In 1972, Okamoto et a/.26 observed strain con­trast around voids in an austenitic stainless steel Fe-18Cr-8Ni-1Si during irradiation in a high — voltage electron microscope. They attributed this con­trast to the segregation strains predicted by Anthony. This is the first reported experimental evidence of RIS. Soon after, a chemical segregation was directly measured by Auger spectroscopy measurements at the surface of a similar alloy irradiated by Ni ions.27

It was then realized that ifthe solute concentration near the point defect sinks reaches the solubility limit, a local precipitation would take place. In 1975, Barbu and Ardell28 observed such a radiation — induced precipitation (RIP) of an ordered Ni3Si phase in an undersaturated Ni-Si alloy.

The analysis of strain contrast and concentration profiles measured by Auger spectroscopy suggested that undersized Ni and Si atoms (which can be more easily accommodated in interstitial sites) were diffus­ing toward point defect sinks, while oversized atoms (such as Cr) were diffusing away. Such a trend, later
confirmed in other austenitic steels and nickel-based alloys,29 led Okamoto and Wiedersich27 to conclude that RIS in austenitic steels was due to the migration of interstitial-solute complexes, and they proposed this new RIS mechanism, in addition to the ones involv­ing vacancies (Figure 2(c)). Then again, Marwick30 explained the same experimental observations by a coupling between fluxes of vacancies and solute atoms, pointing out that thermal diffusion data showed Ni to be a slow diffuser and Cr to be a rapid diffuser in austenitic steels. We will see later that, in spite of many experimental and theoretical studies, the debate on the diffusion mechanisms responsible for RIS in austenitic steels is not over.

Following these debates on RIS mechanisms, it became common to refer to the situation illustrated in Figure 2(a) as segregation by an inverse Kirkendall (IK) effect (the term was coined by Marwick30 in 1977) and to the one in Figure 2(b) as segregation by drag effects, or by migration of vacancy-solute complexes. In the classical Kirkendall effect,31 a gra­dient of chemical species produces a flux of defects. It occurs typically in interdiffusion experiments in A-B diffusion couples, when A and B do not diffuse at the same speed. A vacancy flux must compensate for the difference between the flux of A and B atoms, and this leads to a shift of the initial A/B interface (the Kirkendall plane). The IK effect is due to the same diffusion mechanisms but corresponds to the situa­tion where the gradient of point defects is imposed and generates a flux of solute. The distinction between RIS by IK effect and RIS by migration of defect-solute complexes, initially proposed for the vacancy mechanisms, was soon generalized to inter­stitial fluxes by Okamoto and Rehn.32,33 RIS in dilute alloys, where solute-defect binding energies are clearly defined and often play a key role, is com­monly explained by diffusion of solute-defect complexes, while the IK effect is often more useful to explain RIS in concentrated alloys. This distinction

is reflected in the modeling of RIS (see Section 1.18.3). However, it is clear that RIS can occur in dilute alloys without migration of solute-defect fluxes. Moreover, such a terminology and sharp distinction can be somewhat misleading; the mechanisms are not mutually exclusive. In the case of undersized B atoms, for example, a strong binding between interstitial and B atoms can lead to a rapid diffusion of B by the interstitial (IK effect with DBi > DaO and to the migration of interstitial-solute complexes. More gen­erally, one can always say that RIS results from an IK effect, in the sense that it occurs when a gradient of point defects produces a flux of solute. Nevertheless, because they are widely used, we will refer to these terms at times when they do not create confusion.