## 3 Results and Discussion
distribution of the Ni-Cu-Sn and Ag-Sn intermetallic precipitates can be rationalised through the different thermodynamic behaviour of the alloy systems: in the temperature range near the respective eutectic temperature, liquid Sn has a much lower solubility for Cu and especially for Ni with respect to formation of the binary or ternary intermetallic Ni-Cu-Sn phases than for Ag with respect to formation of Ag\({}_{3}\)Sn [21; 25; 26]. Hence, formation of the Ni-Cu-Sn intermetallics predominantly occurs close to the Cu- and or Ni-source (i.e., the Ni-metallisation layer or the Ag-Cu-In layer) by scallop-type growth of grains into the liquid phase. By contrast, Ag can easily dissolve into the liquid Sn phase, and Ag\({}_{3}\)Sn then only precipitates from the liquid phase at considerably lower temperatures during the cooling stage, thus forming Ag\({}_{3}\)Sn-islands within the Sn solder layer. For Si, with a pronouncedly higher thermal conductivity than Al\({}_{2}\)O\({}_{3}\) (but lower heat capacity, Table 1), a similar microstructure is found (Figure 4c). A marked difference is the less pronounced formation of Ni-Cu-Sn scallops along the interface between Ni-metallisation and solder layer, and an even more localised formation of Ni-Cu-Sn intermetallics along the boundary between solder layer and reactive foil, due to the shorter time at less elevated temperatures. Finally, in case of Cu, a distinctly different joint microstructure can be observed (Figure 4d): first, along the interface between the Ni-metallisation and the Sn-solder layer, no formation of Ni-Cu-Sn intermetallic scallops takes place, although comparison with the initial microstructure of the galvanically deposited Sn-layer showed that the Sn-layer was indeed completely melted during reactive joining, i.e., temperatures of at least 232 \({}^{\circ}\)C were reached at the solder-substrate interface. Scallops are also absent in the transition region between the solder layer and the reactive foil layer. In addition, formation of the Ag\({}_{3}\)Sn-phase is limited to a narrow region close to the original Ag-Cu-In-layer. Obviously, only very limited intermixing of Ag, Cu and In with the Sn-solder could take place due to the short time at comparatively low temperatures, leading to very localised precipitation of the intermetallic particles. In addition, a substantial amount of sub-micron-sized pores can be observed within the Sn-layer, in contrast to the other substrate types. This indicates that, in accordance with the results of CT (Figure 2), only very limited coarsening of the porosity could take place in case of Cu. Hence, the microstructure developing for each joining material is strongly influenced by the thermal properties and the ensuing specific time-temperature-profile upon reactive joining. The mechanical strength of the reactive solder joints was assessed by shear tests. Although a firm bond was obtained for all four substrate types, the corresponding shear strength of the joints shows a strikingly different behaviour (Figure 5, also see [28]): for borosilicate glass, an apparently very low shear strength of less than 5 MPa was obtained. However, analysis of the specimens after shear testing showed that failure occurred via cracking of the substrate plates, preferentially along the cracks already formed during reactive joining and not by failure of the solder joint. Hence, this value does not represent the actual shear strength of the joint. After optimisation of the joining procedure by considerably increasing the amount of solder in order to avoid the formation of cracks, nominal shear-strength values of about 20 MPa could be obtained (see [10]), also in this case limited by the failure of the brittle glass. For Si, the maximum shear strength was limited by brittle failure of the substrates, as well.