The n = 3 regime, though in principle corresponding to the power-law con­trolled (n = 4-1) creep mechanism, differs from it at a mechanistic level. The power law controlled creep mechanism (as will be discussed in the fol­lowing section) is mostly dislocation climb-controlled commonly noted in pure metals and class-II or metal-class alloys. In contrast the n = 3 regime is dislocation glide-controlled creep usually exhibited by alloys known as

class-I or class-A, and hence the n = 3 regime at times is referred to as alloy type creep behavior. The creep behavior of materials can thus be classified into two groups: class-A and class-M. The creep behavior of solid solutions or class-A alloys at intermediate stresses and for specific material param­eters consists of 3 regimes. As shown in Fig. 3.10 , as stress increases the
stress exponent changes from 5 (region I) to 3 (region II) and back to 5 (region III).

The creep behavior illustrated in Fig. 3.10 is a consequence of a compe­tition between two rate controlling mechanisms: dislocation climb and dis­location glide. Once the dislocations are generated from Frank-Read (FR) sources on parallel glide planes, the leading edge dislocations first glide and then climb to annihilation. In pure metals dislocation glide is relatively faster compared to the diffusion-controlled climb and thus climb becomes the rate controlling process resulting in n = 5. In class-A alloys, the rate of glide is controlled by the diffusion of the solute atoms, thereby leading to a relatively slower rate of glide compared to that of climb whereby the vis­cous glide of dislocations becomes the rate controlling process with n = 3; this mechanism is known as Weertman microcreep.54 Region II, the three power-law creep regime, is also known as the viscous glide regime. Viscous glide is described by

. 0.35 n a

єs = Ds I

s A s I E

where A is an interaction parameter that depends upon the viscous process controlling dislocation glide and Ds is the solute diffusivity.

The viscous process can be of different types. According to Cottrell and Jaswon,58 the dragging force could be due to the segregation of solute atmo­spheres to moving dislocations. The dislocation speed in this case is con­trolled by the rate of migration of the solute atoms. Fisher55 suggested that the viscous process had its origin in the destruction of the short range order in solid solution alloys. The disorder created by dislocation motion would result in the formation of a new interface thereby the interfacial energy becomes the rate controlling process. Suzuki59 suggested that the drag­ging force was an outcome of solute atoms segregating to stacking faults. There are suggestions that the obstacle to dislocation motion could be the stress-induced local ordering of solute atoms. The ordering of the region surrounding a dislocation reduces the total energy of the crystal pinning the dislocation.

The three power-law creep region has usually been observed to occur in solid solutions with a large atom size mismatch. Alloys with higher con­centrations of the solute atoms seem to prefer the three power-law creep regime as a viable creep mechanism. In fact, for very high concentrations of the solute atoms, regime II could be suppressed. In addition, class-A alloys usually exhibit either no or little primary creep or a region characterized by an increasing slope (increasing strain rate). This is in sharp contrast to pure metals and class-M alloys that exhibit a distinct primary creep curve with a decreasing strain rate; distinguishing features of class-A and class-M alloys

2.11 Deformation microstructure in Nb-modified Zr-alloy crept in the three power-law regime.61

were summarized by Murty.60 Some of the different alloys that exhibit three power-law creep behavior are Al-Zn, Al-Ag, and Ni-Fe alloys.

Microstructural features

Since the n = 3 region is dislocation glide-controlled, recovery based pro­cesses (such as climb) are considered to be less important. The deformation microstructures are found to consist of a large number of dislocations as shown in Fig. 3.11. In comparison to class-M alloys, the deformation micro­structures of class-A alloys are devoid of subgrains.61