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Made available in DSpace on 2014-10-09T12:36:50Z (GMT). No. of bitstreams: 0 / Made available in DSpace on 2014-10-09T13:56:00Z (GMT). No. of bitstreams: 1 04366.pdf: 5734662 bytes, checksum: 226af9141e1984ce9dff1230b5f437ed (MD5) / Tese (Doutoramento) / IPEN/T / University of Manchester, England

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

Crevice corrosion experiments have been conducted on Inconel 600 and Inconel 625 exposed to two principle test solutions of 1 M NaCl and 1 M NaCl + 0.01 M Na₂S₂0₃ (Sodium Thiosulphate) at three temperatures, 22°C, 55 °C and 80°C. The crevice corrosion tests were performed in a corrosion cell which was constructed from PTFE (Polytetrafluoroethylene, Teflon) and Pyrex glass. Features of the cell included the utilization of an artificial Teflon-metal crevice and provisions to monitor crevice corrosion current, active crevice corrosion potential and active crevice pH. Additional experiments included potentiodynamic anodic polarization tests on pure Ni, Alloy 600, and Alloy 625 in bulk solution environments and in simulated crevice solutions. Crevice corrosion morphology and compositional analysis of the corrosion products was studied using a scanning electron microscope equipped with an X-ray energy dispersive spectroscopy (EDS) system. Results show that crevice corrosion rates increase with increasing temperature for Alloy 600 in both principle test solutions. X-ray EDS analysis indicated that an insoluble nickel sulphide corrosion product formed on Alloy 600 in a solution of 1 M NaCl + 0.01 M Na₂S₂0₃. For the Alloy 600, in a solution of 1 M NaCl + 0.01 M Na₂S₂0₃, initiation times were significantly reduced and crevice corrosion propagation rates enhanced, as compared to Alloy 600 in 1 M NaCl. The decrease in initiation times has been attributed to the destabilizing nature of the S₂O₃⁻² species on the passive oxide film. Enhanced propagation rates have been attributed to the presence of H₂S in the crevice solution and the formation of an adsorbed species Ni(H₂S)ads which enhances the anodic dissolution reaction. The H₂S in the active crevice solution originated from the thermodynamically favoured electrochemical reduction of the S₂0₃⁻² species in the active crevice solution. Experiments on Alloy 625, which is alloyed with molybdenum, (Mo), show that it was virtually immune to crevice corrosion as compared to Alloy 600 which is not alloyed with Mo. The resistance of Alloy 625 to crevice corrosion initiation has been attributed to the stabilizing nature of MoO₂ in the passive oxide film. For an actively corroding system, the formation of the molybdate species MoO₄⁻² may act as an anodic inhibitor and effectively enhance the repassivation of the passive film. / Applied Science, Faculty of / Mining Engineering, Keevil Institute of / Graduate

Circumferentially notched rods of three types of austenitic stainless steel were stress corroded under freely corroding conditions at their yield stress in boiling 154°C MgCl2, and boiling MgCl2 with additions of HCl, CoCl2, and FeCl3. Alloy types 304, 316, and 310 were chosen because of their known different stress corrosion susceptibility. The corrosion products formed on the stress corrosion fracture surface were analyzed by electron diffraction and energy dispersive x-ray analysis. Electron diffraction of corrosion products, both in situ and stripped from the fracture surface, showed the corrosion product was a spinel oxide in all cases. Qualitative x-ray analysis of corrosion products, in situ and stripped from the fracture surface, indicated the corrosion product formed on all alloys was enriched in chromium and contained lesser amounts of the elements iron, nickel, silicon, molybdenum, magnesium, phosphorous and chlorine. Observations led to the conclusion that the corrosion product formed in the cracks of the various alloys was similar, being predominantly a chromium enriched oxide spinel with possible traces of metal chlorides or a corundum type oxide. The presence of the spinel oxide was consistent with anticipated E-pH equilibrium within the crack. However, it was felt the variation in s.c.c. behaviour between the different alloys could not be adequately accounted for in terms of the composition of the oxide. / Applied Science, Faculty of / Materials Engineering, Department of / Graduate

A thesis submitted to the Faculty of Engineering and Built Environment, University of the Witwatersrand, in fulfilment of the requirements for the degree Doctor of Philosophy (Materials and Metallurgy) Johannesburg, February 2019 / The first corrosion map of South Africa (SA) was published in 1991. Since then only minor variations have been made to this map. However, due to the lack of differentiation of inland locations, overstatement of the corrosivity of environments, changes in international standards regarding the measurement of corrosive atmospheres, increased industrialisation of SA’s large metropolitan areas, global climate change effects and improved mapping techniques, this chart has become outdated. The present study focused on the development of a new corrosion map for SA, with the emphasis placed on the provision of more detail concerning the atmospheric corrosivity of the Greater Johannesburg Metropolitan Area (GJMA) – SA’s economic heartland. In the study, historical published and unpublished corrosion data for South Africa were assimilated and analysed. The atmospheric corrosivity of the GJMA was also measured using ASTM G116 wire-on-bolt and ISO 9226:1992 open aluminium, zinc and copper wire helix devices, including ISO 9226:2012 flat mild steel and hot-dip galvanised steel coupons. Coated mild steel specimens were also exposed to determine the corrosive effects of the GJMA’s atmosphere on organic materials. The morphologies and chemical compositions of the mild steel corrosion products, using SEM-EDS, FTIR and Raman spectroscopy, as well as pH and water-soluble salt measurement techniques were furthermore investigated, along with correlations between the measured corrosion data, and general meteorological and pollution parameters for the GJMA area. The study showed that the corrosivity of SA’s coastal environments decreases rapidly within the first 150 m from the ocean and that for most inland locations, very low (C1) to potentially high (C4) corrosive conditions may be expected, as per the ISO 9223 rating scheme. High correlations were also found between the first-year corrosion rates of mild steel, hot-dip galvanised steel, zinc, aluminium and copper. Moreover, it was confirmed that the corrosivity of the GJMA is influenced by precipitation, humidity, PM10 and PM2.5 particulate matter, as well as SO2, NO2, NO, CO and O3 levels in the atmosphere, including wind direction and daily temperatures (maximum and minimum). Other factors found to affect the corrosivity of the GJMA’s atmosphere are: elevations above sea and ground level; the presence of vegetation and large water bodies; topography; shielding and shading effects; the occurrence of an El Niño or La Nina event; and acid rain. Moreover, it was established that the GJMA is most corrosive during spring and summer and that at least 90% of the area can be rated Upper-C2 (low to medium) corrosive. The results furthermore revealed high correlations between the ISO 9223 (1992 and 2012) corrosion monitoring devices and that wire-on-bolt (CLIMAT) units are better indicators of the impact of atmospheric pollutants on the 12- month corrosion rate of hot-dip galvanised steel than uncoated steel. A strong linear correlation was also found regarding the average corrosion rate of mild steel with every 20 mg/m2 rise in the concentration of water-soluble salts in the corrosion product. The East Rand of the GJMA rated most corrosive, with corrosion trouble spots identified at Aeroport, Bonaero Park, Brakpan, Chloorkop, Dalpark, Dunnottar, and Nigel. The Kagiso-Randfontein area was the only other area, outside the East Rand, that rated more corrosive. Extrapolations of the long-term corrosion rates of mild steel and hot-dip galvanised steel in the GJMA were additionally made based on logarithmic regressions of the 6-, 12-, 18- and 24-month corrosion data. Finally, geoprocessed (metal specific) corrosion maps were developed for the GJMA, subsequently incorporated into several corrosion maps for SA (also metal specific), to provide better clarity regarding SA’s inland areas. Keywords: Corrosion, atmosphere, South Africa, Greater Johannesburg, inland Metropolitan, wire-on-bolt, CLIMAT, ISO 9223, mapping, pollution, mild steel, hot-dip galvanised steel, aluminium, zinc, copper, coatings / E.K. 2020

Student Number : 9900051W - MSc dissertation - School of Chemical and Metallurgical Engineering - Faculty of Engineering and the Built Environment / When specifying materials for use in exhaust systems, it is imperative that they exhibit sufficient corrosion resistance for the specific conditionsto which exhaust components are exposed, since up to 80% of all failures is attributed to corrosion and oxidation. It is therefore neccesary to establish the corrosion behaviour of the materials in conditions and environments to which the exhausts would typically come into contact with. Most car manufacturers, exhaust manufacturers and material providers have specific corrosion testing methods which they use to determine the corrosion resistance of candidate materials, but there appears to be no standard procedure. A summary comparing all the existing systems is given in section 2.7. The corrosion testing methods utilise a wide range of conditions, testing temperatures and stages. However, careful investigation of the tests show some similarities, and it was possible to identify eleven key tests, that cover internal corrosion, external corrosion and oxidation for both diesel and petrol engines. Eight of these tests were used to rank the corrosion and oxidation resistance of selected stainless steels, namely AISI type 304, 321, 409, 434 and DIN 1.4509. It appears that the austenitic stainless steels perform better in the cold end conditions, while the ferritic types are more resistant in the hot end high temperature conditions. Of all the eight test performed, only the electrochemical tests for external corrosion of cold end components did not give reproducible results. The rest of the tests could be used to screen materials for exhaust system applications. In the internal condition of the cold end, the results of the elctrochemical tests indicated that they can be used as a possible replacement for the long exposure tests. The key tests also highlighted the the presence of NH4+ ions in an exhaust gas is benificial to the corrosion resistance od stainless steels in internal cold end application. Its inhibiting effect was more pronounced for the ferritic stainless steels. The project indicated that external corrosion due to salt environments is not the major cause of the failure of cold end components, but rather that internal corrosion due to the condensate is the most detrimental.