## 3 Results and Discussion
### Microstructure Examination of K465 Alloy after Interrupted Creep Test  
After interrupted creep, samples for microstructure examination were cut longitudinally from the gauge section of crept specimens. The SEM micrographs shown in Figure 3 demonstrate that the evolution of \(\gamma^{\prime}\) precipitates in the DC with different creep strains is quite pronounced. Some of the \(\gamma^{\prime}\) particles lost their cuboidal shape, and the coarsening of \(\gamma^{\prime}\) precipitate was obvious even in the sample with a relatively low creep strain of 0.03. The size of \(\gamma^{\prime}\) in this sample increased to 1 \(\upmu\)m. With the increase of creep strain to 0.05, the morphology of all \(\gamma^{\prime}\) precipitates had become irregular, and they had coalesced with each other. Namely, the rafted \(\gamma^{\prime}\) microstructure was completely formed at this creep stage. In the sample with accumulated strain of 0.07, broadening of matrix channels was also apparent apart from the rafted microstructure. The analogous evolution of \(\gamma^{\prime}\) precipitates also occurred in the ID region. One typical microstructure in the interdendritic region after creep is shown in Figure 4. It can be seen that similar microstructural evolution also took place in the interdendritic region during creep, even though the original precipitation of \(\gamma^{\prime}\) was different in these two regions.  
Figure 3: \(Cont.\)  
Figure 2: Typical creep time vs. strain curve (**a**) and strain vs. strain rate curve (**b**) of K465 alloy, dashed lines showing the interrupted test with different creep strain.  
TEM micrographs in Figure 5 show the dislocation configuration of different samples. It can be seen that a large amount of dislocations were produced and the motion of these dislocations was confined by the \(\gamma/\gamma^{\prime}\) interfaces. As a consequence, dense dislocation networks were generated in the matrix channels even in the sample with creep strain of 0.03. This observation also demonstrated that creep deformation is mainly concentrated in the \(\gamma\) matrix, which is a softer phase compared to \(\gamma^{\prime}\) phase. The formation of these networks will change the \(\gamma/\gamma^{\prime}\) interfaces and retard the motion of dislocations in the matrix channels. This phenomenon was commonly observed in all of the samples, and it is the fundamental mechanism for creep deformation of Ni based superalloys [28,29]. In terms of dislocation networks, no apparent difference can be detected with respect to the creep strain. However, the creep deformation is not a significant effect.  
Figure 4: SEM micrograph showing the morphology of \(\gamma^{\prime}\) phase in interendritic region of the sample with creep strain of (**a**) 0.03, (**b**) 0.05, and (**c**) 0.07.  
Figure 3: SEM micrographs showing the morphology of \(\gamma^{\prime}\) phase in the DC with creep strain of (**a**) 0.03, (**b**) 0.05, and (**c**) 0.07.  
with the accumulation of creep strain, a larger number of dislocations were found to penetrate into \(\gamma^{\prime}\) precipitates, as indicated by the arrows in Figure 5c, which displays that these \(\gamma^{\prime}\) particles were sheared by dislocations in order to accommodate higher creep strain.