X-ray diffraction (XRD) and transmission electron microscopy (TEM) have been used to characterize microstructural and microchemical changes produced by neutron irradiation in zirconium and zirconium alloys. Zircaloy-2, Zircaloy-4, and Zr-2.5Nb alloys with differing metallurgical states have been analyzed after irradiation for neutron fluences up to 25 × 1025 n.m-2 (E > 1 MeV) for a range of temperatures between 330 and 580 K. Irradiation modifies the dislocation structure through nucleation and growth of dislocation loops and, for cold-worked materials in particular, climb of existing network dislocations. In general, the a-type dislocation structure tends to saturate at low fluences (< 1 × 1025 n.m-2). The c-component dislocation structure, however, may evolve over long periods of irradiation (for fluences >10 × 1025 n.m-2 in some cases). The phase structure is also modified by irradiation. The common alloying/impurity elements, Fe, Cr, and Ni, are relatively insoluble in the α-phase but are dispersed into the α-phase during irradiation irrespective of the state of the phase initially containing these elements, i.e., metastable β-phase or stable intermetallic precipitate. The stable intermetallic particles may undergo structural changes dependent on their composition and the temperature. For the metastable dual-phase α/β-alloys (Zr-2.5Nb alloy), the β-phase structure is modified during irradiation, but the change is complex, being a combination of thermal decomposition and radiation-induced mixing.

X-ray diffraction (XRD) line-broadening analysis has been used to determine dislocation densities in zirconium alloys with hexagonal closepacked (hep) crystal structures and a complex distribution of dislocations reflecting the plastic, anisotropy of the material. The validity of the technique has been assessed by comparison with direct measurements of dislocation densities in deformed polycrystalline and neutron-irradiated single crystal material using transmission electron microscopy (TEM). The results show that-there is good agreement between the XRD and TEM for measurements on the deformed material whereas there is a large discrepancy for measurements on the irradiated single crystal; the XRD measurements significantly underestimating the TEM observations.

Intergranular residual stresses can exist in zirconium alloys, especially when there is a large distribution of grain orientations. The stresses result from the anisotropic plasticity and thermal expansion of the hexagonal close-packed crystal structure of α-zirconium. Apart from complicating the characterisation of materials using lattice parameter measurements, the intergranular stresses can significantly affect material behaviour, especially in nuclear reactor environments, and there is therefore a great deal of interest in their measurement.The effects of specimen preparation and surface relaxation on X-ray diffraction measurements of lattice parameters of zirconium alloys have been investigated by comparing bulk neutron diffraction with X-ray diffraction on identical materials. The results show that: (i) intergranular or interphase residual stresses exist in dual-phase Zr-2.5Nb pressure tubes; (ii) the stresses normal to the surface of an X-ray diffraction specimen are not relieved completely when there are intergranular residual stresses in the material. One can conclude that intergranular stresses have to be considered when determining chemical compositions from lattice parameter measurements and also when measuring macroscopic residual stress using X-ray diffraction.

The rearrangement of dislocations in a deforming metal during a stress reversal is thought to contribute to a lowering of the yield stress in the reverse direction. Part of the stress-strain curve of an annealed polycrystalline OFHC Cu specimen, deformed first to -118 MPa in compression and then in tension, is shown in fig.1. The workhardening rate, θp = dσ/dεp where σ is the true stress and εp the true plastic strain, is plotted against |σ| in fig.2. The compressive workhardening rate, for stresses beyond -118 MPa (point A) where the reversal was made, was obtained from a second specimen that was deformed entirely in compression. Figures 1 and 2 reveal the following: (a) the reverse proportional limit (σr) is ∽40 MPa lower than the “forward” peak stress (σf), and it requires ∽1% εp before the reverse flow stress achieves a value equal to the previous foreward maximum; (b) θp-reverse is initially higher than θp-forward, then falls below, and eventually becomes equal to θp-forward at a much higher stress than the peak forward stress.