| High purity, single-phase materials were employed to simulate transformation characteristics and evaluate the microstructural evolution along the fusion boundary in dissimilar metal welds (DMWs). High purity iron, Type 409 ferritic stainless steel, and 1080 pearlitic steel were selected as the base metal substrates and Monel (70Ni-30Cu) as the filler metal. The resulting weld metal microstructures produced in iron/monel system are analogous to those observed in engineering materials used in cladding and dissimilar metal welding. Weld metal microstructures change from fully martensitic, to an austenite plus martensite, to a fully austenitic microstructure as base metal dilution of the filler metal decreases. Microanalysis of welds produced on Type 409 with Monel filler exhibited a fully austenitic weld microstructure. In both systems, evidence of typical epitaxial growth observed in homogenous weld metal combinations was not evident. However, in 1080 steel, evidence of normal epitaxial growth was observed at the fusion boundary where solidification and HAZ grain boundaries converged.; Misorientation analysis using transmission electron diffraction Orientation Imaging Microscopy analyses revealed various orientation relationships (OR) between HAZ and weld metal grains at the fusion boundary. Welds between iron and Monel exhibited trends toward the Bain, and Kurdjumov-Sachs or Nishyama-Wassermann ORs, the Bain being the strongest tendency. These orientation relationships are a result of a mobile fusion boundary within the austenitic temperature range, and subsequent nucleation and growth of {dollar}alpha{dollar}-ferrite in the HAZ within the alpha ferrite temperature range. The ORs or grain boundary types along the fusion boundary in Type 409/Monel combination welds exhibit a more random distribution, resulting from heterogeneous nucleation at the partially melted HAZ substrate.; A new theory is proposed for the evolution of the type II boundaries so frequently associated with fabrication related failures. It is proposed that the type II boundaries form within the austenite temperature range, where the fusion boundary becomes an {dollar}gamma{dollar}-{dollar}gamma{dollar} boundary which subsequently migrates into the weld metal. Grain boundary migration is driven by composition and thermal gradients, and local strain energies that exist adjacent the fusion boundary in DMWs. Likewise, elimination of grain boundary area, mainly solidification grain boundaries, contributed to the driving force behind this {dollar}gamma{dollar}-{dollar}gamma{dollar} fusion boundary migration. |
