By using the same ‘form factors,’ the moderator brick mean weight loss can be estimated, assuming that weight loss is proportional to burnup or fluence.

The gas temperature will vary roughly linearly in the axial (vertical) direction from the inlet tempera­ture T to the outlet temperature T2. A more detailed profile may be calculated using a thermohydraulics
code. The radial temperature can be assumed to follow the radial flux profile. Thus, an approximate mean gas temperature for an individual moderator brick may be obtained.

In a Li/V blanket, it is believed that the interior ofthe wall needs to be coated with insulator ceramics for mitigating the pressure drop caused by magnetohydro­dynamic effects (see also Chapter 4.21, Ceramic Coatings as Electrical Insulators in Fusion Blan­kets). Corrosion of vanadium alloys in liquid Li might not be a concern if the entire inner wall is covered with the insulating ceramic coating. However, since the idea to cover the insulator ceramic coating again with a thin vanadium or vanadium alloy layer was presented for the purpose of preventing liquid lithium from intruding into the cracks in the ceramics coating, the corrosion of vanadium alloys in liquid lithium again attracted attention. It is known that the corrosion of vanadium alloys in liquid lithium is highly dependent on the alloy composition and lithium chemistry. Espe­cially, the N level influences the corrosion in complex manners.28,29 Figure 13 shows a summary of the weight

 

deformation in Li with that in vacuum.25 However, the correlation of creep data is subject to the alloy heat and manufacturing processes as well as test methods and environments. Figure 12 shows the comparison of the NIFS-HEAT-2 creep strain rate versus creep strain data for tests in vacuum and Li environments at 1073 K, for the same batch of NIFS — HEAT-2 creep tubes.25,26 The figure clearly shows reduced strain rate in Li environments. A possible factor could be N pick-up from Li and the resulting

Ti

Cr

Figure 13 The compositions of V-Ti-Cr alloys (wt%) with increase (area I) and decrease (area II) of mass (g cm-2) after holding of samples in Li at 973 K, 500 h. Reproduced from Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P. J. Nucl. Mater. 2000, 283-287, 1282-1286.

gain and loss in V-xCr-yTi systems in Li.30 High Ti alloys showed a weight increase by forming a TiN layer and high Cr alloys exhibited a weight loss as a result of the dissolution of Cr-N complexes. As the boundary of the two contradictory changes, Ti:Cr^2:1 was observed.

Recently, a corrosion test using monometallic thermal convection Li loop made of V-4Cr-4Ti was conducted at 973 K for 2355 h. Because of the temperature gradient, weight loss and weight gain of V-4Cr-4Ti samples occurred at the hot leg and cold leg, respectively. However, the loss rate corresponded to only < 1 pm year- and the degradation of the mechanical properties were shown to be small.31

V-4Cr-4Ti alloys have been developed mainly for use in Li environments, which are extremely reducing conditions. For the use of vanadium alloys in oxidizing conditions, a different alloy optimization may be necessary. The corrosion of vanadium alloys in oxi­dizing environments is of interest both for the perfor­mance of the pipe exterior out of the breeding blanket and application in non-Li coolant systems such as gas and water systems. Oxidation kinetics of vanadium alloys were studied and showed either parabolic or linear kinetics.32, As the surface oxide layer is not formed or, if formed, not protective to the internal oxidation, alloying with other oxide-formers is neces­sary for improvement. The addition of Si, Al, or Y was shown to significantly suppress the weight gain during exposure to air above 873 K as shown in Figure 14.34

50

40

♦

—1—— ▼—- *—

200 300 400

Hydrogen concentration (wppm)

Figure 15 Total elongation as a function of hydrogen concentration for V-4Cr-4Ti alloys with different O levels. Modified from DiStefano, J. R.; Pint, B. A.; DeVan, J. H.

J. Nucl. Mater. 2000, 283-287, 841-846; Chen, J. M.; Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y. J. Nucl. Mater. 2004, 325, 79-86.

However, the addition of these elements was not effec­tive in suppressing corrosion in water. Increase in Cr level was shown to be effective, instead.

The effects of oxygen level on hydrogen embrittle­ment have been investigated. Figure 15 compares elon­gation versus hydrogen concentration for V-4Cr-4Ti
alloys with various O levels. The loss of ductility by hydrogen charging was shown to be enhanced

35,36

by impurity oxygen.

It is important to note that material modification by radiation arises from two primary spectral-related processes. In addition to the neutron-induced dis­placement of atoms there can be a chemical and/or isotopic alteration of the steel via transmutation. With the exception of helium production, transmuta­tion in general has been ignored as being a significant contributor to property changes ofstainless steels and nickel-base alloys. In this chapter, transmutation is shown to be sometimes much more important than previously assumed.

Both the displacement and transmutation pro­cesses are sensitive to the details of the neutron flux-spectra, and under some conditions each can synergistically and strongly impact the properties of the steel during irradiation. In addition to the brief summary presented below on flux-spectra issues rel­evant to stainless steels, the reader is referred to various papers on transmutation and its consequences in different reactor spectra.5- ’21-23

Transmutation may be subdivided into four cate­gories of transmutants. Three of these are relevant to fission-derived or fusion-derived spectra, and the fourth is associated with spallation-derived spectra. The first three are solid transmutants, gaseous trans­mutants, and ‘isotope shifts,’ the latter involving pro­duction of other isotopes of the same element. While the latter does not change the chemical composition of stainless steels, it is an underappreciated effect that is particularly relevant to nickel-containing alloys such as stainless steels and nickel-base alloys when irradiated in highly thermalized neutron spectra.

Whereas the first three categories arise from dis­crete nuclear reactions to produce discrete isotopes of specific elements, the spallation-induced transmuta­tion arising in accelerator-driven devices involves a continuous distribution ofevery conceivable fragment of the spalled atom, producing every element below that of the target atom across a wide range of isotopes for each element. While individual solid transmutants in spallation spectra are usually produced at levels that do not change the alloy composition significantly, the very wide range of elements produced allows the possibility that deleterious impurities not normally found in the original steel may impact its continued viability. This possibility has not received sufficient attention and should be examined further ifspallation devices continue to be developed.

Another consequence of spallation-relevant trans­mutation is that the induced radioactivity per unit mass is correspondingly much higher than that pro­duced per dpa in other spectra. The majority of the spalled fragments and their daughters/granddaugh- ters are radioactive with relatively short half-lives, leading to materials that are often much more diffi­cult to examine than materials irradiated in fission spectra.

Most importantly, there is a very strong produc­tion of hydrogen and helium in spallation spectra at levels that are one or two orders of magnitude greater than produced in most fission or fusion spectra.5,6,21

While there is a tendency to view displacement and transmutation processes as separate processes, it will be shown later that under some circum­stances the two processes are strongly linked and therefore inseparable in their action to change alloy behavior.

Research into the effects of irradiation on nickel — based alloys peaked during the fast reactor develop­ment programs carried out in the 1970s and 1980s. Interest in these materials focused on their high resis­tance to radiation-induced void swelling compared to austenitic steels, though a perceived susceptibility to irradiation embrittlement limited their application to some extent. Nevertheless, the Nimonic alloy PE16 was successfully used for fuel element cladding and subassembly wrappers in the United Kingdom, and Inconel 706 was utilized for cladding in France. Both of these materials are precipitation hardened and consequently have high creep strength, and much research and development of alternative alloys was directed toward maintaining swelling resistance and creep strength while aiming to alleviate, or at least understand, irradiation embrittlement effects. There has been some revival of interest in nickel-based alloys for nuclear applications, and various aspects of radiation damage in such materials have recently been reviewed by Rowcliffe et al‘ in the context of Gener­ation IV reactors, and by Angeliu et at2 in consider­ation of their use for the pressure vessel of the Prometheus space reactor. Nickel-based alloys are also candidate structural materials for molten salt reactors, for which resistance to corrosion by molten fluoride salts and high-temperature creep strength are prime requirements, though intergranular attack by the fission product tellurium and irradiation embrittlement due to helium production are poten­tially limiting factors for this application.3

This chapter focuses on the void swelling behav­ior, irradiation creep, microstructural stability, and irradiation embrittlement of precipitation-hardened nickel-based alloys. Fundamental to all of these effects are the basic processes of damage production
by the creation of vacancies and interstitial atoms in displacement cascades, and the ways in which these point defects migrate and interact with, causing the redistribution of, solute atoms. Detailed discus­sions of damage processes and radiation-induced segregation are beyond the scope of this chapter but these topics will be introduced where necessary, particularly in relation to void swelling models. More detailed reviews are given in Chapter 1.01, Fundamental Properties of Defects in Metals; Chapter 1.03, Radiation-Induced Effects on Microstructure; Chapter 1.11, Primary Radiation Damage Formation; Chapter 1.12, Atomic-Level Level Dislocation Dynamics in Irradiated Metals and Chapter 1.18, Radiation-Induced Segregation.

Typical compositions of nickel-based alloys and some precipitation-hardened steels, which are con­sidered in this chapter, are shown in Table 1. Alloy compositions are generally given in weight percent throughout this chapter unless stated otherwise. Precipitation-hardened alloys may be utilized in a number of different heat-treated conditions, which are generally abbreviated here as ST (solution trea­ted), STA (solution treated and aged), and OA (over­aged). Further information on the material properties of nickel alloys is given in Chapter 2.08, Nickel Alloys: Properties and Characteristics.

Neutron fluences are generally given for E > 0.1 MeV unless indicated otherwise. Atomic dis­placement doses (dpa) are generally given in NRT (Fe) units, although the half-Nelson (N/2) model was
widely used particularly in the United Kingdom in the 1970s6. The exact relationship between these units will vary depending on the neutron spectrum (which may differ, not only from one reactor to another, but also depending on location within a reactor), but approximate conversion factors for fast reactor core irradiations are

1026n m~2(E > 0.1MeV)

= 5dpa NRT(Fe) = 6.25dpa (N/2)

Tantalum and its alloys have historically been exam­ined for high-temperature nuclear applications, par­ticularly in the various space reactor programs. For reasons similar to those of Nb and its alloys, various alloying combinations of Ta were examined, particu­larly in the late 1950s to 1960s. Much of this effort emphasized the development of solid solution (W and Re additions) and dispersion-strengthened (Hf addi­tion) alloys. While Ta-alloys pay a penalty in higher density over, for example, Nb, and decreases the low temperature density-compensated strength to com­parable values on Nb-base alloys. The higher melting temperature of Ta (3290 k) results in better strength retention above 1000 K and in density-compensated

creep strength.12,41

Early work on substitutional solid solution — strengthened Ta-10W for aerospace applications42 led to limited examination of this alloy for irradiation environments. The improved strengthening by addi­tion of a maximum of 10 wt% allows the retention of suitable nonirradiated ductility and weldability.43,44

However, the use of Ta-10W in space reactor appli­cations where liquid alkali coolants are considered was unacceptable because of the lack of oxide getter — ing elements such as Hf that form stable dispersion — strengthened structures. The T-111 (Ta-8%W-2% Hf) alloy, with its demonstrated compatibility with liquid alkali metals and improved strength over pure Ta while retaining ductility and weldability, has been a lead candidate alloy in space reactor systems since the 1960s.4 Though a considerable effort has been made on the Ta-10W and T-111 alloys, the irradiation properties database is very small. Irra­diated mechanical property behavior follows typical bcc alloys in which radiation hardening effects including limit ductility appear and are expected at temperatures ^0.3 Tm (987 K).3

4.08.7.1 Sodium Compatibility

It is essential to evaluate the environmental effects of sodium on the mechanical strength properties of ODS steels to ensure their structural integrity throughout their design life-time in SFR. ODS-steels basically display superior compatibility with sodium. For 9Cr — ODS steel (M93) and 12Cr-ODS steel (F95), which are potential cladding materials for SFR, their UTS at 700 °C after exposure to sodium in a stagnant state is shown in Figure 30.53 Both show almost constant strength after exposure to sodium, and it was confirmed that there is no degradation up to 10 000 h. For conven­tional ferritic steel without Y2O3, a clear strength reduction occurs above 600 °C due to decarburization phenomena in sodium. ODS steel does not show such a clear strength reduction because the fine Y2O3 oxide particles remain stable in steel, thereby maintaining the strength of the steel.

Figure 31 shows the results of creep-rupture tests with internally pressurized specimens in a stagnant sodium environment.54 The creep-rupture strength of 9Cr-ODS steel (M11) in sodium is equal to its strength in air, and no impact from a sodium environ­ment was observed. However, under a flowing sodium condition of 4.5 ms, the element nickel penetrates the surface of ODS steel cladding, where an increase in nickel concentration and decrease in chromium con­centration were observed at 700 °C. These results suggest that the effects of a sodium environment can be ignored under stagnant conditions; however, as fuel cladding is utilized in an environment with a high flow rate of sodium, the effects of the microstruc­ture change associated with nickel diffusion into the cladding surface need to be considered.53

Based upon the evidence from UK and US creep experiments, Davies and Bradford49,68 suggest the following:

• The strain induced change in CTE is not a func­tion of secondary creep strain, but saturates after a dose of-30 x 1020 n cm 2 EDN (-3.9 dpa).

• There is evidence, from both thermal and irradia­tion annealing, for a recoverable strain several

 

times that of primary creep, and a lower associated secondary creep coefficient that has been previ­ously assumed. • The dose at which the recoverable strain saturates bears a striking similarity to that of the saturation of the CTE change. Davies and Bradford49,68 proposed a new creep model (the M2 model) without the term reflecting changes in CTE due to creep, but containing one additional term, recoverable creep:

 

Figure 19 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 600 °C under a compressive stress of 13.8 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46-54.

 

♦ Experimental creep strain “ ■ True creep strain — — — CTE correction strain —— Predicted apparent creep strain

 

Figure 20 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 600 °C under a compressive stress of 20.7 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46-54.

where ec is the total creep strain; s, the applied stress; l, X, and b, are the empirical fitting parameters; ki and k2, the primary and recoverable dose constants respectively; and W is the oxidation change factor (with respect to Young’s modulus) and is analogous to the structure factor. The terms in the eqn [26] are proportional to esu and the effects of structural changes and radiolytic oxidation (gasification of graphite by an activated species that occurs in CO2 cooled reactors) are also included. The rates of satu­ration of the primary and recoverable creep compo­nents are controlled by the dose constants ki and k2. The first and last terms in eqn [26] are primary and secondary creep as in the prior UK creep model, with the middle term being recoverable creep.

Primary creep is still fast acting, but in the AGR temperature range of 400-650 °C, appears to act on a longer fluence scale equivalent to that associated in the United Kingdom with the Young’s modulus pinning,69 ki = 0.1, and saturates at 1 esu (a = 1). The irrecoverable creep is synonymous with second­ary creep, but with a coefficient, b, derived from the irrecoverable strain postthermal anneal, as

0. 15 per 1020 n cm 2 EDN (-1.3 dpa) in the AGR temperature range. The lateral creep strain ratios for primary and recoverable creep are assumed to be the Poisson’s ratio and secondary creep is assumed to occur at constant volume.

Figure 24 shows the performance of the M2 models applied to some high dose ATR-2E tensile creep data52 when irradiated at 500 °C in high flux reactor (HFR), Petten. The prediction matches the observed data well up to significant fluence of -160 x i020ncm~2 EDN (-21 dpa). Only beyond
this fluence does the new model prediction deviate from the data with a delay in the increase in creep strain at high doses that is often referred to as the ‘tertiary’ creep phase.

Figure 25 shows the corresponding compressive creep data,52 irradiated at 550 °C. The model over predicts the data slightly but follows the trend remark­ably well up to a significant fluence of -i60 x i020 n cm~2 EDN (-21 dpa). Beyond this fluence, the compressive prediction also indicates a ‘tertiary’ creep, but the data does not extend into this region. The data52 also indicates a possible difference between tensile and compressive creep (seen more clearly in Figure 18).

Saturation of CTE with creep strain as reported by Davies and Bradford49,68 is not however in agree­ment with other published data. Gray70 reported CTE behavior with creep strain (up to 3%) for three different graphites at irradiation temperatures of 550 and 800 °C. Saturation of the CTE in the

manner described by Davies and Bradford49,68 for UK AGR graphite was not observed.

Thermal conductivity in nuclear graphite is usually determined by measuring thermal diffusivity using the laser flash method at ^30 °C. The mechanism for thermal conductivity in graphite over the tem­peratures of interest in nuclear reactors is lattice vibration (phonon) conductance. There is a pro­nounced reduction in thermal conductivity with increased temperature attributed to phonon — phonon scattering. At low irradiation fluence, there

9
6
5
CD О
4
О
3 1 0-1———— 0 10

Weight loss (%)

Figure 42 Coefficient of thermal expansion of thermally oxidized graphite. Modified from Hacker, P. J.; Neighbour, G. B.; McEnaney, B. J. Phys. D Appl. Phys. 2000, 33, 991-998.

is a significant decrease in thermal conductivity due to fast neutron irradiation, attributed to an increase in scattering in the damaged lattice. This decrease in thermal conductivity (or increase in thermal resis­tivity) saturates in the medium fluence range. At very high fluence, there is a secondary decrease attributed to microcracking due to high crystallite strains. There is also evidence of change in temper­ature dependence with irradiation.

The thermal conductivity in crystallite basal plane is much larger than that perpendicular to basal plain; thus, Ka ^ Kc. The thermal resistivity can be described by the equation below:

1 _ 1 1 1

Kg_ KB+KU+KD

• U — Umklapp scattering (German for turnover/ down) or phonon-phonon scattering (due to increase in temperature)

• D — scattering due to defects (caused by irradiation)

• B — boundary scattering (structural effects)

Changes to these resistances will be reflected in the thermal conductivity of polycrystalline graphite. Ther­mal conductivity is significantly decreased by radio­lytic oxidation. Data is usually presented as the reciprocal of conductivity, that is, thermal resistivity.

4.11.16.1 Pile Grade A

Although PGA is significantly anisotropic over the range of interest to the Magnox reactors, the change in thermal resistivity can, for practical purposes, be considered as invariant to grain direction. Changes
in thermal resistivity in PGA graphite are given in Figure 43. There is a significant change in the rate of increases in thermal resistivity between 250 and 300 °C, giving a similar trend to the change in crystal growth rates between these two temperatures. The increase in thermal resistivity is significant (a factor of 100) at 150 °C, for a relatively low fluence. The low irradiation temperature data for PGA in Figure 43 do not reach a high enough fluence to saturate. However, at the higher temperatures, data is near saturation.

In general, the performance of NPP safety-related concrete structures has been very good. However, there have been several isolated incidences that, if not remedied, could challenge the capacity ofthe contain­ment and other safety-related structures to meet future functional and performance requirements. Table 2 presents a summary of local degradation mechanisms that have been observed by one organization during condition surveys ofvarious concrete structures at both United States and foreign NPPs located in areas having several different climatic conditions.25 Some general observations derived from these results were that vir­tually all NPPs have experienced cracking of the con­crete structures that exceeds typical acceptance criteria for width and length; numerous NPPs had groundwater intrusion occurring through the power block or other subsurface structures; and aging con­cerns exist for subsurface concrete structures, as their physical condition cannot be easily verified. Collec­tively, it was concluded in this study that the general performance of the NPP concrete structures, has been quite favorable and proper evaluation and treatment of observed degradation at an early stage is both a cost — effective and necessary approach to long-term plant operations.

Initially, degradation of NPP concrete structures in the United States occurred early in their life and has been corrected.26-28 Causes were primarily related either to improper material selection and construction/design deficiencies or environmental

effects. Examples of some of the problems attributed to these deficiencies include low 28-day concrete com­pressive strengths; voids under the post-tensioning tendon-bearing plates resulting from improper con­crete placement; cracking of post-tensioning tendon anchor heads due to stress corrosion or embrittlement; and containment dome delaminations due to low- quality aggregate materials and absence of radial steel reinforcement or unbalanced prestressing forces.29-31 Other construction-related problems included occur­rence of excessive voids or honeycomb in the concrete, contaminated concrete, cold joints, cadweld (steel reinforcement connector) deficiencies, materials out of specification, higher than code allowable concrete temperatures, misplaced steel reinforcement, post­tensioning system button-head deficiencies, and water-contaminated corrosion inhibitors.26 Although continuing the service of a NPP past the initial operating license period is not expected to be limited by the concrete structures, several incidences of age — related degradation have been reported.28-33 Examples of some of these problems include corrosion of steel reinforcement in water intake structures, corrosion of post-tensioning tendon wires, leaching of tendon gal­lery concrete, low prestressing forces, and leakage of corrosion inhibitors from tendon sheaths. Other related problems include cracking and spalling of con­tainment dome concrete due to freeze-thaw damage,
low strengths of tendon wires, contamination of corro­sion inhibitors by chlorides, and corrosion of concrete containment liners. As the plants age, the incidences of degradation are expected to increase, primarily due to environmental effects. A listing of documented concrete problem areas by plant, type reactor, and degradation is available.34 Documented information on problem areas experienced with NPP concrete struc­tures in other countries has also been assembled.35 Figure 4 presents examples of occurrences of degrada­tion that have been observed at NPPs. Anchor head failure and containment dome delamination shown in the figure represent occurrences related to materials selection and design, respectively, with the remainder representing aging-related occurrences.

To the first-order most researchers concentrate on the cold-work level as the primary way to delay void swelling, although it is known that increasing cold work beyond a certain level specific to each alloy yields diminishing returns, with the optimum level usually chosen to be 20-25% for austenitic alloys. Larger levels are often counter-productive in that the additional stored energy at higher cold-work levels sometimes induces recrystallization during irradia­tion, often resulting in higher swelling.1

Additionally, in some alloys and metals it is diffi­cult to nucleate voids under some combinations of temperature and dpa rate due to the difficulty to establish a stable dislocation network. Cold working in some cases can actually shorten the transient by providing a stable glissile dislocation network and thereby accelerate swelling, as observed in model Fe-Cr-Ni alloys and simple metals such as nickel

132-134

and iron.

The starting thermal-mechanical condition ofthe alloy plays an important role in determining the transient duration via its influence on the starting dislocation density, but more importantly in deter­mining the distribution or chemical activity of the active elements. For instance, aging of an alloy at intermediate temperatures that encourage carbide precipitation, for instance, is the most effective way to produce the shortest transient and the highest swelling.1

There are many other examples. For instance, the chemical activity of an element like phosphorus is very sensitive to the inter-pass annealing tempera­ture range employed in producing cold-worked tub­ing by multiple drawings. It is speculated that

phosphorus can be either in solution or existing as small invisible precipitates oflesser chemical activity depending on the inter-pass annealing temperature or tube feed rate through the furnace.1,135

As carbon plays a role in both carbide and inter­metallic phase evolution, and its chemical activity can be strongly affected by thermal and mechanical his­tory, it exerts a strong and often complex effect on the transient duration. One aspect of this complexity is the often-observed two-peak swelling behavior ver­sus temperature that strongly varies with thermal — mechanical treatment.1 This effect is so strong that the swelling valley between the two peaks often occurs at the peak flux position. Cold-working tends to suppress the low temperature peak more than the high temperature peak due to its effect to delay and homogenize carbide formation. Removing almost all carbon into precipitates by aging erases the double peak behavior and usually produces the largest amount of swelling, as shown in Figure 52.