No description available.

No description available.

No description available.

No description available.

No description available.

No description available.

Note:

Due to the relatively low tensile strength of concrete, cracks are inevitable in reinforced concrete structures. Therefore, studying the cracking behaviour of reinforced concrete elements and controlling the width of cracks are necessary objectives both in research and in design. The introduction of higher strength reinforcing steel has exacerbated the problem of crack control. Using higher strength steel, means less steel is required for a given structure to satisfy the strength requirements. The stiffness after cracking is reduced and wider crack widths will occur under normal service loads. Unserviceable cracking may encourage corrosion in the reinforcement and surface deterioration, and may lead to long term problems with durability. Indeed excessive cracking results in a huge annual cost to the construction industry because it is the most common cause of damage in concrete structures. In this study cracking caused by both shrinkage and external loads in reinforced concrete members is examined experimentally and analytically. The mechanisms associated with cracking and the factors affecting the time-varying width and spacing of both direct tension cracks due to restrained shrinkage deformation and flexural cracks due to the combined effects of constant sustained service loads and shrinkage are examined. Laboratory tests on eight fully restrained slab specimens were conducted for up to 150 days to measure the effects of drying shrinkage on the time-dependent development of direct tension cracks due to restrained deformation. The effect of varying the quantity, diameter, and spacing of reinforcing steel bars was studied. In addition, an analytical model previously developed without experimental verification by Gilbert (1992) to study shrinkage cracking was modified and recalibrated. A second series of tests on twenty four prismatic, singly reinforced concrete beams and slabs subjected to monotonically increasing loads or to constant sustained service loads for up to 400 days, were also conducted. The effects of steel area, steel stress, bar diameter, bar spacing, concrete cover and shrinkage were measured and quantified. An analytical model is presented to simulate instantaneous and time-dependent flexural cracking. The tension chord model (Marti et al, 1998) is modified and used in the proposed model to simulate the tension zone of a flexural member and the time-dependent effects of creep and shrinkage are included. The analytical predictions of crack width and crack spacing are in reasonably good agreement with the experimental observations.

Most stress corrosion cracking (SCC) mechanisms initiate from localised corrosion (pitting, intergranular attack, de-alloying), which provides local stress concentration. Alloys are generally more susceptible to SCC than pure metals because selective dissolution or oxidation is possible. De-alloying involves the selective dissolution of the less noble (LN) component from an alloy. The more noble (MN) component enriches on the surface forming a brittle, metallic, nanoporous layer. In noble metal alloys and brass, SCC shows correlation with the threshold LN content below which de-alloying stops (the parting limit). In Fe-Ni-Cr engineering alloys de-alloying may be responsible for Cl-SCC, although this has not been proven explicitly. Initial indications show that de-alloying causes SCC in hot, caustic environments. In some cases, Ni enrichment and porosity are associated with cracks in stainless steel after long-term service in nuclear high temperature water environments, but it is unclear if this plays a causal role in cracking. Here the de-alloying mechanism (primarily the effect of Ni (MN) content) and its relationship to SCC in Fe-Ni-Cr materials (Fe10Ni, 310SS and Alloy 800) is examined using a hot caustic environment, and compared to classical de-alloying systems. De-alloyed layers formed on all materials, although Alloy 800 required a higher temperature. Increasing Ni content improved de-alloying resistance according to classical theory. Unlike classical systems, de-alloying occurred with concurrent MN dissolution and, at open circuit potential (OCP), the layers retained significant Fe and Cr (LN) instead of being ‘almost pure’ MN. Layers formed with applied anodic potential were friable and highly LN depleted. This behaviour was successfully modelled in Kinetic Monte Carlo simulations. Recently, it has been shown that SCC in noble element alloys depends on the mechanical integrity (quality) of the de-alloyed layer; a finding that was reflected here. At 140 °C at OCP the layer on 310SS was too thin to promote SCC and Alloy 800 did not de-alloy significantly. Layers formed with anodic potential did not result in SCC. In 50% NaOH at 280 °C, severely stressed 310SS cracked where thick de-alloyed layers formed. However, the thin layer formed on Alloy 800 was associated with SCC, even with low residual stress.

Most stress corrosion cracking (SCC) mechanisms initiate from localised corrosion (pitting, intergranular attack, de-alloying), which provides local stress concentration. Alloys are generally more susceptible to SCC than pure metals because selective dissolution or oxidation is possible. De-alloying involves the selective dissolution of the less noble (LN) component from an alloy. The more noble (MN) component enriches on the surface forming a brittle, metallic, nanoporous layer. In noble metal alloys and brass, SCC shows correlation with the threshold LN content below which de-alloying stops (the parting limit). In Fe-Ni-Cr engineering alloys de-alloying may be responsible for Cl-SCC, although this has not been proven explicitly. Initial indications show that de-alloying causes SCC in hot, caustic environments. In some cases, Ni enrichment and porosity are associated with cracks in stainless steel after long-term service in nuclear high temperature water environments, but it is unclear if this plays a causal role in cracking. Here the de-alloying mechanism (primarily the effect of Ni (MN) content) and its relationship to SCC in Fe-Ni-Cr materials (Fe10Ni, 310SS and Alloy 800) is examined using a hot caustic environment, and compared to classical de-alloying systems. De-alloyed layers formed on all materials, although Alloy 800 required a higher temperature. Increasing Ni content improved de-alloying resistance according to classical theory. Unlike classical systems, de-alloying occurred with concurrent MN dissolution and, at open circuit potential (OCP), the layers retained significant Fe and Cr (LN) instead of being ‘almost pure’ MN. Layers formed with applied anodic potential were friable and highly LN depleted. This behaviour was successfully modelled in Kinetic Monte Carlo simulations. Recently, it has been shown that SCC in noble element alloys depends on the mechanical integrity (quality) of the de-alloyed layer; a finding that was reflected here. At 140 °C at OCP the layer on 310SS was too thin to promote SCC and Alloy 800 did not de-alloy significantly. Layers formed with anodic potential did not result in SCC. In 50% NaOH at 280 °C, severely stressed 310SS cracked where thick de-alloyed layers formed. However, the thin layer formed on Alloy 800 was associated with SCC, even with low residual stress.