Abstract Scope |
Introduction: Face-centered cubic (FCC) alloys such as nickel and austenitic stainless steels are important to many industries, most notably nuclear power generation, turbine manufacturing, and petrochemical. These alloys are prone to ductility-dip cracking (DDC), a high-temperature, solid-state cracking phenomenon. Susceptible metals experience an abnormal intermediate temperature ductility loss which leads to cracking upon applying sufficient restraint, such as the restraint introduced by large, multipass welds. Cracking is typically found in underlying weld metal which has been reheated numerous times. A unified mechanism for DDC has been elusive. Various literature sources do not agree on whether sulfur causes DDC or if it simply deteriorates a metal’s resistance to DDC. To learn more, several plates of high-purity alloy 690 were manufactured at varying sulfur levels. One plate was kept at high purity as a control. Two levels of sulfur were desired: 0.003<wt% S<0.015 and wt% S>0.015. Sample composition was evaluated via combustion analysis to precisely determine the sulfur levels after the plates were made. Each plate was machined into sets of smaller specimens for thermal cycling designed to “prime” the microstructure for the formation of DDC. Microstructural evolution was tracked via SEM imaging and EBSD analysis. Experimental Procedures: These sulfur studies were conducted with alloy 690 owing to its relative simplicity, high purity, and known susceptibility to DDC. Plates were manufactured in a casting process where ingot melts of high purity alloy 690 received controlled additions of sulfur, or none in the case of the control ingot. These ingots were then hot-rolled and pressed to be used in the laboratory experiments. Combustion analysis showed an initial composition of <0.001wt%S in the control plate (L), 0.01 in medium plate 1 (M1), 0.02 in medium plate 2 (M2), and 0.6 in the high plate (H). The samples machined from each plate were welded so the analyzed microstructures would be similar to production conditions where DDC is found. To begin the priming trials, a sample from each plate was prepared for SEM and EBSD analysis to record the beginning microstructure in each plate. Then, a set of four samples was selected from each plate. These four samples would undergo one of the following numbers of thermal cycles: 1, 2, 5, 10. The thermal cycles are based on recorded weld metal heat affected zone thermal histories in similar production welds in high-chromium nickel-based alloys, and any melting is avoided to eliminate solidification confounding with other microstructural changes. Next, samples would be prepared for the SEM as before and analyzed to observe how this thermal cycling evolved the microstructure and if sulfur’s presence affected this evolution. The SEM analysis consisted of imaging and a complete EBSD analysis was conducted to examine grain size, boundary type, sigma distribution, etc. It was determined EDS would be unnecessary as the interaction volumes are too high to observe sulfur diffusing preferentially to grain boundaries. Results and Discussion: The overarching hypothesis is if DDC is typically found in only underlying weld metal and not weld metal towards the surface, then the reheating this weld metal experiences must be priming it for the formation of DDC on subsequent weld passes. Thus, it is predicted the microstructural changes which take place during thermal cycling will be forming the conditions in which DDC is typically observed. That is, long, straight, and migrated grain boundaries in higher-energy regions such as boundaries which have high coincidence site lattice (CSL) misorientations. Since sulfur is a low-solubility impurity and an embrittling element, it is expected to affect this microstructure evolution in some way. Past research indicates it has a proclivity to diffuse to grain boundaries and disrupt nano-scale bond lengths by reacting with nickel atoms to form nickel sulfide molecules. This introduces atomic lattice strains which over a myriad of repeated occurrences
adds up to cause micro-scale embrittlement. This may appear in the EBSD analysis as a higher concentration of high-energy boundary morphologies, higher boundary strains, and changes in boundary migration and/or thermal faceting. Such faceting has been observed in past research involving DDC fracture in sulfur-containing nickel-based weld metal 52M. It has never been implicated as a causal or contributing factor toward the formation or propagation of DDC in the past, but as a local grain boundary phenomenon which may be affected by sulfur’s presence, it will be examined in this study. Conclusion: The presence of sulfur at elevated levels, even those allowable by the alloy 690 specification, affects microstructural evolution in ways which may increase susceptibility to DDC. This study aims to provide a more concise explanation for why sulfur does this. The explanation will come in terms of grain size, grain boundary morphology, grain boundary type distribution, and grain boundary strains. There has often been disagreement about what role sulfur plays here, and hopefully this will help to clarify how sulfur behaves in this scenario. Past research has shown the formation of DDC in both sulfur-containing metals and those with high purity. Perhaps it is the same microstructural evolution happening in both cases, but with sulfur causing accelerated development towards a microstructure which is more susceptible to DDC. |