Y. Dai

Paul Scherrer Institut, Villegen PSI, Switzerland

G. R. Odette and T. Yamamoto

University of California, Santa Barbara, CA, USA © Introduction and Overview

This chapter reviews the profound effects of He on the bulk microstructures and mechanical properties of alloys used in nuclear fission and fusion energy systems. Helium is produced in these service envir­onments by transmutation reactions in amounts rang­ing from less than one to thousands of atomic parts per million (appm), depending on the neutron spec­trum, fluence, and alloy composition. Even higher amounts of H are produced by corresponding n, p reactions. In the case of direct transmutations, the amount of He and H are simply given by the content weighted sum of the total neutron spectrum averaged energy dependent n, a and n, p cross-sections for all the alloy isotopes ((sn, a)) times the total fluence (ft). The spectral averaged cross-sections for a specified neutron spectrum can be obtained from nuclear data­base compilations such as SPECTER,1 LAHET,2 and MCNPX3 codes. He and H are also produced in copious amounts by very high-energy protons and neutrons in spallation targets of accelerator-based nuclear systems (hereafter referred to as spallation proton-neutron (SPN) irradiations, SPNI).4, The D-T fusion first wall spectrum includes 14MeV neutrons («20%), along with a lower energy spec­trum («80%). The 14MeV neutron energy is far above the threshold for n, a («5 MeV) and n, p («1 MeV) reactions in Fe.6 Note that some impor­tant transmutations also take place by multistep nuclear reactions. For example, thermal neutrons (nth) generate large amounts of He in Ni-bearing alloys by a 58Ni(nth, g)59Ni(nth, a) reaction sequence. These various irradiation environments also produce a range of solid transmutation products.

High-energy neutrons also produce radiation- induced displacement damage in the form of vacancy and self-interstitial atom (SIA) defects. Vacancies and SIA are the result of a neutron reaction and scattering — induced spectrum of energetic primary recoiling atoms with energies ranging from less than 1 keV, in neutron irradiations, up to several MeV in SPN irra­diations.7 The high-energy primary recoils create cascades of secondary displacements of atoms from their crystal lattice positions, measured in a calcu­lated dose unit of displacements per atom (dpa). As in the case of n, a transmutations, dpa production can also be evaluated using spectral averaged dis­placement cross-sections8 that are calculated using the codes and nuclear database compilations cited above.

Typical operating conditions of various fission, fusion, and spallation facilities are summarized in Table 1. Notably, He (and H) generation in fast fission (He/dpa << 1), fusion (He/dpa « 10), and spallation proton-neutron (He/dpa up to 100) environments differs greatly and this is likely to have significant effects on the corresponding microstructural and mechanical property evolutions.

The primary characteristic of He, which makes it significant to a wide range of irradiation damage phenomena, is that it is essentially insoluble in solids. Hence, in the temperature range where it is mobile, He diffuses in the matrix and precipitates to initially form bubbles, typically at various microstructural trapping sites. The bubbles can serve as nucleation sites of growing voids in the matrix and creep cavities on grain boundaries (GBs), driven by displacement damage and stress, respectively. While He effects are primarily manifested as variations in the cavities, all microstructural processes taking place under irradia­tion are intrinsically coupled; hence, difference in the He generation rate can also affect precipitate, dislo­cation loop, and network dislocation evolutions as well (see Section 1.06.3).

Figure 1, adopted from Molvik et al.,9 schemati­cally illustrates the effects of high He as a function of lifetime-temperature limits in a fusion first wall structure for various irradiation-induced degrada­tion phenomena. At high temperatures, lifetimes (green curve) are primarily dictated by chemical compatibility, fatigue, thermal creep, creep rupture, and creep-fatigue limits. In this regime, He can fur­ther degrade the tensile ductility and the other high — temperature properties, primarily by enhancing grain boundary cavitation, in some cases severely. In aus­tenitic stainless steels (AuSS), high-temperature He embrittlement (HTHE) has been observed at concentrations as low as 1 appm.10,11 In contrast, 9Cr ferritic-martensitic steels (FMS), which are cur­rently the prime candidate alloy for fusion structures, are much more resistant to HTHE.12,13

Figure 1 Illustration of the materials design window for the fusion energy environment, as a function of temperature. Reproduced from Molvik, A.; Ivanov, A.; Kulcinski, G. L.; et al. Fusion Sci. Technol. 2010, 57, 369-394.

At intermediate temperatures (blue curve), growing voids form on He bubbles, and He accumulation largely controls the incubation time prior to the onset of rapid swelling (see Section 1.06.3). FMS are also much more resistant to swelling than stan­dard austenitic alloys,14,15 although the microstruc­tures of the latter can be tailored to be more resistant to void formation by He management schemes.16 High He concentrations can also extend irradiation hardening and fast fracture embrittlement to inter­mediate temperatures.17

At lower temperatures (red curve), where irradiation hardening and loss of tensile uniform ductility are severe, high He concentrations enhance large positive shifts in the ductile-to-brittle transition temperature (DBTT) in bcc (body-centered cubic) alloys.18-20 This low-temperature fast fracture embrittlement phenomenon is believed to be primarily the result of

He-induced grain boundary weakening, manifested by a very brittle intergranular (IG) fracture path, inter­acting synergistically with irradiation hardening. , High concentrations also increase the irradiation hardening at dpa levels that would experience satu­ration in the absence of significant amounts of He.17 A significant concern for fusion is that the dpa- temperature window may narrow, or even close, for a practical fusion reactor operating regime.

What is sketched above is only a very broad-brush, qualitative description of some of the important He effects. The quantitative effects of He, displacement damage, temperature and stress, and their interactions, which control the actual positions of the schematic curves shown in Figure 1, depend on the combination of all the irradiation variables, as well as details of the alloy type, composition, and starting microstructure (material variables). The effects of a large number of interacting variables, the complex interactions of a plethora of physical mechanisms, and the implica­tions to the wide range of properties of concern are not well understood; and even if they were, such
complexity would beg easy description. Therefore, a first priority is to develop a good understanding of and models for the transport and fate of He at the point when it is effectively immobilized in bubbles and voids, often at various microstructural sites. Such insight provides a basis for developing microstruc­tures that can manage He and thus mitigate its dele­terious effects. To this end we next briefly outline key radiation damage processes, including the role ofHe.

Figure 2 schematically illustrates the combined effects of He and displacement damage on irradiation- induced microstructural evolutions.2 Figure 2(a) shows a molecular dynamics simulation of primary displacement damage produced in displacement cas­cades. Most of the initially displaced atoms return to a lattice site (self-heal). Residual cascade defects include single and small clusters of vacancies and SIA. In the temperature range of interest, vacancies (red cir­cles) and SIA (green dumbbells) are mobile. SIA clusters, in the form of dislocation loops, are also believed to be mobile in some cases, undergoing one-dimensional diffusion on their glide prisms.

However, the cascade loops may also be trapped by interactions with solutes. Small cascade vacancy clus­ters may coarsen in the cascade region by Ostwald ripening and diffusion coalescence mechanisms. Both isolated and clustered defects interact with alloy solutes forming cascade complexes. The cascade vacancy clusters dissolve over a time associated with cascade aging, which depends strongly on tempera­ture. The concentration of cascade vacancy clusters, which act as sinks (or recombination centers) for migrating vacancies and SIA, scales directly with the damage rate. Thus, the overall defect production microstructures can be viewed as being composed of steady-state concentrations of diffusing defects, small loops, and cascade vacancy clusters; the latter are important if the irradiation time is much less than the cluster annealing time. Vacancy-SIA recombina­tion at clusters, in the matrix and at vacancy trapping sites, can give rise to important damage rate, or flux, effects.

Figure 2(b) shows that SIA can recombine with diffusing and trapped vacancies, in this case one trapped on a precipitate interface. Figure 2(b) also shows that both bubbles (blue part circle) and voids (orange part circle) often form on precipitates. Figure 2(c) shows that dislocation loops (green hexagon) nucleate and grow due to preferential absorption of SIA (bias). Preferential accumulation of SIA also takes place at network dislocation segments (inverted green T), causing climb. Loop growth and dislocation climb can lead to creation (loops and Herring-Nabarro sources) and annihilation (of oppositely signed network segments) of disloca­tions, ultimately leading to quasi-steady-state densi­ties, as is observed in the case of AuSS.

Figure 2(d) shows that He precipitates to form bubbles (larger blue circles) at various sites, in this case in the matrix. Small bubbles are stable since they absorb and emit vacancies in net numbers that exactly equal the number of SIA that they absorb; thus bubbles grow only by the addition of diffusing He atoms (small blue circle). However, Figure 2(e) shows that when bubbles reach a critical size they convert to unstably growing underpressurized voids (large orange circle containing blue He atoms) due to an excess flux of vacancies over SIA arising from the dislocation bias for the latter defect. Figure 2(e) shows the corresponding growing creep cavities transformed from critical He bubbles on stressed GBs. Designs of microstructures that mitigate, or even fully suppress, these various coupled evolutions are described in Section 1.06.6 and discussed in references.2 ,

Therefore, a master overarching framework for measuring, modeling, and managing He effects must be based on developing and understanding the domi­nant mechanisms controlling its generation, trans­port, fate, and consequences, as mediated by the irradiation conditions and the detailed alloy micro­structure. Figure 3 illustrates such a framework for He generation, transport, and fate. In this framework, experiments and models can be integrated to estab­lish how He is transported to various microstructural trapping (-detrapping) features and how He locally clusters to form bubbles at these sites, as well as in the matrix. The master models must incorporate para­meters that describe He diffusion coefficients under irradiation, binding energies for trapping at the vari­ous sites and He-vacancy cluster and other interac­tion energies.

Given the length and comprehensive character of this chapter, it is useful to provide the reader a guide to what follows. Notably, we have tried to develop useful semi-standalone sections.

Section 1.06.2 describes the various experimental approaches to studying He effects in structural alloys including both neutrons and various types of charged-particle irradiations (CPI).

Section 1.06.3 reviews the historical knowledge base on He effects, which has been developed over the past 40 years, with emphasis on bubble evolution, void swelling, and HTHE processes. While less of current interest, the examples included here pri­marily pertain to standard AuSS, discussions of experiment and modeling are closely integrated to emphasize the insight that can be derived from such coupling. Particular attention is paid to the critical bubble mechanism for the formation of growing voids and grain boundary cavities and the corresponding consequences to swelling and creep rupture. The implications of the coupled models and experimental observations to designing irradiation-tolerant alloys that can manage He are discussed in some detail.

Section 1.06.4 focuses on a much more recent body of observations on He effects in SPNI. The emphasis here is on descriptions of defect and cavity microstruc­tures in both FMS and AuSS irradiated at low to intermediate temperatures and the corresponding effects on their strength, ductility, and fast fracture resistance. Similarities and differences between the SPNI effects and those observed for fission irradia­tions are drawn where possible.

Section 1.06.5 summarizes some key examples of atomistic modeling of He behavior, which has been the focus of most recent modeling efforts. Insight into mechanisms and critical parameters provided by these models will form the underpinning of the com­prehensive master models of He transport, fate, and consequences.

Section 1.06.6 builds on the discussion in Section 1.06.3 regarding managing He by trapping it in a population of small stable bubbles. A specific example comparing FMS to a new class of high — temperature, irradiation-tolerant nanostructured fer­ritic alloys (NFA) irradiated in a High-Flux Isotope Reactor (HFIR) at 500 °C to 9 dpa and 380 appm He is described. The results of this study offer proof in principle of the enormous potential for developing irradiation-tolerant NFA that could turn He from a liability to an asset. Section 1.06.6 again couples these experimental observations with a master multi­scale model of the transport and fate of He in both

FMS and NFA. The predictions of the master model, that is both microstructurally informed and parame­terized by atomistic submodels, are favorably com­pared to the HFIR data.

Section 1.06.7 briefly summarizes the status of understanding of He effects in structural alloys and concludes with some outstanding issues. Reading this summary first may be helpful to general readers who then can access the more detailed information at their own discretion.