published_or_final_version / Dentistry / Master / Master of Dental Surgery

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Graduation date: 1991

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No description available.

Mg-V and Mg-Zr alloys with nominal compositions 1, 6, 17.5, 27 wt% V and 2, 8.6 and 10.6 wt% Zr respectively were produced by PVD. All deposits exhibited compositional inhomogeneity, columnar microstructures and a strong basal texture. The solid solubilities of V and Zr in Mg were extended approximately to 17 wt% V and 10 wt% respectively. Grain refinement occurred with increasing solute content. The solid solution break up temperature decreased as the V and Zr content in the alloys increased. Pure V precipitated when the extended solid solubility of was exceeded. Both c and a lattice parameters, as well as the c/a ratio decreased with increasing V content in the Mg-V alloys. The slight increase of the a-lattice parameter and the decrease of the c one led to a decrease of the c/a ratio with increasing Zr additions in the Mg-Zr alloys. The air-formed oxide on the surfaces of the Mg-V alloys consisted predominantly of hydromagnesite at the outermost surface with Mg(OH)2 in excess of MgO underneath. No evidence of V oxide in the surface film was found. Magnesium oxide was also found between the grains of the deposits. The air-formed oxide on the surfaces of the Mg-Zr alloys consisted of ZrO2, MgO and possibly Zr sub-oxide. The presence of the oxides beween the columnar grains gave rise to graded metal/oxide interfaces. The outermost surfaces of the Mg-Zr alloys were similar to the Mg-V ones. Analysis of changes of the Auger parameters of the Mg-V and Mg-Zr alloys was also undertaken in order to investigate the electronic changes that take place upon alloying Mg with V and Zr. Charge transfer between 0.09 and 0.11 electrons/atom from Mg to V as well as changes in the V d charge were calculated by measuring the Mg and V Auger parameters and using the charge transfer model of Thomas and Weightman. Electron transfer between 0.02 and 0.03 electrons/atom from Mg to Zr was also found to occur upon alloying Mg with Zr. The electron transfer has been related to changes in crystal structure. The Mg-V and Mg-Zr alloys were examined after immersion in 3 wt% NaCl solution for 5 and 15 minutes, 9 hours and 7 days. The dramatic increase in the corrosion rate of the Mg-V alloys was attributed to the precipitation of pure V. The unsatisfactory corrosion performance of the Mg-V alloys was attributed to the absence of compositional uniformity through the thickness of the Mg-V deposits and the low thermodynamic stability of the corrosion products in the saline environment. Hydromagnesite at the outermost surface and Mg(OH)2, MgO and V2O4 in the bulk of the corrosion layer were the corrosion products. MgH2 and areas enriched in metallic V within the bulk of the corrosion products were also detected. The low corrosion rates of the Mg-Zr alloys, the lowest ever reported for Mg alloys, were attributed to the nature of the corrosion products and particularly the Zr contribution. The corrosion products were enriched in Zr, and were non-porous and in many cases well adherent. X-ray and electron diffraction suggested the existence of only Mg(OH)2 and MgO in the corrosion products, indirectly implying the participation of zirconium oxide/hydroxide in an amorphous/nanocrystalline state. Surface analysis indicated that a Zr oxide coexisted with Mg(OH)2 and MgO below a magnesium carbonate overlayer and also suggested the existence of Zr hydrous oxide (hydroxide). The repetition of the substrate pattern, as well as the fact that Zr hydroxide was replaced with ZrO2 and Zr sub-oxide as the metal-oxide interface was approached, implied a corrosion mechanism involving inwards diffusion of the anionic species.

The stress corrosion behaviour of precipitation hardened Al-9Mg, Al-22Zn and Al-3Mg-6Zn alloys has been studied in aqueous environments and ethanol. The stress corrosion susceptibility defined as the reciprocal of failure time has been investigated as a function of alloy-environment system, isothermal aging treatment, microstructure, applied tensile stress, and temperature using smooth and notched specimens. Constant load tests, load-relaxation tests and tensile tests in different environments have been used to evaluate the stress corrosion characteristics of aluminum alloys. A limited study of Mg-9Al has also been carried out in aqueous environments. The process of stress corrosion generally consisted of three parts: 1) A slow initiation stage 2) a rapid propagation stage 3) mechanical fracture due to tensile overload. With a few exceptions, the initiation time was greater than the propagation time. The crack initiation and propagation rates were stress and thermally activated and could be expressed by a general equation of the form Rate = [formula omitted] where α is the applied tensile stress, Q is the apparent activation energy of the rate controlling process and A(0) and n are constants for a given alloy-environment system. The apparent activation energy of the rate controlling process was different in the two environments. It also changed between initiation and propagation stages. The aluminum alloys when ranked in order of increasing susceptibility were: 1) Al-3Mg-6Zn, 2) Al-9Mg, 3) Al-22Zn. The alloys which were given heat treatments correlating to the presence of coherent or partially coherent phases, were found to be most prone to stress corrosion cracking. The environments placed in an order of increasing aggressiveness were dessicant-dried air, double distilled water, ethanol, ambient air, deionized water and NaCl/K₂CrO₄solution. The ductility of susceptible aluminum alloys was found to be significantly decreased by NaCl/K₂CrO₄and deionized water at low strain rates and enhanced by dessicant-dried air. Fractography showed the cracking to be intergranular in aluminum alloys and transgranular in the Mg-Al alloy. The stress corrosion surface was characterised by a rough or corroded appearance while the mechanically fractured surface exhibited slip steps and dimples caused by void formation. The hydrogen mechanism of cracking was examined in light of hydrogen charging experiments and other evidence and was found to be unsatisfactory. Models involving either dissolution or deformation alone were also inadequate in explaining the present results. Therefore a new model was postulated which involves the generation of a continuous path of chemical heterogeneity by shearing and link up of coherent precipitates followed by their dissolution. The rate controlling step in the deformation process is believed to change during the transition from initiation to propagation. The postulated model is consistent with the present results but its further development must await better knowledge of the kinetics of dissolution of precipitates and that of deformation processes at the crack tip. / Applied Science, Faculty of / Materials Engineering, Department of / Graduate

Single crystals and bicrystals have been used to study pitting corrosion and intergranular corrosion of Al and Al-Cu alloys in 0.5M NaCl solution. The critical pitting potential and pit density were examined as a function of a number of factors. These included crystallographic orientation; the bulk solution chemistry including CI- concentration, NO₃- addition and pH; the effect of Cu alloying; and the effects of homogenizing and aging on the alloy crystals. The susceptibility for pitting corrosion was found to depend on crystallographic orientation in pure Al with {111} showing maximum pitting and {011} and {001} exhibiting progressively less pitting. This crystallographic effect was not observed in the Al-3 wt% Cu alloy. The addition of Cu to pure Al was found to raise the Epit and produce a higher pit density on the surface. The increase of CI⁻ concentration was found to enhance pitting corrosion, producing a higher pit density and lowering the Epit. Addition of NO₃- to the solution decreases pitting corrosion, reduces the pit density and substantially shifts the Epit to a more noble potential. A model of pitting corrosion is proposed, based on a local kinetic balance between the repassivation process and the dissolution process at the bare metal surface at the base of a preexisting oxide flaw on the crystal surface. The model successfully accounts for the observed effects of the Cu alloy addition, and the solution composition variations on pitting corrosion. In the alloy bicrystals, it was observed that pitting corrosion in the grain boundary region was dependent on the composition and thermal history of the crystal. In most of the homogenized Al-Cu bicrystals, the presence of the grain boundary did not influence the pitting corrosion. In a 0.1 wt% Cu alloy with a tilt boundary of 28° about the <001> direction preferential pitting along the grain boundary was observed. The preferential pitting is attributed to nonequilibrium depletion of Cu at the high angle tilt boundary. Preferential attack is also observed at grain boundaries in as-grown and in aged bicrystals. This is attributed to Cu segregation in the crystals and the lower value of Epit associated with the Cu depleted regions. / Applied Science, Faculty of / Materials Engineering, Department of / Graduate

by Ka-Man Mak. / Thesis (M.Phil.)--Chinese University of Hong Kong, 1998. / Includes bibliographical references. / Text in English; abstract also in Chinese. / by Ka-Man Mak. / Acknowledgement --- p.i / Abstract --- p.ii / List of tables --- p.v / List of figures --- p.vi / Table of contents --- p.xi / Chapter Chapter One --- Introduction --- p.1-1 / Chapter 1.1 --- History of intermetallics --- p.1-1 / Chapter 1.2 --- Properties of intermetallic compounds --- p.1-5 / Chapter 1.2.1 --- Magnetic properties --- p.1-5 / Chapter 1.2.2 --- Chemical properties --- p.1-6 / Chapter 1.2.3 --- Semiconducting properties --- p.1-7 / Chapter 1.2.4 --- Superconducting properties --- p.1-7 / Chapter 1.2.5 --- Hydrogen storage --- p.1-8 / References --- p.1-9 / Chapter Chapter Two --- Background --- p.2-1 / Chapter 2.1 --- Some Behaviours of Intermetallics / Chapter 2.1.1 --- Intergranular and cleavage fracture --- p.2-1 / Chapter 2.1.2 --- Corrosion --- p.2-3 / Chapter 2.1.3 --- Oxidation in high-temperature intermetallics --- p.2-5 / Chapter 2.1.4 --- Hot corrosion --- p.2-8 / Chapter 2.2 --- Nickel aluminides --- p.2-9 / Chapter 2.2.1 --- Ni3Al --- p.2-10 / Chapter 2.2.2 --- NiAl --- p.2-12 / References --- p.2-14 / Chapter Chapter Three --- Oxidation --- p.3-1 / Chapter 3.1 --- Introduction --- p.3-1 / Chapter 3.2 --- Specimens preparation --- p.3-1 / Chapter 3.3 --- Experiment process --- p.3-5 / Chapter 3.3.1 --- Instrumentation --- p.3-5 / Chapter 3.3.2 --- Choosing of experimental temperature --- p.3-9 / Chapter 3.3.3 --- Methodology --- p.3-9 / Chapter 3.4 --- Results and Discussions --- p.3-10 / Chapter 3.4.1 --- Dependence of time --- p.3-10 / Chapter 3.4.2 --- Dependence of temperature --- p.3-14 / Chapter 3.4.3 --- Dependence of composition --- p.3-15 / Chapter 3.4.4 --- Activation energy of oxidation --- p.3-15 / Chapter 3.4.5 --- Oxidation morphology and mechanism --- p.3-16 / Chapter 3.5 --- Conclusions --- p.3-20 / References --- p.3-21 / Chapter Chapter Four --- Hot corrosion --- p.4-1 / Chapter 4.1 --- Introduction --- p.4-1 / Chapter 4.2 --- Specimens preparation --- p.4-1 / Chapter 4.3 --- Experiment process --- p.4-3 / Chapter 4.3.1 --- Instrumentation --- p.4-3 / Chapter 4.3.2 --- Choosing of experimental environment and temperature --- p.4-5 / Chapter 4.3.3 --- Methodology --- p.4-6 / Chapter 4.3.3.1 --- Experiment --- p.4-6 / Chapter 4.3.3.2 --- Experimental setup --- p.4-8 / Chapter 4.4 --- Results and discussions --- p.4-9 / Chapter 4.4.1 --- Dependence of time --- p.4-9 / Chapter 4.4.2 --- Dependence of temperature --- p.4-10 / Chapter 4.4.3 --- Comparison between hot corrosion with oxidation --- p.4-11 / Chapter 4.4.4 --- Dependence of composition --- p.4-12 / Chapter 4.4.4.1 --- Comparison between S1 and S2 --- p.4-12 / Chapter 4.4.4.2 --- Comparison between S3- S7 --- p.4-12 / Chapter 4.4.5 --- Results from XRPDS --- p.4-13 / Chapter 4.4.6 --- Study of microstructure --- p.4-13 / Chapter 4.4.6.1 --- Dependence on time --- p.4-14 / Chapter 4.4.6.2 --- Dependence on temperature --- p.4-14 / Chapter 4.4.6.3 --- Dependence on composition --- p.4-14 / Chapter 4.5 --- Corrosion mechanism --- p.4-15 / Chapter 4.5.1 --- Chemical reactions --- p.4-15 / Chapter 4.5.2 --- Corrosion process --- p.4-16 / Chapter 4.5.2.1 --- Temperature effect --- p.4-16 / Chapter 4.5.2.2 --- Composition dependence --- p.4-17 / Chapter 4.6 --- Conclusions --- p.4-18 / References --- p.4-19 / Chapter Chapter Five --- Conclusionsand suggestions for further studies --- p.5-1 / Chapter 5.1 --- Oxidation --- p.5-1 / Chapter 5.2 --- Hot corrosion --- p.5-2 / Chapter 5.3 --- Further development --- p.5-3