings. By comparing calculated and measured lattice parameters (corrected for the internal stresses present) it became evident that the chromium was in solid solution in TiCx, whereas the iron was not. This was confirmed by micro Auger electron spectroscopy and X-ray diffraction phase analyses. The carbon to metal ratio,x, of the TiCx coatings decreased with increasing distance to the coating/substrate interface. The effect of iron on the X-ray diffraction line broadening and hardness of the coatings was large (in contrast with the effect of chromium) and increased with increasing distance to the coating/substrate interface because of a decreasing iron particle size. The TiCx crystallite size was small and constant throughout the thickness of the coatings. The chromium present in the substrate region adjacent to the TiCx coatings influenced the microstructure of the substrate by formation of iron, chromium-carbides and reduced the growth rate of the coatings.

Residual stresses in chemically vapour deposited (CVD) TiC and TiN coatings on various substrates were measured using X-ray diffraction. A quantitative prediction of the vaiue of the strain (and stress) was obtained from the difference in thermal shrinkage between layer and substrate on cooling from the CVD temperature to room temperature. The sign of the strains measured always agreed with the prediction, but in a number of cases the absolute value of the elastic strain measured was smaller than that predicted. For layers on iron and Fe-1.0 wt-%C substrates this difference is caused by strain accommodating processes in the substrate, mainly during phase transformations. For TiN layers on molybdenum and tungsten substrates, strain accommodating processes occur in the layer, and these may originate in the microstructure of the layer.