Os moinhos de bolas são equipamentos presentes em plantas de mineração, sendo uma máquina importante no circuito de cominuição de minérios. Em função de inúmeros fatores como, por exemplo, projeto, fabricação, sobrecarga no equipamento, falta de manutenção e procedimentos de operação inadequados, são desenvolvidas descontinuidades nos componentes estruturais desse equipamento. Os componentes estruturais dos moinhos, basicamente, corpo, tampas e munhões, além do custo elevado, possuem prazos de fabricação que podem variar de dois a três anos dependendo da demanda do mercado. Portanto é cada vez mais necessário que as descontinuidades detectadas nesses componentes sejam corretamente avaliadas. Neste trabalho analisaram-se pela mecânica da fratura a criticalidade de descontinuidades, tais como trincas constatadas em um munhão de um moinho de bolas e compararam-se os valores teóricos de taxa de propagação (crescimento) dessas descontinuidades com valores reais obtidos por meio de inspeções periódicas realizadas neste componente. A nucleação das trincas foi causada por falta de lubrificação nos mancais do moinho, gerando esforços térmicos circunferenciais no munhão, em que a temperatura estimada do contato munhão e bucha atingiu a faixa de 100 °C a 150 °C. No período analisado, os resultados obtidos por meio da norma BS7910 mostraram-se mais próximos dos valores reais do que a norma ASME Seção XI, Apêndice A. O tamanho da trinca mais crítica foi aceito pelos critérios da norma BS7910 e aprovado apenas na condição de emergência da norma ASME Seção XI, Apêndice A. / The grinding ball mills are equipments present in mining plants, being important in the ore comminution circuits. Depending on numerous factors, such as, for example, design, manufacturing, overloads, poor maintenance and inadequate operating procedures, flaws are developed in the structural components of this equipment. The structural components of a mill, basically, shell, heads and trunnions, besides high costs, have lead times that might vary from two to three years, according to market demand. Therefore, it becomes increasingly necessary that any flaws in those components to be properly evaluated. This paper analyzed the fracture mechanics of the criticality of flaws, such as cracks observed in a ball mill trunnion and compared the theoretical values of growth rate of these defects with actual values obtained through periodic inspections performed in this component. The cracks nucleation was caused by lack of lubrication in the trunnion bearings, generating circumferential thermal stresses, thus the estimated temperature of the trunnion and bushing contact achieved the range 100 °C to 150 °C. During the analyzed period, the results obtained by the standard BS7910 proved to be closest to the actual values than the standard ASME, Section XI, Appendix A. The most critical crack size was accepted by the criteria of the BS7910 and approved only on emergency condition of ASME Section XI, Appendix A.

Steam treatments reduce porosity(1) of sintered iron and sintered iron-base alloy products and provide them with increased compressive strength, hardness, resistance to wear and corrosion. The aims of the present investigation were: (a) to examine the influence of copper content and process variables on the response of sintered iron to steam treatment. (b) to explore the possibility of combining steam treatment with precipitation hardening in sintered iron/copper alloys. (c) to survey the influence of process parameters on the mechanical and wear properties of these materials. Samples prepared from - 100# A.S.C. iron powder and - 100# pre-alloyed powder containing 2,4,6 and 8% copper, pressed to densities of 6.0, 6.4, 6.8 g/cc and sintered for 1 hour at 1120 oC were treated in superheated steam and inert gas at 525 oC for periods ranging from 5 to 100 minutes and under vacuun for 500 hours. The main results of the investigation were: (i) Addition of copper in amounts of 2 to 8% slightly improves the resistance of iron to steam oxidation. (ii) During ageing in steam, the hardness increased due to simultaneous precipitation of copper from solution and formation of and void filling by layers of iron oxide. The results indicated that maximum benefit is for Fe - 2% Cu alloy and steam ageing of higher copper content alloys reduces the effect of precipitation of copper. (iii) Steam treatment of sintered. pure iron improvesthe wear properties but the effect of this treatment on Fe-Cu. alloys is not pronounced and sometimes disadvantageous. (iv) Steam oxidation of sintered pure iron improved radial-crushing stress (R.C.S.) but this treatnent reduced the R.C.S. of iron-copper alloys.

The aim of this study is to determine the influence alloy diffusion processes have on corrosion.

No description available.

For polymer materials to compete with metals in engineering applications, they must possess adequate mechanical properties. Some crystalline plastics are able to offer enhanced properties but these tend to be very expensive. Often a viable alternative is to modify commodity polymers through further processing techniques which introduce molecular orientation, thereby producing high modulus polymers.

An investigation into the effect of indentation on the fatigue behaviour of an aluminium alloy (A1 - 2.5wt% Cu - 1.2wt% Mg) has been carried out. Room-temperature aged (RTA) and high-temperature aged (HTA) specimens were cycled in tension-compression under conditions of constant plastic strain in the range epsilonp = 1.0 x 10 -4 to 1.4x10-3 in a 20 kN Mand servohydrolic machine. The fatigue behaviour was cyclic hardening in the RTA state and cyclic softening in the HTA state. During fatigue, the RTA specimen surfaces were quickly covered with coarse slip lines, which after prolonged cycling developed into PSBs. Unlike the slip lines, the PSBs did not always traverse the entire grain, but started off in patches which steadily grew larger and more intense. Dislocation structures were studied using TEM and were found to develop from a structure consisting of quenched-in loops and helices into slip bands concentrated on groups of {111} planes, often dominated by dislocations of a single Burgers' vector. Further cycling resulted in more concentrated slip in localized bands leading to a misorientation of these bands relative to the matrix. The development of surface relief in HTA specimens was much slower than in the case of the RTA specimens and took place over thousands of cycles. The slip lines were shorter and broader, and in late stages more wavy in appearance, both features typical of cross-slip. The dislocation structure consisted of some loop formation but mostly dislocations tangled with the S precipitates and were densely packed at the precipitate-matrix boundary. There was no evidence for the break-up of S precipitates during cyclic deformation. Surface damage was introduced in the form of Vickers micro-hardness indentations. The indentations were made in the middle of the grains in order to obtain a single-crystal response. Optical microscope studies of the surface were carried out using Nomarski contrast and scanning laser microscopy, and the dislocation configurations were studied by TEM. Analytical models were used to see whether a correlation could be found between the observations of slip and calculations of resolved shear stresses. A reasonable correlation was found, which prompted the extension of the modelling to fatigue slip near indentations. The development of fatigue damage was studied both before and after various stages of fatigue, and it was found that the presence of indentations had a significant effect on the fatigue behaviour of the indented grain. During the initial fatigue cycles, slip lines were observed near indentations in the RTA specimens, and modelling indicated that the effect could be due to the residual stresses surrounding the indentation. In all cases the slip appeared in trace directions of planes containing slip systems with high Schmidt factors, but in many cases on planes not containing the active slip system in fatigue. In the HTA specimens a similar effect was noted at the stage when surface slip started to appear, but only associated with slip on the slip plane containing the active slip system. Cracking was initiated near most indentations, but none were found to lead to the ultimate failure of the specimens. This could be mainly due to stress concentrations at the specimen shoulders. Cracks initiated along the ridge of the indentations oriented approximately perpendicular to the stress axis and appeared at the corresponding corners of the indentations approximately half-way into the fatigue life. In the case of the RTA specimens, the cracks propagated along traces of {111} slip planes, whereas in the HTA specimens the cracks were more irregular and sometimes followed the traces of {100} planes. Modelling suggested that the preferred path for crack propagation was associated with a concentration of shear stress on highly stressed slip planes. The cracks were never observed to stop propagating, and some cracks propagated into neighbouring grains, and it is thought that the condition for damage to lead to failure is the size of the indentation in combination with the magnitude of the applied stress/strain.

The indentation behaviour and erosion properties of two commercially available ceramic materials, a silicon carbide (Hexoloy SA) and silicon carbide-titanium diboride composite (Hexoloy ST) have been investigated over a range of experimental conditions. The microstructure of both materials has been examined using reflected light and scanning electron microscopy. Hexoloy SA is a single phase material with a grain size typically ranging from 4 to 8 mum, while Hexoloy ST is a two phase particulate composite, containing about 16 vol% of discrete titanium diboride particles in a silicon carbide matrix, with a more uniform grain size. The materials have both been shown to have weak grain boundaries. The silicon carbide-titanium diboride interface is weak and this is believed to be due to tensile residual stresses arising from the mismatch in coefficients of thermal expansion of the two phases. Vickers indentation testing indicated that both materials have similar hardness values, but that the composite was significantly tougher than the monolithic material. Sub-surface crack profiles have been examined with a particular regard to radial and lateral cracking. It was found that the scale of lateral cracking was not directly proportional to the length of radial cracks in these materials. Indeed, lateral cracks were not seen when the radial/median system was fully formed, but only when it was partially formed. This is an important observation since one of the fundamental assumptions of two models of erosion is that radial and lateral length are directly proportional. Another important finding of the indentation study was that lateral cracking occurred to a greater extent in the composite than in the monolithic materials at low loads, indicating that wear of the composite may be relatively more extensive for the smaller erodent sizes. Erosion testing has been performed using a gas blast apparatus. Different sizes of silica and silicon carbide erodent have been used for tests from room temperature to 1000°C. With the silica erodent, material loss progressed by small scale cracking. The mechanisms of material removal involved grain boundary cracking in the monolithic material and grain boundary cracking and cracking along the particuiate-matrix interface in the composite. For the silicon carbide erodent, lateral cracking has been shown to be the dominant mechanism of material removal. In the monolithic SIC the lateral cracking scales with erodent size, while in the composite the TiB2 particles inhibit growth of the laterals generated by the largest erodent, but proved to be detrimental when using the smallest erodent. This observation was consistent with the observations from quasi-static Indentation. The presence of an easily removed oxide on the surface of the TiB2 particles has led to an increase in the erosion rate of the composite at temperatures greater than 800 °C for the silica erodent. At lower temperatures both materials behaved similarly. When using the silicon carbide erodent, increasing the temperature resulted in an increase in the erosion rate for both materials although at the lower temperatures, the composite was more erosion resistant than the monolithic material. As the temperature increased, the erosion rates converged, suggesting that the toughening mechanisms of the composite were decreasing in effectiveness. Thus, it has been shown that the presence of TiB2 particles can lead to increased or decreased erosion resistance relative to the monolithic material, depending on the precise erosion conditions. In general, the composite has the lower wear rate at lower temperatures and larger erodent sizes. Also, it has been shown that cracking due to quasi-static indentation using a sharp indenter is consistent with the damage produced by hard, sharp erodent particles at room temperature.

The corrosion of SiCp / 6061 A1 in 3% NaCl solution, at pH values of 3.8, 7.2 and 10.2, has been studied using SEM and combined scanning Auger electron/ X-ray microscopy. The MMC, fabricated by powder metallurgy, contains approximately 15 percent (volume) SiC as < 10 mum particles. Scatter diagrams are applied to analyze elemental maps created by scanning Auger microscopy. The experiments reveal that some impurity elements such as Fe, Ti etc., introduced during fabrication, form intermetallic compounds and that these are the dominant factors causing micro-galvanic corrosion. The Auger maps show, by inspection, that oxygen is enriched after exposure around these intermetallics. This has been confirmed by using scatter diagrams to analyze the data contained within the map. The phenomenon is probably caused by the dissolution of aluminium, resulting in the deposition of porous corrosion products such as Al(OH)3 and AlOOH. Because of a low overpotential for oxygen reduction, the intermetallic compounds are expected to act as cathodes. This is confirmed by the finding of magnesium hydroxide on the surface of the intermetallics after exposure of the specimen in MgCl2 solution. It is shown that the onset of this deposition can be used to estimate the cathodic current density at the cathode. Current density at the SiC particles, estimated by this method is negligible, suggesting that the principal cause of corrosion is found in the interaction between the aluminium alloy and the intermetallics. The same method is used for SiC, / 6061 A1 MMC. Intermetallic particles are found much more likely to be located at a fibre / matrix interface, and they are still a sensitive factor for localized corrosion. There is some Cr in the Fe-containing intermetallics in this matrix, this type of intermetallic is not very susceptible to micro-galvanic corrosion. Another different aspect is that the reinforcing material in SiCf / 6061 A1 MMC is directly associated with the micro-corrosion cell. The reinforcing SiC fibre belongs to the SCS series which consists of three different layers. The SiC is grown on a carbon core by chemical vapour deposition and a protective coating is applied on the outside of the fibre. This coating is mainly composed of carbon which is slightly enriched in silicon at the surface. In this investigation we found that carbon core and carbon coating act as cathodes and form a galvanic cell with surrounding aluminium alloy. This is confirmed by the magnesium decoration method. Electrochemical techniques are also used. Several kinds of SiC fibres and graphite fibres are cathodically polarized in 3% NaCl solution. By comparison with the anodic polarization curve of A1 alloy, the galvanic cells between A1 alloy and carbon core or protective coating are confirmed. There is no galvanic effect between A1 and pure SiC. Al3Fe is extracted from intermetallic enriched A1 alloy, and the electrochemical performance of this intermetallic is investigated.

Matrix damage and its effects on mechanical properties have been examined for SiC (Nicalon1) fibre reinforced glass and glass ceramic matrix composites under quasi-static and fatigue loading conditions. Nicalon/Pyrex laminates of different lay-ups have been tested under quasi-static tension. The elastic moduli have been measured and matrix damage monitored as a function of applied strain. The mechanical properties are strongly influenced by the presence of crystalline regions in the matrix which promote microcracking. Laminated plate theory is used to provide bounds to the moduli of the laminates. For unidirectional and simple crossply Nicalon/CAS2 laminates the quasi-static stress/strain behaviour and associated matrix damage accumulation have been examined in detail. The damage development with applied stress was quantified by counts of crack density (in both longitudinal and transverse plies), stiffness loss and cumulative residual strain. The quasi static stress/strain behaviour during continuous tests (accumulating damage) and discontinuous tests (constant damage) have been modelled using a stress analysis based on Aveston, Cooper and Kelly (ACK) theory. The continuous stress/strain behaviour of (0/90) crossply laminates has been modelled using a shear-lag analysis developed previously to describe the transverse ply cracking behaviour of polymer matrix composites. The analysis is modified to account for longitudinal ply cracking. Matrix damage development in unidirectional and (0/90) crossply laminates under quasistatic cycling and high frequency fatigue loading have been studied. For unidirectional laminates stable stress/strain hysteresis loops were obtained during quasi-static cycling, corresponding to stable matrix damage states. These and similar loops obtained after high frequency fatigue are modelled using, the discontinuous stress/strain analysis. It is suggested that the effect of high frequency fatigue is to decrease the interfacial shear strength.

The evaluation of the dynamic fracture properties with the current impact testing techniques is subject of some controversy, because the crack initiation conditions cannot be established with accuracy. Besides, some of them are limited to low loading velocities and the evaluations are made following a static approach. In consequence, a fundamental study is required with two aims, namely, a development of a new impact testing technique and the analysis of the extension of the fracture mechanics concepts to crack initiation under dynamic loading. These are the objectives of this research project. The impact testing technique proposed is based on the shock tube technique and is capable to test the specimens, which are central cracked circular plates, under static and dynamic conditions. The findings show that the specimens are loaded uniformly with blast waves, which have an initial ramp that increases steadily without oscillations. The magnitude of the load is controllable and reproducible and the fracture analysis should be based on a dynamic calibration. This technique has been applied to the evaluation of the dynamic fracture toughness of PMMA following an experimental-numerical procedure. The crack initiation load and the strain rate were measured experimentally and the fracture toughness was calculated numerically with standard solutions and with numerical calculations of the J-Integral with the finite element code ABAQUS. The problem of crack initiation under pressure load has not been completely solved. Consequently, the significance of the static and dynamic test results was established with the findings of additional experimental and numerical work, which showed that the modulus of elasticity and fracture stress increases with the strain rate, that there is plastic deformation at the crack tip of those specimens tested under static conditions and that the dynamic calculations with ABAQUS involves the interaction of the stress waves with the boundaries. Furthermore, it has to be noticed that the crack tip is under bending stresses. Thus, this test is not a standard plane strain material dynamic test. Finally, conclusions are drawn on the use of the this experimental-numerical technique and its further application.