Unlike most thermoplastics, polycarbonate can undergo large plastic deformations without cracking or breaking. As a result, it can be processed and formed at room temperature using sheet metal techniques, such as bending on a brake. Even for sharp angle bends with a tight radius, heating may not be necessary. This makes it valuable in prototyping applications where transparent or electrically non-conductive parts are needed, which cannot be made from sheet metal. PMMA/Acrylic, which is similar in appearance to polycarbonate, is brittle and cannot be bent at room temperature.
In the automotive industry, injection-molded polycarbonate can produce very smooth surfaces that make it well-suited for sputter deposition or evaporation deposition of aluminium without the need for a base-coat. Decorative bezels and optical reflectors are commonly made of polycarbonate.Its low weight and high impact resistance have made polycarbonate the dominant material for automotive headlamp lenses. However, automotive headlamps require outer surface coatings because of its low scratch resistance and susceptibility to ultraviolet degradation (yellowing). The use of polycarbonate in automotive applications is limited to low stress applications. Stress from fasteners, plastic welding and molding render polycarbonate susceptible to stress corrosion cracking when it comes in contact with certain accelerants such as salt water and plastisol. It can be laminated to make bullet-proof "glass", although "bullet-resistant" is more accurate for the thinner windows, such as are used in bullet-resistant windows in automobiles. The thicker barriers of transparent plastic used in teller's windows and barriers in banks are also polycarbonate.
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An analysis of the literature on bisphenol A leachate low-dose effects by vom Saal and Hughes published in August 2005 seems to have found a suggestive correlation between the source of funding and the conclusion drawn. Industry-funded studies tend to find no significant effects whereas government-funded studies tend to find significant effects.[31]
Stress corrosion cracking (SCC) of UNS S31603 austenitic stainless steel (ASS), UNS S32205 duplex stainless steel (DSS) and UNS S32750 super duplex stainless steel (SDSS) was investigated. SCC tests were carried out at 110 ºC for 500 h under drops of synthetic seawater (DET, drop evaporation test). Two loading conditions were investigated: 50 % and 100 % of the experimental yield strength of each steel. DSS and SDSS specimens showed no susceptibility to SCC under loading of 50 % of their yield strength, contrary to ASS, but all steels fractured at the highest load. SCC nucleated under the salt deposit formed on the surface of all specimens. SCC propagation was mainly transgranular, but SCC propagation of DSS also featured crack ramification in the austenite phase. In addition, SDSS also presented crack propagation along the ferrite/austenite interfaces. Transgranular cleavage fracture was also observed in all fractured specimens, but DSS also presented ferrite/austenite interfacial brittle fracture, while SDSS also featured intergranular brittle fracture.
In Brazil, ninety-five percent of the petroleum is extracted from offshore wells, requiring a large range of engineering components and materials, which should present high corrosion resistance to marine environments. In this sense, ASS, DSS and SDSS are usually selected for the manufacturing of offshore pressure vessels1, due to their high corrosion resistance and good mechanical properties. These offshore components are usually submitted to service conditions combining tensile stress and chloride environments, which may induce their premature failure by stress corrosion cracking (SCC), even under tensile values below the yield strength2-4. A comparison of the general properties and characteristics of the ASS and DSS and SDSS is shown in Table 1. Although DSS and SDSS are more expensive than ASS, their values of fracture toughness, yield strength and pitting resistance equivalent number (PREN) are comparatively much higher5. Despite their wide and growing usage of DSS and SDSS components, they might be subjected to various types of mechanical and environmentally induced failures during their life cycle, especially when these components are exposed to temperatures in the range of 300 ºC to 900 ºC, which might promote the precipitation of stable and deleterious phases5-8.
The stable crack propagation of ASS, DSS and SDSS usually presents a ductile behavior under non aggressive environments, but the action of SCC promotes a brittle behaviour in these steels, leading to the formation of intergranular or cleavage-like transgranular brittle fractures15. It is interesting to observe that the exposure to a corrosive environment might decrease the fracture toughness of ASS from KIC equals to 72 MPa.m(1/2) to KISCC (acidic solution) equals to 53 MPa.m(1/2)5-8. There are at least four atomic-level mechanisms to explain the stable brittle crack propagation of metals and alloys caused by SCC4,5-10,15-22:
The tensile decohesion of the atomic bounds of the metallic material ahead of the crack tip (cleavage) is promoted by the adsorption of ions of the corrosive solution on the metallic surfaces near the crack tip. This adsorption is followed by surface diffusion of these ions into the region ahead of the crack tip, causing local embrittlement due to the decohesion of the atomic bounds (see Figure 1)4,5-7,23-24;
The creation of vacancies on the metallic surfaces near the crack tip is caused by the removal of the elements of the crystal lattice by the corrosive solution. These vacancies preferentially diffuse towards the region ahead of the crack tip due to the presence of a stress gradient. The metallic ion surface mobility is promoted by the presence of contaminants in the corrosive solution. The stable step-like brittle propagation of the crack takes place when these vacancies reach the crack tip (see Figure 2)5-7,25-26;
The formation of a brittle surface film on the metallic surfaces near the crack tip locally reduces the fracture toughness, promoting the stable step-like brittle propagation of the crack. According to this model, the formation of the brittle film is induced by the environment, which controls the kinetics of the stable crack propagation. In this sense, the next step of the stable brittle propagation of the crack will proceed when another layer of brittle film is formed ahead of the crack tip (see Figure 3)4,5-7,27-28;
The shear strain in the region located ahead of the crack tip is promoted by the adsorption of atomic hydrogen on the surfaces of the crack (see Figure 4). These hydrogen atoms diffuse into the FCC lattice, preferentially into the region located ahead of the crack-tip. The presence of interstitial hydrogen atoms in this region locally increase the plasticity of the metal. In the case of FCC iron, the hydrogen atom (atomic radius equals to 0.48 Ã…) preferentially occupies its octahedral interstitial sites (size of approximately 0.52 Ã…) without introducing any lattice elastic strain. The presence of interstitial hydrogen in this lattice decreases its shear modulus and the repulsion forces between edge dislocations, consequently reducing the distance between these dislocations and increasing both the density of the dislocation pile-up and the stress field ahead of the leading dislocation. This stress field will promote the nucleation of a secondary crack at the dislocation slip obstacle (such as a grain-boundary) and this crack will propagate towards the crack tip by a brittle mechanism due to the high dislocation density between the obstacle and the tip. Additionally, this same stress field will activate the slip systems of the adjacent grains, repeating the process of nucleation and brittle propagation of secondary cracks (see Figure 4)5-7,15-20.
Figure 4 SCC mechanism in FCC metals: hydrogen induced plasticity, promoting denser dislocation pile-up and higher stress field, causing the nucleation and brittle propagation of secondary cracks from the obstacle towards the tip of the primary crack. The stress field will also activate the slip system of the adjacent grain, repeating the process of nucleation and brittle propagation of secondary cracks4.
Although, there are several mechanisms to explain the brittle crack propagation during SCC, none of them allows a straightforward approach for the microstructural design of SCC resistant SSs. Therefore, the experimental determination of the critical stress for the activation of the SCC during service conditions is an important technique to investigate the SCC susceptibility of metals and alloys. SCC tests usually use smooth or pre-cracked specimens under static or dynamic mechanical load exposed to a particular environment (corrosive solution and temperature). In some cases, a pre-loaded specimen is immersed in a standard solution with controlled temperature, while in other SCC tests a standard solution drips on pre-loaded specimens with controlled temperature5-7,29-32. The drop evaporation test (DET), for instance, simulates the external environment of offshore components by dripping synthetic seawater on pre-loaded specimens with controlled temperature. This method is performed using different combinations of tensile stress, temperature and standard solutions to determine the critical parameters for the onset of SCC30. Seawater DET at temperatures above 105 ºC forms a saline deposit on the surface of the specimen, which act as a barrier to the diffusion of oxygen, locally promoting the crevice corrosion and SCC nucleation and growth underneath the deposit31-32.
DET was carried at 110 ºC for 500 h and the dripping rate of synthetic sea water was kept equal to (10 1) drops per minute. After DET40, the microstructures of the DETed SSs were analyzed using the same etching procedures as described previously. Optical microscopy using an Olympus microscope model BX51M and a Leica stereomicroscope model M205 C were used during the metallographic and macrofractographic examination. Microfractographic examination of the cracked/ fractured specimens used a scanning electron microscope (SEM FEI Quanta 400) with acceleration voltage of 20 kV and current of 1 nA. FEI Quanta 3D FEG dual beam microscope was also used as focused ion beam (FIB) microscopy in order to characterize the microstructure under the surface of SCC (cross-sectional examination). 2ff7e9595c