Comprehensive nuclear materials 4 02 radiation damage in austenitic steels

Comprehensive nuclear materials 4 02 radiation damage in austenitic steels Comprehensive nuclear materials 4 02 radiation damage in austenitic steels Comprehensive nuclear materials 4 02 radiation damage in austenitic steels Comprehensive nuclear materials 4 02 radiation damage in austenitic steels Comprehensive nuclear materials 4 02 radiation damage in austenitic steels Comprehensive nuclear materials 4 02 radiation damage in austenitic steels

F A Garner

Radiation Effects Consulting, Richland, WA, USA

ß 2012 Elsevier Ltd All rights reserved.

4.02.5 Evolution of Radiation-Induced Microchemistry and Microstructure 444.02.6 A Cross-Over Issue Involving Radiation-Induced Microstructural Evolution and

4.02.9.5.2 Dependence of creep and creep relaxation on neutron spectra 84

MW in Obninsk, Russia BWR Boiling water reactor CAGR Commercial Advanced Gas Reactor CANDU Registered trademark for Canadian

Deuterium Uranium Reactor DFR Dounreay Fast Reactor in Dounreay,

Scotland

33

DMTR Dounreay Materials Test Reactor in

Dounreay, Scotland

EBR-II Experimental Breeder Reactor-II in

Idaho Falls, Idaho

FFTF Fast Flux Test Facility, fast reactor in

Richland, WA

HFIR High Flux Isotope Reactor at Oak Ridge

National Laboratory

HFR High Flux Reactor in Petten, Netherlands

IASCC Irradiation-assisted stress corrosion

cracking

IGSCC Intergranular stress corrosion cracking

JMTR Japan Material Testing Reactor in Oarai,

Japan

NRU National Research Universal Reactor in

Chalk River, Canada

ORNL Oak Ridge National Laboratory:

ORR Oak Ridge Research Reactor in Oak

Ridge, Tennessee

PWR Pressurized water reactor

T/F Thermal-to-fast neutron ratio

VVER Russian acronym for water-cooled,

water moderated energetic reactor

Austenitic stainless steels are widely used as

struc-tural components in nuclear service in addition

to being employed in many other nonnuclear

engineering and technological applications The

description of these steels and their as-fabricated

properties is covered in Chapter 2.09, Properties

of Austenitic Steels for Nuclear Reactor

Applica-tions This chapter describes the evolution of both

microstructure and macroscopic property changes

that occur when these steels are subjected not only

to prolonged strenuous environments but also to the

punishing effects of radiation While various nuclear

environments involve mixtures of charged particles,

high-energy photons and neutrons, it is the latter

that usually exerts the strongest influence on the

evolution of structural steels and thereby determines

the lifetime and continued functionality of structural

components

To describe the response of austenitic stainless

steels in all neutron environments is a challenging

assignment, especially given the wide range of

neutron spectra characteristic of various neutron

devices This review of neutron-induced changes in

properties and dimensions of austenitic stainless

steels in all spectral environments has thereforebeen compiled from a series of other, more focusedreviews directed toward particular reactor types1–8and then augmented with material from a recentlypublished textbook9and journal articles It should benoted, however, that many of the behavioral char-acteristics of iron-based stainless steels followingneutron irradiation are also observed in nickel-based alloys Whenever appropriate, the similaritiesbetween the two face-centered-cubic alloy systemswill be highlighted A more comprehensive treat-ment of radiation effects in nickel-base alloys isprovided in Chapter 4.04, Radiation Effects inNickel-Based Alloys

This review is confined to the effects of neutronexposure only on the response of irradiated steels anddoes not address the influence of charged particleirradiation While most of the phenomena induced

by neutrons and charged particles are identical, thereare additional processes occurring in charged par-ticle studies that can strongly influence the results.Examples of processes characteristic of charged par-ticle simulations are the injected interstitial effect,10,11strong surface effects,12,13 dose gradients,14,15 andatypical stress states.16,17 Chapter 1.07, RadiationDamage Using Ion Beams addresses the use ofcharged particles for irradiation

Austenitic stainless steels used as fuel cladding orstructural components in various reactor types mustoften withstand an exceptionally strenuous and chal-lenging environment, even in the absence of neutronirradiation Depending on the particular reactor type,the inlet temperature during reactor operation canrange from
50 to
370C The maximum temper-ature can range from as high as 650 to 700C forstructural components in some reactor types, althoughmost nonfueled stainless steel components reachmaximum temperatures in the range of 400–550C.During operation, the steel must also withstand thecorrosive action of fission products on some surfacesand flowing coolant on other surfaces The coolantespecially may be corrosive to the steel underoperating conditions Some of these environmentalphenomena are synergized or enhanced by the effect

of neutron irradiation

Dependent on the nature of the component andthe length of its exposure, there may also be sig-nificant levels of stress acting on the component.Stress not only influences cracking and corrosion(seeChapter5.08, Irradiation Assisted Stress Cor-rosion Cracking) but can also impact the dimen-sional stability of stainless steel, primarily due to

thermal creep and irradiation creep, and also from

the influence of stress on precipitation, phase

stabil-ity, and void growth, some of which will be discussed

later However, it will be shown that neutron

irradia-tion can strongly affect both the microstructure and

microchemistry of stainless steels and high-nickel

alloys, with strong consequences on physical

proper-ties, mechanical properproper-ties, dimensional stability, and

structural integrity

Stainless steels are currently being used or have

been used as structural materials in a variety of

nuclear environments, most particularly in

sodium-cooled fast reactors, water-sodium-cooled and water-moderated

test reactors, water-cooled and water-moderated

power reactors, with the latter subdivided into light

water and heavy water types Additionally, there are

reactor types involving the use of other coolants

(helium, lithium, NaK, lead, lead–bismuth eutectic,

mercury, molten salt, organic liquids, etc.) and other

moderators such as graphite or beryllium

The preceding reactor types are based on the

fission of uranium and/or plutonium, producing

neutron energy distributions peaking at
2 MeV

prior to moderation and leakage effects that produce

the operating spectrum However, there are more

energetic sources of neutrons in fusion-derived

spectra, with the source peaking at
14 MeV and

especially from spallation events occurring at

ener-gies of hundreds of MeV, although most spallation

spectra are mixtures of high-energy protons and

neutrons It is important to note that in each of

these various reactors, there are not only significant

differences in neutron flux-spectra but also

signifi-cant differences in neutron fluence experienced by

structural components These differences in fluence

arise not only from differences in neutron flux

characteristic of the different reactor types but also

the location of the steel relative to the core For

instance, boiling water reactors and pressurized

water reactors have similar in-core spectra, but

stainless steels in boiling water reactors are located

much farther from the core, resulting in a factor of

reduction of
20 in both neutron dose rate and

accumulated dose compared to steels in pressurized

water reactors

What are the nature and origins of neutron-induced

phenomena in metals? The major underlying driving

force arises primarily from neutron collisions withatoms in a crystalline metal matrix When exposed todisplacive irradiation by energetic neutrons, the atoms

in a metal experience a transfer of energy, which iflarger than several tens of eV, can lead to displacement

of the atom from its crystalline position The ments can be in the form of single displacementsresulting from a low-energy neutron collision with asingle atom or a glancing collision with a higher energyneutron More frequently, however, the ‘primaryknock-on’ collision involves a larger energy transferand there occurs a localized ‘cascade’ of defects thatresult from subsequent atom-to-atom collisions.There are several other contributions to displace-ment of atoms from their lattice site, but these areusually of second-order importance The first ofthese processes involve production of energetic elec-trons produced by high-energy photons via thephotoelectric effect, Compton Effect, or pair produc-tion.18These electrons can then cause atomic displa-cements, but at a much lower efficiency than thatassociated with neutron-scattering events The sec-ond type of process involves neutron absorption by

displace-an atom, its subsequent trdisplace-ansmutation or excitation,followed by gamma emission The emission-inducedrecoil of the resulting isotope often is sufficient todisplace one or several atoms In general, however,such recoils add a maximum of only several percent

to the displacement process and only then in highlythermalized neutron spectra.4 One very significantexception to this generalization involving nickel will

be presented later

For structural components of various types ofnuclear reactors, it is the convention to express theaccumulated damage exposure in terms of the calcu-lated number of times, on the average, that each atomhas been displaced from its lattice site Thus, 10 dpa(displacements per atom) means that each atomhas been displaced an average of 10 times Doses

in the order of 100–200 dpa can be accumulatedover the lifetimes of some reactor components invarious high-flux reactor types The dpa concept isvery useful in that it divorces the damage processfrom the details of the neutron spectrum, allowingcomparison of data generated in various spectra,providing that the damage mechanism arises primar-ily from displacements and not from transmutation.The use of the dpa concept also relieves research-ers from the use of relatively artificial and sometimesconfusing threshold energies frequently used todescribe the damage-causing portion of the neutronspectrum Neutrons with ‘energies greater than

X MeV,’ where X is most frequently 0.0, 0.1, 0.5, or

1.0 MeV, have been used for different reactor

con-cepts at different times in history The threshold

energy of 0.1 MeV is currently the most widely

used value and is most applicable to fast reactors

where large fractions of the spectra lay below 0.5

and 1.0 MeV Many older studies employed the total

neutron flux (E > 0.0) but this is the least useful

threshold for most correlation efforts Caution

should be exercised when compiling data from

many older studies where the neutron flux was not

adequately identified in terms of the threshold

energy employed

There are rough conversion factors for

‘displace-ment effectiveness’ for 300 series austenitic steels that

are useful for estimating dpa from>0.1 MeV fluences

for both in-core or near-core spectra in most fission

spectra Examples are
7 dpa per 1022

n cm2(E >

0.1) for most in-core light water spectra with lower

in-core values of
5 dpa per 1022

n cm2(E > 0.1) formetal fueled fast reactors and
4 dpa per 1022

n cm2(E > 0.1) for oxide-fueled fast reactors.4

Such version factors should not be trusted within more

con-than (10–15%), primarily due to spatial variations

across the core resulting from neutron leakage For

fast reactor spectra, E > 1.0 conversion factors are

completely unreliable

WhenE > 1.0 fluxes are employed in light water

reactor studies, the conversion factor increases

from
7 dpa per 1022

n cm2(E > 0.1) to
14 dpaper 1022 n cm2 (E > 1.0) In Russia, a threshold

energy of >0.5 MeV is popular for light water

reactors with
9 dpa per 1022

n cm2(E > 0.5) All

of these conversion factors assume that within severalpercent pure iron is a good surrogate for 300 seriesalloys Note that other metals such as Cu, Al, W, etc.will have different conversion values arising fromdifferent displacement threshold energies and some-times different displacement contributions

A standard procedure for calculating dpa has beenpublished,19 although other definitions of dpa wereused prior to international acceptance of the ‘NRTmodel’ where the letters represent the first letter

of the three author’s last name (see Garner1 fordetails on earlier models) Caution must be exercisedwhen compiling doses from older studies wheredisplacement doses were calculated using other mod-els (Kinchin-Pease, Half-Nelson, French dpa, etc.)sometimes without clearly identifying the modelemployed Conversion factors between the NRTmodel and various older models of dpa are provided

in Garner,1but all models agree within
23%.While sometimes controversial with respect tohow far the dpa concept can be stretched to coverthe full range of spectral differences for neutron andespecially for charged particle environments, it appearsthat the dpa concept is very efficient to stretch overlight water, heavy water, fusion, and spallation spectra,providing that all energy deposition and displacementprocesses are included Note inFigure 1how well thedpa concept collapses the data on neutron-inducedstrengthening of stainless steel into one responsefunction for three very different spectra (light waterfission, pure D–T fusion and ‘beam-stop’ spallation).20

Neutron fluence, E > 0.1 MeV

dpa

100 150 200

250 LASREF, 40 C

Figure 1 Radiation-induced yield stress changes of 316 stainless steel versus (left) neutron fluence (n cm2E > 0.1 MeV), and (right) displacements per atom Reproduced from Heinisch, H L.; Hamilton, M L.; Sommer, W F.; Ferguson, P.

J Nucl Mater 1992, 191–194, 1177, as modified by Greenwood, L R J Nucl Mater 1994, 216, 29–44.

4.02.2.2 Transmutation

It is important to note that material modification by

radiation arises from two primary spectral-related

processes In addition to the neutron-induced

dis-placement of atoms there can be a chemical and/or

isotopic alteration of the steel via transmutation

With the exception of helium production,

transmuta-tion in general has been ignored as being a significant

contributor to property changes of stainless 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

pro-cesses 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

rel-evant to stainless steels, the reader is referred to

various papers on transmutation and its consequences

in different reactor spectra.5–8,18,21–23

Transmutation may be subdivided into four

cate-gories 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

trans-mutants, and ‘isotope shifts,’ the latter involving

pro-duction 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

dis-crete nuclear reactions to produce disdis-crete isotopes of

specific elements, the spallation-induced

transmuta-tion arising in accelerator-driven devices involves a

continuous distribution of every 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 if spallation

devices continue to be developed

Another consequence of spallation-relevant

trans-mutation is that the induced radioactivity per unit

mass is correspondingly much higher than that duced per dpa in other spectra The majority of thespalled fragments and their daughters/granddaugh-ters are radioactive with relatively short half-lives,leading to materials that are often much more diffi-cult to examine than materials irradiated in fissionspectra

pro-Most importantly, there is a very strong tion of hydrogen and helium in spallation spectra atlevels that are one or two orders of magnitude greaterthan produced in most fission or fusion spectra.5,6,21While there is a tendency to view displacementand transmutation processes as separate processes,

produc-it will be shown later that under some stances the two processes are strongly linked andtherefore inseparable in their action to change alloybehavior

SpectraThere are significant differences in neutron spectrafor water-cooled, sodium-cooled, and other types offission-based reactors It should be noted that there is

a conventional but slightly misleading practice todifferentiate between ‘fast’ and ‘thermal’ reactors.Thermal reactors have a significant portion of theirspectra composed of thermal neutrons Thermalizedneutrons have suffered enough collisions with themoderator material that they are in thermal equilib-rium with the vibrations of the surrounding atoms.Efficient thermalization requires low-Z materialssuch as H, D, and C in the form of water, graphite,

or hydrocarbons At room temperature the meanenergy of thermalized neutrons is 0.023 eV

The designation ‘fast’ reactor, as compared to

‘thermal’ reactor, refers to the portion of the neutronspectrum used to control the kinetics of ascent to fullpower for each type of reactor As shown later, thispractice incorrectly implies to many that fast reactorshave ‘harder’ neutron spectra than do ‘softer’ thermalreactors Actually, the opposite is true

Examples of typical flux-spectral differences infission-based reactors are shown in Figures 2–5.The local spectrum at any position is determinedprimarily by the fuel (U, Pu) and fuel type (metal,oxide, carbide, etc.), the coolant identity and density,the local balance of fuel/coolant/metal as well as theproximity to control rods, water traps, or core bound-aries Additionally, it is possible to modify the neutronspectra in a given irradiation capsule by including in it

or enclosing it with a moderator or absorber Metal

hydrides are used in fast reactors to soften the

spec-trum, while in mixed-spectrum reactors the

thermal-to-fast ratio can be strongly reduced by incorporating

elements such as B, Hf, Gd, and Eu

The most pronounced influence on neutron

spectra in fission reactors arises from the choices of

coolant and moderator, which are often the same

material (e.g., water) Moving from heavy liquid metals

such as lead or lead–bismuth to lighter metals such as

sodium leads to less energetic or ‘softer’ spectra

Use of light water for cooling serves as a

much more effective moderator Counterintuitively,

however, this leads to both more energetic and lessenergetic spectra at the same time, producing a two-peaked ‘fast’ and ‘thermal’ distribution separated by awide energy gulf at lower fluxes

Such two-peaked spectra are frequently called

‘mixed spectra.’ The ratio of the thermal and fastneutron fluxes in and near such reactors can varysignificantly with position and also with time.4Usingheavy water, we obtain a somewhat less efficientmoderator that does not absorb neutrons as easily aslight water, but one that produces an even more pro-nounced two-peak spectral distribution where thethermal-to-fast neutron ratio can be very large.These spectral differences lead to strong varia-tions between various reactors in the neutron’s ability

to displace atoms and to cause transmutation ing on the reactor size and its construction detailsthere can also be significant variations in neutronspectra and ‘displacement effectiveness’ within agiven reactor and its environs, especially wheremore energetic neutrons can leak out of the core.Examples of these variations of displacement effec-tiveness for fast reactors are shown in Figures 6and 7 Compared to fission-derived spectra, thereare even larger spectral differences in various fusion

Depend-or spallation neutron devices

The reader should note the emphasis placed here

on flux-spectra rather than simply spectra If we focusonly on light water-cooled reactors for example,there are in general three regimes of neutron flux ofrelevance to this review First, there are the relativelylow fluxes typical of many experimental reactors that

Neutron energy (MeV) 1.E - 9

1.E - 7 1.E - 5 1.E - 3 1.E - 1 1.E + 1

Neutron energy (MeV)

EBR II ORR

Figure 2 Difference in neutron flux-spectra of two

water-cooled test reactors (high-flux HFIR and lower-flux

ORR) and one high-flux sodium-cooled fast reactor (EBR-II).

Upper core plate

Figure 3 Typical neutron flux-spectra of internal

components of a pressurized water reactor, having a

thermal-to-fast neutron ratio smaller by factors of 10–20

than that of typical light water test reactors Reproduced

from Garner, F A.; Greenwood, L R In 11th International

Conference on Environmental Degradation of Materials

in Nuclear Power Systems – Water Reactors; 2003;

pp 887–909.

can produce doses of 10 dpa or less over a decade.

Second, there are moderate flux reactors that are

used to produce power that can introduce doses as

high as 60–100 dpa maximum over a 30–40 year

life-time and finally, some high-flux thermal reactors that

can produce 10–15 dpa year1in stainless steels

Most importantly, fast reactors also operate in

the high-flux regime, producing 10–40 dpa year1

Therefore, the largest amount of published

high-dpa data on stainless steels has been generated in

fast reactors Some phenomena observed at high

exposure, such as void swelling, have been found to

be exceptionally sensitive to the dpa rate, while

others are less sensitive (change in yield strength)

or essentially insensitive (irradiation creep) These

sensitivities will be covered in later sections

For light water-cooled reactors, the various flux

regimes need not necessarily involve large

differ-ences in neutron spectra, but only in flux However,

the very large dpa rates characteristic of fast reactors

are associated with a significant difference in

spec-trum This difference is a direct consequence of the

fact that fast reactors were originally designed to

breed the fissionable isotope239Pu from the relatively

nonfissile isotope 238U, which comprises 99.3% of

natural uranium

In order to maximize the breeding of 239Pu, it

is necessary to minimize the unproductive capture

of neutrons by elements other than uranium One

strategy used to accomplish this goal is to avoidthermalization of the reactor neutrons, which requiresthat no low atomic weight materials such as H2O, D2O,

Be, or graphite be used as coolants or as moderators.For this purpose, sodium is an excellent coolant with

a moderate atomic weight The use of sodium results

Axial position (cm)

Row 4 Row 2

4.0 4.5 5.0 5.5

dpa

Figure 6 Displacement effectiveness values of dpa per

1022n cm2(E > 0.1 MeV) across the small core (30 cm tall and
30 cm diameter) of the EBR-II fast reactor, showing effects of neutron leakage to soften the spectrum near the core axial boundaries Near core center (Row 2) the spectrum and displacement effectiveness are dictated primarily by the use of metal fuel, producing a maximum of
5.2 dpa per 10 22

>0.821 MeV

>0.111 MeV

Figure 5 Variation in fast and thermal fluxes in HFIR as a function of radial position at mid-core at 85 MW, also showing change in thermal population with burn-up (Source: ORNL website).

in a neutron spectrum that is nominally single-peaked

rather than the typical double-peaked (thermal and

fast) neutron spectrum found in light water or heavy

water reactors The single-peaked fast reactor

spec-trum is significantly less energetic or softer, however,

than that found in the fast peak of light water reactors

Depending on the fuel type (metal vs oxide) the

mean energy of fast reactor spectra varies from
0.8

to
0.5–0.4 MeV while light water-cooled reactors

have a fast neutron peak near
1.2 MeV

One consequence of attaining successful breeding

conditions is that the spectrum-averaged

cross-section for fission is reduced by a factor of 300–400

relative to that found in light water spectra To reach

a power density comparable to that of a light water

power-producing reactor, the fast reactor utilizes two

concurrent strategies: increases in fissile enrichment

to levels in the order of 20% or more, and most

importantly, an increase in neutron flux by one or

two orders of magnitude

Thus, for a given power density, the fast reactor

will subject its structural materials to the punishing

effects of neutron bombardment at a rate that is

several orders of magnitude greater than that in

light water reactors At the same time, however, the

softer ‘fast’ spectrum without thermalized neutrons

leads to a significant reduction in transmutation

compared to typical light water spectra, at least for

stainless steels and nickel-base steels

Stainless SteelsFor most, but not all fission-derived spectra, stainlesssteels are relatively immune to transmutation, espe-cially when compared to other elements such asaluminum, copper, silver, gold, vanadium, tungsten,and rhenium,5,21,24–27 each of which can rapidlybecome two or three component alloys via transmu-tation arising from thermal or epithermal neutrons.Whereas the properties of these metals are particu-larly sensitive to formation of solid transmutationproducts, stainless steels in general do not changetheir composition by significant amounts compared

to preexisting levels of impurities, but significantamounts of helium and hydrogen can be produced

in fission-derived spectra, however

In stainless steels the primary transmutantchanges that arise in various fission and fusion reactorspectra involve the loss of manganese to form iron,loss of chromium to form vanadium, conversion ofboron to lithium and helium, and formation of heliumand hydrogen gas.4,28While each of these changes insolid or gaseous elements are produced at relativelysmall concentrations, they can impact the evolution

of alloy properties and behavior

For instance, vanadium is not a starting nent of most 300 series stainless steels, but whenincluded it participates in the formation of carbide

compo-BC 6.0

5.0 dpa

displacement effectiveness values are lower, determined primarily by the absence of fuel and the balance of sodium and steel.

precipitates that change the distribution and

chemi-cal activity of carbon in the alloy matrix Carbon

plays a number of important roles in the evolution

of microstructure1and especially in grain boundary

composition The latter consideration is very

impor-tant in determining the grain boundary cracking

behavior, designated irradiation-assisted stress

corro-sion cracking (IASCC), especially with respect to the

sensitization process.29

The strong loss of manganese in highly

therma-lized neutron spectra has been suggested to degrade

the stability of insoluble MnS precipitates that tie up

S, Cl, and F, all of which are elements implicated

in grain boundary cracking.30 Late-term

radiation-induced release of these impurities to grain

bound-aries may participate in cracking, but this possibility

has not yet been conclusively demonstrated

In some high-manganese alloys such as XM-19

manganese serves to enhance the solubility of

nitrogen which serves as a very efficient matrix

strengthener In highly thermalized spectra the loss

of manganese via transmutation has been proposed to

possibly lead to a decrease in the strength of the alloy

and perhaps to induce a release of nitrogen from

solution to form bubbles.31

The overwhelming majority of published

trans-mutation studies for stainless steels and high-nickel

alloys steels have addressed the effects of He/dpa

ratio on mechanical properties and dimensional

instabilities Much less attention has been paid to

the effect of H/dpa ratio based on the long-standing

perception that hydrogen is very mobile in metals

and therefore is not easily retained in steels at

reactor-relevant temperatures As presented later,

this perception is now known to be incorrect,

espe-cially for water-cooled reactors

The focus of most published studies concerned

the much higher helium generation rates anticipated

in fusion spectra (
3–10 appm He/dpa) compared to

the lower rates found in fast reactors (
0.1–0.3 appm

He/dpa).32It was later realized that in some highly

thermalized test reactors, such as HFIR, very large

generation rates could be reached (
100 appm

He/dpa), and even in pressurized water reactors

the rate could be very high (
15 appm He/dpa).33

In heavy water reactors the rate can be much larger,

especially in out-of-core regions.34,35

While some helium arises from (n, a) reactions

with thermal and epithermal neutrons interacting

with the small amounts of boron found in most

stainless steels, the major contribution comes initially

from high-energy threshold-type (n, a) reactions

with the major alloy components This type of tion occurs only above high neutron threshold ener-gies (>6 MeV) Figure 8 shows that nickel is themajor contributor to helium production by (n, a)reactions,36 and thus the helium generation ratescales almost directly with nickel content for a largenumber of commercial steels

reac-A similar behavior occurs for production ofhydrogen by transmutation via high-energy neutrons,where nickel is also the major source of hydrogencompared to other elements in the steel.4,7 In thiscase, the threshold energy is around 1 MeV with58

Ni being the major contributor

This generality concerning nickel as the majorsource of He and H is preserved in more energeticfusion-derived spectra, although the He/dpa andH/dpa generation rates in fusion spectra are muchlarger than those of fast reactor spectra Whenmoving to very energetic spallation-derived neutronand proton spectra, however, the observation thatnickel accounts for most of the helium and hydrogen

is no longer correct Iron, nickel, chromium, cobalt,and copper produce essentially the same amounts ofhelium and hydrogen for energies above
100 MeV

as shown inFigure 9.6Another very important helium-generation pro-cess also involves nickel Helium is producedvia the two-step 58Ni(n, g)59Ni(n, a)56Fe reactionsequence.37,38 This sequence operates very strongly

in mixed-spectrum reactors 59Ni is not a naturallyoccurring isotope and is produced from58Ni Thus,this helium contribution involves a delay relative to

0.14 0.12 0.10

Fe Ti

Cr

20 Energy (MeV)

Figure 8 Cross-sections for (n, a) reactions as a function

of neutron energy for common elements used in stainless steels Reproduced from Mansur, L K.; Grossbeck, M L.

J Nucl Mater 1988, 155–157, 130–147 Nickel dominates the production of helium at higher neutron energies.

that of single-step threshold (n, a) reactions Since

both steps of the sequence involve cross-sections that

increase with decreasing energy and the second step

exhibits a resonance at 203 eV, the generation rate

per dpa in fast reactors increases near the core

boundaries and out-of-core areas

It is in thermalized neutron spectra characteristic

of light and heavy water-cooled reactors, however,

where the59Ni(n, a) reaction can produce He/dpa

generation rates that are significantly larger than

those characteristic of fusion-derived spectra

Nickel has five naturally occurring stable isotopes

with58Ni comprising 67.8% natural abundance,60Ni

comprising 26.2%, and
6.1% total of61

Ni,62Ni, and64

Ni There is no natural59Ni or63Ni at the beginning

of radiation During irradiation in a highly

therma-lized neutron spectrum, all nickel isotopes are

strongly transmuted, primarily to the next higher

isotopic number of nickel 59Ni has a half-life of

76 000 years and is progressively transmuted to60Ni

while58Ni is continuously reduced in concentration

Therefore, the59Ni concentration rises to a peak at a

thermal neutron fluence of 4 1022

n cm2where the59/58 ratio peaks at
0.04 and then declines, as shown

inFigure 10

This transmutation sequence in nickel is an

exam-ple of the isotopic shift category of transmutation

defined earlier For other elements used to make

stain-less steels, there are no consequences to such a shift

since the total amount of the element is unchanged

and isotope shifts induce no significant consequences.However, in the case of nickel there is an intimatelinkage between the displacement and transmutationprocesses that arises from the isotope shift

The recoil of the 59Ni upon emission of thegamma ray produces only about five displacementsper event, and usually is not a significant addition tothe displacement dose However, the isotope 59Niundergoes three strong reactions with thermal andresonance (
0.2 keV) neutrons, two of which areexceptionally exothermic and can significantly add

to the dpa level These reactions, in order ofhighest-to-lowest thermal cross-section, are (n, g) toproduce60Ni, followed by (n, a) and (n, p) to producehelium and hydrogen, respectively

Even at relatively low thermal-to-fast neutronratios, the reaction sequence can produce significantamounts of helium For example, He/dpa ratios inthe order of
3–8 appm dpa1 can be experiencedalong the length of a 316 stainless baffle bolt inthe baffle-former assembly of a pressurized water

Figure 10 Transmutation-induced evolution of three nickel isotopes during irradiation in thermalized neutron spectra Reproduced from Garner, F A.;

Greenwood, L R In 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2003; pp 887–909.

Reproduced from Garner, F A.; Griffiths, M.; Greenwood,

L R.; Gilbert, E R In Proceedings of the 14th International Conference on Environmental Degradation

of Materials in Nuclear Power Systems – Water Reactors; American Nuclear Society, 2010;

Figure 9 Measured amount of helium in alloys and pure

metals that were irradiated by a mixed spectrum of high

energy neutrons and protons produced by 800 MeV proton

irradiation of tungsten rods There is some significant

uncertainty in the dpa assignment for Inconel 718 at the

highest dose Otherwise the He/dpa ratio appears to be

independent of composition Reproduced from Garner, F A.;

Oliver, B M.; Greenwood, L R.; James, M R.; Ferguson, P D.;

Maloy, S A.; Sommer, W F J Nucl Mater 2001, 296,

66–82.

reactor4,33,39while comparable rates in fast reactors

are in the order of 0.1–0.2 appm dpa1 In

therma-lized spectra the latter two reactions can quickly

overwhelm the gas production produced by nickel

at high neutron energies

As mentioned previously, the thermal neutron

reactions of59Ni are quite exothermic in nature and

release large amounts of energy, thereby causing

increases in the rate of atomic displacements, and

con-comitant increases in nuclear heating rates Nuclear

heating by elastic collisions with high-energy

neu-trons is usually too small to be of much significance

The59Ni(n, a) reaction releases 5.1 MeV,

produc-ing a 4.8 MeV alpha particle which loses most of its

energy by electronic losses, depositing significant

thermal energy but producing only
62 atomic

dis-placements per each event However, the recoiling

56

Fe carries 340 keV, which is very large compared to

most primary knock-on energies, and produces an

astounding
1701 displacements per event

The thermal (n, p) reaction of 59Ni produces

about one proton per six helium atoms, reflecting

the difference in thermal neutron cross-sections of

2.0 and 12.3 barns, and is somewhat less energetic

(1.85 MeV), producing a total of
222 displacements

per event.7,40 In addition, approximately five

dis-placed atoms are created by each emission-induced

recoil of60Ni This reaction occurs at six times higher

rate compared to the 59Ni(n, a) reaction, resultingfrom a thermal neutron cross-section of 77.7 barns Ineffect, the dpa rate increases during irradiation due tothe three 59Ni reactions even though the neutronflux-spectrum may not change

The major point here is that use of standardizedcomputer codes to calculate dpa does not trackshifts in isotopic distribution and therefore willunderpredict the dpa level when59Ni production is

an important consideration

A strong example of this time-dependent increase

in dpa rate in highly thermalized light water spectra

is shown for pure nickel inFigure 11for a to-fast ratio of 2.0 Note that the calculated increase

thermal-in this figure addresses only the 59Ni(n, a) reaction.Additional increases occur as a result of the59Ni(n, p)and59Ni(n, g) reactions, resulting in almost doubling

of dpa by the three59Ni reactions before a calculateddose of
40 dpa is attained

Recently, however, an even stronger example ofthe linkage of the59Ni transmutation effect and thedisplacement process has been observed.34,35In-corethermal-to-fast ratios in heavy water-moderatedreactors such as CANDUs are in the order of
10,but far from the core the ratio can be near
1000.Compression-loaded springs constructed of high-nickel alloy X-750 were examined after 18.5 years ofoperation far from the core and were found to be

of the 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; American Nuclear Society, 2010; pp 1344–1354.

completely relaxed Calculating the 59Ni

contribu-tion, it was deduced that full relaxation occurred

in
3–4 years rather than the 650–700 years one

would predict based on dpa calculated without taking

into account the59Ni contribution

Therefore, in this case59Ni contributed
95% of

the dpa Additionally, 1100 appm of helium was

cal-culated to have been produced at the mid-section of

the spring in
3 years, with
20 000 appm helium

having been produced when the spring was examined

after 18.5 years of exposure

There is another consequence of the59Ni sequence

that causes the temperature to increase during

irradia-tion At the peak59Ni level reached at 4 1022

n cm2,the nuclear heating rates from the energetic (n, a) and

(n, p) reactions are 0.377 and 0.023 W g1of nickel,

significantly larger than the neutron heating level of

0.03 W g1 of natural nickel Thus, an increase in

nuclear heating of
0.4 W g1 of nickel must be

added to the gamma heating rate at the peak 59Ni

level Fractions of the peak heating rates that are

pro-portional to the current59Ni level should be added at

nonpeak conditions Depending on the nickel level of

the steel and the level of gamma heating, which is the

primary cause of temperature increases in the interior

of thick plates, this additional heating contribution

may or may not be significant

Gamma heating is also a strong function of the

thermal-to-fast (T/F) neutron ratio and the neutron

flux, being
54 W g1in the center of the HFIR test

reactor where the T/F ratio is
2.0 In pressurized

water reactors at the austenitic near-core internals,

however, the T/F ratios are lower by a factor of 2–10,

depending on location, and the gamma heating rates

in the baffle-former assembly are
1–3 W g1 In this

case, an additional 0.4 W g1of nuclear heating can

be a significant but time-dependent addition to total

heating, especially for high-nickel alloys

It should be noted that thermal neutron populations

can vary during an irradiation campaign with

conse-quences not only on59Ni production but also on gamma

heating levels In PWRs boric acid is added to the water

as a burnable poison at the beginning of each cycle As

the 10B burns out the thermal neutron population

increases, leading to an increase in gamma heating and

transmutation.3,4Over successive cycles there is a

saw-tooth variation of gamma heating rate in the

baffle-former assembly and therefore in DT, with the latter

reaching values as large as20C in the worst case.

Additionally, another concern may arise in that

small radiation-induced nickel-rich phases such as

g0, Ni-phosphides, and G-phase may become less

stable This concern arises due to cascade-induced

dissolution as the 56Fe from the 59Ni(n, a) reactionrecoils within the precipitates, thereby altering thephase evolution in thermalized neutron spectra com-pared to nonthermalized spectra typical of fast reac-tors These precipitates are known to form as a directresult of irradiation and contribute to hardening,swelling, and irradiation creep processes.1 The size

of these precipitates at PWR-relevant temperatures(290–400C) is often comparable to or smaller thanthe
80 nm range of the recoiling56

Fe atom.Finally, another significant source of helium canarise from the implantation of energetic heliumresulting from collisions with neutrons into the sur-face layers of helium gas-pressurized or gas-cooledcomponents, often involving hundreds and oftenthousands of appm of injected helium In gas-cooledreactors helium injection has been investigated as apossible degradation mechanism of alloy surfaces.41

In fast reactor fuel cladding helium was found to

be injected into the inner surface, coming from twomajor sources, ternary fission events (two heavy fis-sion fragments plus an alpha particle) in the fuel andfrom helium recoiling from the pins’ helium covergas as a result of collisions with neutrons.42

The injection rates from these two sources ofinjected helium are slowly reduced during irradia-tion, however, as heavy fission gases build up in thespace between the fuel pellet and the cladding.These gases slow down the energetic helium atoms,reducing their energy sufficiently to prevent most ofthem from reaching the cladding Helium injection

at high levels was also found on the inner surface ofhelium-pressurized creep tubes.42 Although heliuminjection tends to saturate in fuel pin cladding withincreasing dose, it does not saturate in pressurizedtubes due to the lack of increasing fission gases toreduce the range of helium knock-ons in the gasphase

Some studies have cited this early source of helium

as contributing to the embrittlement of fuel pin ding and its poor performance during transient heatingtests,43although more recent studies have linked themajor mechanism to delayed grain boundary attack bythe fission products cesium and tellurium.44,45

microchemistry commences that is dependent

pri-marily on the alloy starting state, the dpa rate, and

the temperature, and secondarily dependent on

vari-ables such as He/dpa rate and applied or internally

generated stresses

In general, the starting microstructure and

micro-chemistry of the alloy determine only the path taken

to the radiation-defined quasi-equilibrium state,

and not the final state itself If an alloy experiences

enough displacements, it effectively forgets its

start-ing state and arrives at a destination determined only

by irradiation temperature and dpa rate This

quasi-equilibrium or dynamic-quasi-equilibrium state consists of

microstructural components existing at relatively

fixed densities and size distributions, but individual

dislocations, loops, precipitates, or cavities at any one

moment may be growing, shrinking, or even

disap-pearing by shrinkage or annihilation

The displacement process produces two types

of crystalline point defects, vacant crystalline

posi-tions (vacancies) and displaced atoms in interstitial

crystalline positions (interstitials) These two defect

types are both mobile, but move with different

dif-fusional modes and at vastly different velocities,

with interstitials diffusing much faster than

vacan-cies Therefore it is obvious that all diffusion-driven

processes will be strongly affected by radiation

Both defect types have the ability to recombine

with the opposite type (annihilation) or to form

agglomerations of various types and geometries

These agglomerations and their subsequent evolution

alter both the microstructure and elemental

distribu-tion of the alloy

It is important to note that interstitial

agglomera-tions are constrained to be two-dimensional, while

vacancies can agglomerate in both two-dimensional

and three-dimensional forms This dimensional

dis-parity is the root cause of the void swelling

phenom-enon covered in a later section

The developing ensemble of various defect

agglomerations with increasing dose induces

signifi-cant time-dependent and dose-dependent changes

in physical and mechanical properties, as well as

resulting in significant dimensional distortion Most

importantly, under high displacement rates stainless

steels and other alloys are driven far from equilibrium

conditions as defined in phase diagrams, affecting not

only phase stability but also all physical, mechanical,

and distortion processes that involve phase changes in

their initiation or evolution

During irradiation, the phase evolution can be

significantly altered, both in its kinetics and in the

identity and balance of phases that form.46,47Phases

can be altered in their composition from that found inthe absence of irradiation, and new phases can formthat are not found on the equilibrium phase diagram

of a given class of steels In 300 series stainless steelsthese new or altered phases have been classified asradiation-induced phases, radiation-modified phases,and radiation-enhanced phases.48–51These classifica-tions are equally applicable to phases formed in otherclasses of steel

Radiation-induced alterations of microstructureand microchemistry occur because new driving forcesarise that do not occur in purely thermal environ-ments The first of these new driving forces is thepresence of very large supersaturations of pointdefects, especially at relatively low irradiation tem-peratures (250–550C) Not only are vacanciespresent in uncharacteristically high levels, therebyaccelerating normal vacancy-related diffusional pro-cesses, but interstitials are also abundant Solutes thatcan bind with either type of point defect tend to flowdown any microstructurally induced gradient of thatdefect, providing a new mechanism of solute segre-gation referred to as solute drag.52 This mechanismhas been proposed to be particularly important forbinding of smaller solute atoms such as P and Si, andsometimes Ni, with interstitials

A second new driving force is the inverseKirkendall effect53whereby differences in elementaldiffusivity via vacancy exchange lead to segregation

of the slowest diffusing species at the bottom of induced vacancy gradients This mechanism is par-ticularly effective in segregating nickel in austeniticFe–Cr–Ni alloys at all sinks which absorb vacancies,leading to nickel-rich shells or atmospheres ongrain boundaries and other preexisting or radiation-produced microstructural sinks This type of segre-gation arises because the elemental diffusivities ofFe–Cr–Ni alloys are significantly different, withDCr> DFe> DNiat all nickel levels.54–57

sink-A third new driving force results from the action ofthe other two driving forces when operating onmicrostructural sinks that are produced only in irra-diation environments These are Frank interstitialloops, helium bubbles, and voids that may have devel-oped from helium bubbles Precipitates are oftenobserved to form and to co-evolve on the surface ofsuch radiation-induced sinks Examples of typicalradiation-induced microstructures in stainless steelsare shown inFigures 12–15 These microstructuralsinks have been implicated as participating in theevolutionary path taken by the precipitates andthereby influencing the microchemical evolution ofthe matrix.1,58–60

Minor solute elements such as Si and P have muchhigher diffusivities than those of Fe, Ni, and Cr andalso participate in the segregation process Addition-ally, these elements increase the diffusivities of themajor elements Fe, Ni, and Cr.54

When the solute drag mechanism, operatingbetween interstitials and smaller size Si and P atoms,combines with nickel segregation via the inverseKirkendall mechanism, phases that are rich in nickel,silicon, or phosphorus often form (g0, G-phase andNi2P for example), although in 300 series stainlesssteels these phases do not form thermally Otherphases that are normally stable in the absence ofradiation (carbides, intermetallics) can be forced dur-ing irradiation to become enriched in these elements.1The removal of nickel, silicon, and phosphorusfrom the matrix by radiation-induced precipitationexerts a large effect on the effective vacancy diffusiv-

Figure 12 Frank loops observed in a 316 stainless flux thimble from a PWR power reactor (a) 70 dpa, 315C and (b) 33 dpa,

290C imaged edge-on on one set of the four (111) planes using the dark-field relrod technique Reproduced from Edwards, D J.; Garner, F A.; Bruemmer, S M.; Efsing, P G J Nucl Mater 2009, 384, 249–255 The image in (c) is from Frank loops that are slightly inclined to the beam direction imaged using a relrod in the diffraction pattern.

G-phase

50 nm

Figure 13 Electron micrograph of radiation-induced

voids in annealed ‘PCA’ stainless steel irradiated in the

ORR water-cooled test reactor at 500C to 11 dpa.

The largest voids have radiation-induced G-phase

particles attached to them that are rich in Ni, Si, and Ti.

Reproduced from Maziasz, P J J Nucl Mater 1989,

169, 95–115.

shown to exert an even larger effect on the effective

vacancy diffusivity57 and its removal into Ni2P and

other precipitates has a strong influence on matrix

diffusion Silicon is the next most effective element

on a per atom basis As the effective vacancy diffusioncoefficient falls with decreasing matrix levels of Ni,

Si, and P, conditions for void nucleation becomemore favorable

The radiation-induced evolution of diffusionalproperties has been strongly implicated in determin-ing the transient duration before void swelling accel-erates.1 This evolution often does not necessarilyproceed by only one path but occurs in several inter-active stages Some phases such as nickel phosphidesand TiC, especially when precipitated on a veryfine scale, are thought to be beneficial in resistingthe evolution of nickel silicide type phases.59,62,63

It has been shown, however, that continued induced segregation eventually overwhelms thesephases by removing critical elements such as Ni and

radiation-Si from solution, causing their dissolution andreplacement with nickel-rich and silicon-rich phasesthat coincide with accelerated swelling.63–65

In high-nickel alloys that normally form the g0and

g00 ordered phases, irradiation-induced segregationprocesses do not significantly change the identity orcomposition of the phases, but can strongly changetheir distribution, dissolving the original distributionbut plating these phases out on voids, dislocations,and grain boundaries, with the latter often leading tosevere grain boundary embrittlement.66,67

The original dislocation microstructure quicklyresponds to mobile displacement-generated pointdefects, increasing their mobility and leading toreductions in dislocation density and distribution

in the cold-worked steels most frequently used forfuel cladding and structural components.1 Thesedislocations are quickly replaced by new micro-structural components, often at very high densities,with two-dimensional interstitial Frank loops firstdominating the microstructure, then generating newline dislocations via unfaulting and interaction ofloops In well-annealed alloys there are very few pre-existing dislocations but the same radiation-inducedloop and dislocation processes occur, eventuallyreaching the same quasi-equilibrium microstructurereached by cold-worked alloys

At lower temperatures found in water-cooledtest reactors especially, the microstructural featuresappear to be three-dimensional vacancy clusters

or stacking fault tetrahedra and two-dimensionalvacancy or interstitial platelets, which are probablyalso small dislocation loops These ‘defect clusters’ attemperatures below
300C are usually too small to

be easily resolved via conventional transmission

Figure 15 Reverse contrast image showing void and line

dislocation microstructure in Fe–10Cr–30Mn model alloy

irradiated in FFTF fast reactor to 15 dpa at 520C Average

void sizes are
40 nm Reproduced from Brager, H R.;

Garner, F A.; Gelles, D S.; Hamilton, M L J Nucl Mater.

1985, 133–134, 907–911 Frank loops have unfaulted to

produce a line dislocation network whose segments end

either on void surfaces or on upper and lower surfaces of

the thin microscopy specimen The voids are coated with

ferrite phase due to Mn depletion from their surfaces via the

Inverse Kirkendall effect.

50 nm

Figure 14 Void swelling (
1%) and M 23 C 6 carbide

precipitation produced in annealed 304 stainless steel after

irradiation in the reflector region of the sodium-cooled

EBR-II fast reactor at 380C to 21.7 dpa at a dpa rate of

0.84  10 7 dpa s1 Reproduced from Garner, F A.;

Edwards, D J.; Bruemmer, S M.; et al In Proceedings,

Fontevraud 5, Contribution of Materials Investigation to the

Resolution of Problems Encountered in Pressurized Water

Reactors; 2002; paper #22 Dislocations and dislocation

loops are present but are not in contrast.

electron microscopy and are often characterized as

either ‘black dots’ or ‘black spots.’ These dots are

generally thought to be very small Frank interstitial

loops

The cluster and dislocation loop evolution is

fre-quently concurrent with or followed by the loss

or redistribution of preexisting precipitates Most

importantly, new radiation-stabilized precipitates at

high density often appear with crystal structure and

composition that are not found on an equilibrium

phase diagram for austenitic steels

As a consequence of these various processes the

microstructure at higher doses often develops very

high densities of crystallographically faceted,

vacuum-filled ‘cavities’ called voids, thought to nucleate on

helium clusters formed by transmutation, although

residual gases in the steel often help nucleate voids

at lower concentrations Voids have frequently been

observed in charged particle irradiations where no

helium was introduced

The void phenomenon is not a conservative process and the metal begins to ‘swell’

volume-as the microscopic voids in aggregate contribute

to macroscopic changes in dimension, sometimesincreasing the metal volume by levels of many tens

of percent

Concurrently, the dislocation microstructureresponds to the local stress state, moving mass via avolume-conservative process designated irradiationcreep In general, irradiation creep is not a directlydamaging process but it can lead to componentfailures resulting from distortion that causes localblockage of coolant flow or strong postirradiationwithdrawal forces Both swelling and irradiationcreep are interrelated and are interactive processesthat can produce significant distortions in component

examples of such distortion.68,69Eventually, the microstructural/microchemicalensemble approaches a quasi-equilibrium condition

Figure 16 (top) Spiral distortion of 316-clad fuel pins induced by swelling and irradiation creep in an FFTF fuel assembly where the wire wrap swells less than the cladding Reproduced from Makenas, B J.; Chastain, S A.; Gneiting, B C.

In Proceedings of LMR: A Decade of LMR Progress and Promise; ANS: La Grange Park, IL, 1990; pp 176–183; (middle) Swelling-induced changes in length of fuel pins of the same assembly in response to gradients in dose rate, temperature, and production lot variations as observed at the top of the fuel pin bundle Reproduced from Makenas, B J.; Chastain, S A.; Gneiting, B C In Proceedings of LMR: A Decade of LMR Progress and Promise; ANS: La Grange Park, IL, 1990; pp 176–183; (bottom) swelling-induced distortion of a BN-600 fuel assembly and an individual pin where the wire swells more than the cladding Reproduced from Astashov, S E.; Kozmanov, E A.; Ogorodov, A N.; Roslyakov, V F.; Chuev, V V.; Sheinkman, A G In Studies of the Structural Materials in the Core Components of Fast Sodium Reactors; Russian Academy

of Science: Urals Branch, Ekaterinburg, 1984; pp 48–84, in Russian.

or ‘saturation’ state, usually at less than 10 dpa for

mechanical properties but at higher doses for swelling

As a consequence, the mechanical properties tend to

stabilize at levels depending primarily on

tempera-ture and to a lesser extent on dpa rate The two major

deformation processes, swelling and irradiation

creep, do not saturate but reach steady-state

defor-mation rates when quasi-equilibrium microstructures

are attained This coupling of saturation

microstruc-ture with steady-state behavior has been

character-ized as ‘persistence.’70

Interestingly, the saturation states of each property

change are almost always independent of the starting

thermal–mechanical state of the material.1,70,71 If

irradiation continues long enough, the memory of

the starting microstructural state and the associated

mechanical properties is almost completely lost The

only deformation-induced microstructural component

that succeeds in resisting this erasure process is that of

preexisting, deformation-induced twin boundaries

If this quasi-equilibrium is maintained to higher

neutron exposure no further change occurs in the

steel’s mechanical properties However, some slowly

developing second-order processes are nonsaturable

and are often nonlinear Eventually, these processes

force the system to jump toward a new

quasi-equilibrium These new states usually arise from

either the microstructural or microchemical

evolu-tion, with voids dominating the former and the

latter involving continued segregation, continued

transmutation, or a combination of these factors.70–72

A number of such late-stage changes in

quasi-equilibrium state are discussed later in this paper

Radiation-Induced Microstructural Evolution and TransmutationRecently, it been discovered that significant levels ofhydrogen can be stored in bubbles and voids in bothstainless steels and pure nickel when the hydrogen iscogenerated with helium, especially in light waterspectra where there are also environmental sources

of hydrogen.73–75It was shown in these studies thatthis phenomenon is a direct result of the59Ni nuclearreactions Previously, it was a long-standing percep-tion that such storage could not occur at reactor-relevant temperatures

The retained hydrogen levels are in significantexcess of the levels predicted by Sievert’s Law andappear to be increasing with both cavity volumeand neutron fluence Since these gases are known toassist in nucleation and stabilization of cavities, it isexpected that the nonlinear59Ni reactions discussedearlier may lead to a rapidly developing, nonlinear,cavity-dominated microstructure in stainless steelsirradiated at temperatures characteristic of pressur-ized water reactors

Figure 17presents such a microstructure observed

in a PWR flux thimble tube (cold-worked 316 less steel) at
70 dpa and 330C.76

stain-There is a veryhigh density (>1017

cm3) of nanocavities with meters<3 nm in both the alloy matrix and especially

dia-on grain boundaries The measured cdia-oncentratidia-ons of

600 appm He and 2500 appm H in this specimen arethought to reside primarily within the cavities Mostimportantly, these cavities are essentially invisible

He and 2500 appm H Note the high density of cavities on the grain boundary.

under well-focused imaging conditions and can only

be imaged using very large levels of under-focus

This implies that previous studies on similar

mate-rials may have overlooked such cavity-dominated

structures

When this specimen and near-identical specimens

were subjected to slow strain rate testing after

irradi-ation, the fracture surface was indicative of
100%

intergranular stress corrosion cracking (IGSCC),

with lower doses and gas levels producing

propor-tionally less IGSCC.77As hydrogen is known to be a

contributor to grain boundary cracking, it appears

plausible that hydrogen storage may accelerate the

cracking process and that higher exposure will lead to

an increasing susceptibility to cracking This issue

may therefore become increasingly important as

PWRs previously licensed for 40 years are being

considered for life extension to 60 and possibly

80 years

in Mechanical PropertiesLong before the onset of significant phase evolution

or void swelling is observed, the first manifestation ofthe radiation-induced microstructural/microchemicalevolution appears in changes of the mechanical prop-erties As shown in Figure 18 the stress–strain dia-grams of stainless steels begin to change significantlyeven at very low dpa levels The strength of the alloyincreases, the elongation decreases, and there is aprogressive decrease in work-hardening This behav-ior is dependent somewhat on test temperature but isnot very sensitive to neutron spectrum

Movement of dislocations in metals duringdeformation following irradiation is impeded by themicrostructural components produced by radiation(dislocations, dislocation loops, voids, bubbles, preci-pitates) and therefore the strength of annealed steel

0.0 dpa 1.36 dpa

0.01 dpa 0.1 dpa

0.1 dpa 0.78 dpa

288C, and (c) annealed EC316LN irradiated in the LANSCE spallation neutron and proton spectrum at 60–100C and tested

at 25C Reproduced from Kim, J W.; Byun, T S J Nucl Mater 2010, 396, 10–19.

increases The strength increase usually saturates at

relatively low exposure levels (<10 dpa) as shown

inFigure 19, reflecting a similar saturation of

micro-structural densities Since the concentration of

most radiation-induced microstructural components

decreases with increasing temperature above
300C,

one would expect that the saturation strength would

also decrease with increasing temperature, as is shown

inFigures 20–23

Irradiation of cold-worked steels also leads tostrengthening at lower temperatures but softeningcan occur at higher temperatures if the saturationstrength level at a given temperature is below thestarting strength, as seen inFigure 21 Most impor-tantly, both annealed and cold-worked steels con-verge to the same saturation level when irradiated

at the same dpa rate and temperature as seen in

Figure 22.78Similar convergence behavior has been observed

in the evolution of microhardness.79 Note also thatradiation-induced changes in strength are roughlyindependent of composition within the annealed

300 series stainless steels, especially at lower tion temperatures, as shown byFigure 19

irradia-Such convergence behavior has been observedmany times, but there are exceptions; for example,cold-worked steels converge in their notch tensilestrength, but not to the level reached by annealedsteels.80Such behavior is usually observed in steels thattwin heavily during deformation and were irradiated atlow temperatures that resist recrystallization Twinboundaries are not easily erased by displacements,

so their hardening contribution persists

Concurrent with an increase in radiation-inducedhardening is a loss of ductility,81–83 as shown inFigures 23 and 24

The concept of saturation or persistence ofmechanical properties, especially with respect to tem-perature, applies to the most recent irradiationtemperature, as demonstrated by comparing isother-mal and nonisothermal histories In Figure 25 themechanical properties of three model alloys are seen

to converge during isothermal irradiation withoutbeing affected by composition, He/dpa ratio, andmechanical starting state.84 In Figure 26, however,

an early detour in temperature led to differencesfrom isothermal behavior, but these differences dis-appeared when the intended isothermal temperaturewas reestablished.84

Previous saturation states are soon forgotten, ally by
5 dpa, but only if the hardening componentsare easily erased and replaced at the new tempera-ture If hardening arises primarily from dislocationloops and dislocations, this condition is easily met

usu-If the primary hardening arises from a fine density

of voids and especially bubbles produced at lowertemperatures, then the microstructural memory can-not be easily erased, even at much higher tempera-tures An example is shown inFigure 27where a series

of Fe–Cr–Ni ternary austenitic alloys were irradiated

at 400 and 500C in ORR at high He/dpa ratios

316LN

316

304, 304L

316, 316LN PCA 200

Figure 19 Strengthening of various annealed 300 series

stainless steels versus dpa in various water-cooled reactors

at relatively low temperatures (280–330C) Reproduced

from Pawel, J P.; Ioka, I.; Rowcliffe, A F.; Grossbeck, M L.;

Jitsukawa, S In Effects of Radiation on Materials: 18th

International Symposium; ASTM STP 1325; 1999;

pp 671–688 At these temperatures strengthening saturates

at
10 dpa.

Frank loops Cavities, precipitates Defect clusters

Figure 20 Radiation-induced strengthening of annealed

300 series steels versus irradiation temperature and the

microstructural components causing the strengthening.

Note the peak strengthening at 300C followed by a decline

at higher temperatures Reproduced from Pawel, J P.;

Ioka, I.; Rowcliffe, A F.; Grossbeck, M L.; Jitsukawa, S.

In Effects of Radiation on Materials: 18th International

Symposium; ASTM STP 1325; 1999; pp 671–688.

(27–58 appm dpa1) and 395 and 450C in EBR-II atvery low He/dpa ratios (0.7–1.2 appm dpa1).85Note that there are very significant differences inhardening observed between the two experimentsand that the differences arose primarily from a verylarge difference in cavity density, a difference thatwas too large to be explained in terms of heliumcontent alone It was later shown that the ORRexperiment suffered a very large number (237 over

2 years) of unrecognized negative temperature backs of 1–2 h, with decreases varying from 50 to

set-500C.86 Even though the total dpa accumulatedduring these setbacks was only
1% of the totaldose, the frequent bloom of high densities of smallFrank loops at lower temperatures provided a verylarge periodic increase in nucleation sites for heliumbubbles on the new Frank loops that significantlystrengthened the matrix The loops could subse-quently dissolve but the bubbles could not

In addition to temperature, the most prominentirradiation variable is the dpa rate and it is knownthat the microstructural densities, especially Frankloops and voids, are known to increase in concentra-tion as the dpa rate increases Various radiation-stablephases such as w0are also known to be flux-sensitive,while other phases such as carbides and intermetallicsare more time-sensitive.1

Thus, it is not surprising that some sensitivity todpa rate might be observed in strength properties, as

1100 1000 900 800 700 600

Figure 22 Influence of temperature and neutron exposure

on evolution of yield strength in both annealed and 20%

cold-worked AISI 316 irradiated in EBR-II, showing that the

saturation strength level is independent of starting

condition, converging at doses of 5–15 dpa Reproduced

from Garner, F A.; Hamilton, M L.; Panayotou, N F.;

Johnson, G D J Nucl Mater 1981, 103 and 104,

803–808.

suggested by the behavior shown inFigure 28where

both the transient rate of strength rise and saturation

strength appear to increase with increasing dpa rate

Unfortunately, this figure does not represent a single

variable comparison, and by itself is not sufficiently

convincing evidence of flux sensitivity The data

shown inFigure 29is much closer to a single variable

comparison, indicating that the transient rise may or

not be somewhat flux-sensitive, depending on the

details of the microstructural evolution of each alloy

The authors of this study used microscopy to confirm

the microstructural origins of the observed

differ-ences of behavior as a function of dpa rate

More recently, Chatani and coworkers showedthat at relatively low irradiation temperatures char-acteristic of boiling water reactors, the radiation-induced increments in strength of 304 stainless steelincreased by the 1/4 power of the increase in dparate.87It was demonstrated that the black-spot micro-structure dominated the strengthening It was alsoshown that the concentration of black spots variedwith the square root of the flux as expected, and it isknown that hardening varies with the square root ofthe loop density, thereby producing a fourth-rootdependence Thus, in the absence of any significantmicrochemical or phase stability contributions, it

1200

Test temperature = Irradiation temperature

40 20 10 5

2 1 0.5

Figure 23 Neutron-induced changes in tensile properties of annealed 1.4988 stainless steel irradiated in the DFR fast reactor Reproduced from Ehrlich, K J Nucl Mater 1985, 133–134, 119–126 Ductility declines as strength increases.

0 10 20 30 40 50

20% CW 316 25% CW PCA

appears that radiation-induced strengthening is

affected by dpa rate but not very strongly

The loss of ductility proceeds in several stages,

first involving convergence of the yield and ultimate

strengths as shown inFigures 29 and 30, such that

a loss of work-hardening occurs and very little

uniform elongation is attained As the irradiation

pro-ceeds, there is a progressive tendency toward flow

localization followed by necking As seen inFigure 31

the failure surface shows this evolution with

increas-ing dose

The flat faces observed at highest exposure in

Figure 31are often referred to as ‘channel fracture’

but they are not cleavage faces They are the result

of intense flow localization, resulting from the first

moving dislocations clearing a path of

radiation-produced obstacles, especially Frank loops, and thereby

softening the alloy along that path It is not possible to

remove the voids by channeling but the distorted

voids provide a microstructural record of the flowlocalization as shown in Figure 32 Linkage of theelongated voids is thought to contribute to the failure.Such a failure surface might best be characterized

as ‘quasi-embrittlement’, which is a suppression ofuniform deformation, differentiating it from trueembrittlement, which involves the complete sup-pression of the steel’s ability for plastic deformation.This distinction is made because under some con-ditions quasi-embrittlement can evolve into trueembrittlement

The tendency toward quasi-embrittlement growswith increasing swelling but the alloy is actuallysoftening with increasing swelling rather than hard-ening As shown in Figure 33 brittle fracture(defined as strength reduction with zero plasticity)

of a Fe–18Cr–10Ni–Ti stainless wrapper in BOR-60

at 72 dpa maximum was observed at positions wherepeak swelling occurs.88Some decrease of strength is

~0.5 and ~15 appm He per dpa

on Materials; ASTM STP 1175; 1992, pp 921–939.

observed with increasing irradiation temperature, but

the primary strength reduction for specimens tested at

the irradiation temperature arises from the magnitude

of swelling Testing at temperatures below the

irradia-tion temperature (e.g., 20C) demonstrates the same

dependence on swelling and irradiation temperature,

but the strength and plasticity values are higher As

expected, the strengths for tests conducted at 800C

are uniformly much lower than that observed at lower

temperatures, but there is an absence of any

relation-ship between strength and swelling at this temperature

As shown in Figure 34 failure surfaces at high

swelling levels exhibit transgranular cup-cone

mor-phology where failure proceeded by micropore

coa-lescence arising from stress concentration between

deforming voids.88Similar fracture morphology has

been observed in studies on other stainless steels.1

Although voids and bubbles initially serve toharden the microstructure,78 large swelling levelsallow previously second-order void effects to becomedominant.1,88,89One of these second-order effects isthe strong decrease of elastic moduli at high swellinglevels All three elastic moduli decrease initially

at
2% per each percent of void swelling.90–93

At>10% swelling this leads to significant reduction

in strength

As a consequence, the slope of the elastic region(Young’s modulus) of the stress–strain curve decreases,and more importantly, the barrier strengths of all sinksdecrease as the shear modulus likewise decreases.Therefore, the yield and ultimate strengths decreasewith increasing swelling, even though the elongationstrongly decreases Similar behavior has also beenobserved in pure copper.94

Original series

Isothermal repeat series

Isothermal repeat series

Without

Figure 26 Comparison of isothermal and nonisothermal behavior on convergence behavior The original target

temperature of 495C was maintained for some time but thereafter there was a large, relatively brief over-temperature event, followed by a prolonged and significant under-temperature event Reproduced from Garner, F A.; Hamilton, M L.; Greenwood, L R.; Stubbins, J F.; Oliver, B M In Proceedings of 16th ASTM International Symposium on Effects of Radiation

on Materials; ASTM STP 1175; 1992, pp 921–939 When the target temperature was reestablished in the second and third irradiation segments the mechanical properties returned to the isothermal destination.

The nature of the void-related failure changes from

quasi-embrittlement to true embrittlement for tests at

or near room temperature, demonstrating another

example of a late-term second-order process growing

to first-order importance at higher swelling levels

Hamilton and coworkers observed that above
10%swelling the previously established saturation strengthlevel of 316 stainless steel suddenly increased verystrongly in room temperature tensile tests.95Similarresults were observed in Russian steels.96,97As shown

inFigures 35 and 36 the failure surfaces in such testshad rotated from the expected 45 (relative to thestress axis) to 90 as swelling approached 10%, indi-cating complete brittle failure, as also indicated by thefully transgranular nature of the failure surface Con-currently, the ductility vanished and the tearing mod-ulus plunged to zero, indicating no resistance to crackpropagation Once a crack has initiated it then propa-gates completely and instantly through the specimen.Neustroev and coworkers observed such failures inRussian steels that are subject to greater amounts ofprecipitation and determined that the critical micro-structural condition was not defined solely by the level

of swelling, but by the obstacle-to-obstacle distance ofthe void-precipitate ensemble, indicating that stressconcentration between obstacles was one contributingfactor.96 However, it was the progressive segregation

of nickel to increasing amounts of void surface andthe concurrent rejection of chromium from the sur-faces that precipitated the rather abrupt change infailure behavior.1,95This late-term void-induced micro-chemical evolution induces a martensite instability inthe matrix, as evidenced by the failure surface beingcompletely coated with alpha-martensite.95

MFE-4 experiment in ORR

9.4

Ni 7

15 20 Cr

2.9 2.4

2.3 2.1

Fluence (dpa)

Phénix Rapsodie

Figure 28 Differences in strength change exhibited by

annealed 316 stainless steel after irradiation at 390C in

the PHENIX and RAPSODIE fast reactors Dupouy, J M.;

Erler, J.; Huillery, R In Proceedings International

Conference on Radiation Effects in Breeder Reactor

Structural Materials, Scottsdale; The Metallurgical Society

of AIME: New York, 1977, pp 83–93 Phe´nix operated at a

displacement rate that was
three times higher than that of

RAPSODIE.

The abrupt jump in strength just before failure

observed by Hamilton and coworkers is the result

of a stress-induced blossoming of a high density

of small, thin, epsilon-martensite platelets, as seen

inFigure 37 These platelets are essentially stacking

faults that form under stress as a result of the

influ-ence of both falling nickel level and low

deforma-tion temperature to decrease the stacking fault

energy of the matrix.1 When encountered by the

advancing crack tip, the epsilon-martensite is

con-verted to alpha-martensite in the strain field ahead of

the crack, providing a very brittle path for further

cracking

The correlation between void swelling and both

quasi-embrittlement and true embrittlement is

observed not only in slow tensile tests (Figures 36,

38, and 39) but also in Charpy impact tests as shown

in Figure 39 Figures 40–44 present examples ofswelling-induced failures in components experien-cing a wide range of physical insults The example

of Porollo et al in Figure 44 (top) is particularlynoteworthy in that it results from significant swelling

at
335C, a temperature earlier thought not toproduce significant amounts of swelling

If there are no physical insults experienced bythe component during irradiation, the continuedsegregation of nickel to void surfaces and the con-current rejection of chromium can lead to strongchanges in composition in the matrix during irradia-tion, pushing the matrix toward ferrite rather thanmartensite at higher temperatures, especially for steelswith nickel content of <10% In some observationsvoids encased in austenite shells have been observed

to exist in a pure ferrite matrix.98,99To date, however,

no significant component failure has been reported

to result from this particular late-term instability.Finally, there appears to be another late-termphase instability developing at lower irradiation tem-peratures that involves martensite but does notappear to be due to void swelling Gusev et al haveshown that for irradiation temperatures below

350C a growing tendency for stress-induced tensite formation is occurring in Russian austeniticsteels at doses in the range of 25–55 dpa when tested

mar-at room tempermar-ature.100–102 Surprisingly, this bility results in a restoration of engineering ductility

insta-to preirradiation levels However, the ductility is

392 376 373 371

od f,dpa s –1

399 378 374 372

Figure 29 Strength changes observed in annealed 304

and 316 stainless steels irradiated in EBR-II at 371–426C

and tested at 385C Reproduced from Brager, H R.

Blackburn, L D.; Greenslade, D L J Nucl Mater 1984,

122–123, 332–337 Microscopy showed that the

dependence of microstructure on displacement rate was

consistent with the macroscopic behavior exhibited by each

alloy In AISI 316, the flux dependence of precipitation

canceled the opposite dependence of other microstructural

Figure 30 Convergence of ultimate and yield strengths of annealed 304 stainless steel irradiated in EBR-II and tested

at 370C Reproduced from Holmes, J J.; Straalsund, J L.

In Proceedings of International Conference: Radiation Effects in Breeder Reactor Structural Materials; 1977;

pp 53–63.

regained not because the steel has softened, but

because it becomes exceptionally strong and

hard-ened during deformation As a consequence, the steel

has lost the ability to neck

Uniform elongation

Proportional elastic limit

Proportional elastic limit

Unirradiated plastic dimpling

EBR-II 304 SS Irradiated at 700 F Tested at 700 F

Figure 31 Increase in strength, loss of ductility, and change in failure mode observed during tensile testing in annealed

304 safety and control rod thimbles (SRT and CRT) after irradiation at
370 C in EBR-II Reproduced from Fish, R L.;

Straalsund, J L.; Hunter, C W.; Holmes, J J In Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys; ASTM STP 529; 1973; pp 149–164.

Figure 32 Intense flow localization manifested as

shearing of voids below a ‘channeled’ failure surface in a

304 steel tensile specimen at 40 dpa and
400 C when

tested at 370C There is 100–200% strain in the 0.05 mm

wide deformation band Reproduced from Fish, R L.;

Straalsund, J L.; Hunter, C W.; Holmes, J J In Effects of

Radiation on Substructure and Mechanical Properties of

Metals and Alloys; ASTM STP 529; 1973; pp 149–164.

The swelling was
5% in this specimen.

21, 17,

15, 25, 23,

Instead of necking, a moving wave of deformation

is initiated at the first attempted necking point Thewave front then travels nearly the full length ofthe gage section Initially, there is a local deforma-tion in the order of 40–45%, but as the wave movesforward it leaves a relatively uniform local defor-mation in its wake Everywhere behind the wavefront there is measured 30–35% volume percent ofmartensite, as shown in Figures 45 and 46 Themartensite is not only a product of the wave, butalso the cause of the wave Deformation-inducedmartensite resists further necking and forces thedeformation to be displaced to the adjacent lesserdeformed material The mechanisms that cause thelate-term onset of martensite instability have not yetbeen determined

A property of important engineering interest is thefracture toughness Jc While the fracture toughness

of various unirradiated stainless steels can be quite

10mm

Figure 34 Fracture surface of Fe–18Cr–10Ni–Ti stainless

steel specimen at a swelling level of 26% Reproduced from

Neustroev, V S.; Garner, F A J Nucl Mater 2009,

386–388, 157–160 Micrograph corresponds to open circle

different, it appears that all austenitic steels studied

undergo the same general evolution in toughness

during irradiation Mills has shown that three regimes

of evolution occur.103,104 The first regime involves

a low-dose threshold exposure range (<1 dpa) where

there is essentially no loss of toughness, and the

second regime involves an intermediate exposure

range (1–10 dpa) where toughness decreases rapidly

with exposure, producing an order of magnitude

reduction in Jcand two orders of magnitude

degra-dation in tearing modulus Finally, a saturation regime

is reached, in which increasing exposure does not

produce a further reduction in toughness This

saturation occurs well before any of the void-induced

instabilities discussed above can occur As shown in

Figure 47, the saturation level is remarkably pendent of the original toughness level

inde-Welds in austenitic alloys were shown by Mills toexhibit lower initial toughness values and lower satu-ration toughness levels as well The fracture toughnesslevel is sensitive to the test temperature, however, asshown in Figure 48 At high test temperatures,the fracture mode changes from transgranular to inter-granular in nature, reflecting the effect of test tempera-ture on both matrix strength and also the influence ofhelium embrittlement at grain boundaries.105The level

of helium needed to promote high temperature tlement is not very high, however, and can easily bereached after moderate neutron exposure in fastreactor-irradiated alloys with the lowest nickel level

Test at 20 C Test at irradiation temperature

60

30

30 20

Swelling (%) 10

0

Test at 20 C Test at 20 C

Test at irradiation temperature Test at

irradiation temperature

Figure 36 Influence of swelling on fracture properties during tensile testing of an annealed Fe-18Cr-10Ni-Ti steel irradiated in BOR-60 at 400–500C Neustroev, V S.; Shamardin, V K Atomnaya Energiya 1990, 71(4), 345–348,

in Russian Note that softening and rotation of fracture surface by voids is observed at both room and elevated

temperatures.

4.02.8 Radiation-Induced Changes

in DimensionOne of the most challenging engineering conse-quences of neutron irradiation is the development

of dimensional instability, whereby a structural ponent can shrink or grow in volume and where itcan be distorted in shape, often with both processesoccurring at the same time There are two majorcategories of such changes: conservative of volumeand nonconservative of volume A distinction canalso be made between processes that distributethe resulting strains isotropically or anisotropically.Additionally, a further distinction can be made con-cerning whether the process to the first-order isstress-driven or not, or whether it is stress-sensitive

com-to the second-order

Depending on the crystal structure there are

a variety of such distortion processes, some moreprominent than others in a given crystal system.For austenitic stainless steels the phenomenon ofradiation-induced growth (volume-conservative, an-isotropic distribution of strains in the absence of stress)

is not an issue, whereas for hexagonal close packedalloys based on zirconium and rhenium growth isoften a dominant process.9,106 Austenitic steels also

In Effects of Radiation on Materials: 13th International Symposium (Part II) Influence of Radiation on Material Properties; ASTM STP 956; 1987; pp 245–270.

Figure 38 Influence of swelling on ultimate tensile strength

of 0Khl6N15M3B cladding and Kh18N10T hexagonal ducts

irradiated in BOR-60 Tests were performed at the irradiation

temperature, using specimens cut from regions of maximum

swelling Reproduced from Neustroev, V S.; Shamardin, V K.

Phys Met Metallogr 1997, 83(5), 555–560 The two

steels develop loss of strength with swelling differently,

probably reflecting the very different precipitate

structures of the two steels.

are not very prone to significant

transmutation-induced changes in lattice parameter as sometimes

observed in alloys based on rhenium and

vana-dium.106,107See alsoChapter4.01, Radiation Effects

in Zirconium Alloys

Stainless steels experience three general

cate-gories of radiation-induced strain processes These

are precipitation-related strains, void swelling, and

irradiation creep In general, these three processesare not fully independent but are interrelated andoften synergistic

Stainless steels undergo an evolution of phase ture at reactor-relevant temperatures, even in theabsence of radiation These changes involve the for-mation of various carbides, later followed by variousintermetallic phases.1,108This evolution is accompa-nied by net changes in average lattice parameterarising from differences in partial molar volume ofelements when passing from one phase to another

Uniform elongation

Reproduced from Fissolo, A.; Cauvin, R.; Hugot, J P Levy, V In Effects of Radiation on Materials: 14th International Symposium; STP 1046; 1990; Vol II; pp 700–713.

Figure 40 Failure during mounting in a vise of severely

void-embrittled 316 stainless steel creep tube irradiated in

the EBR-II fast reactor to 130 dpa at
400 C with a hoop

stress of 276 MPa Reproduced from Porter, D L.;

Garner, F A J Nucl Mater 1988, 159, 114–121 Swelling

at the initial failure point was
14%.

Figure 41 Void-induced embrittlement of an annealed

304 steel EBR-II assembly duct after
54 dpa

at
400 C Reproduced from Flinn, J E.; Krajcinovic, D.;

Phipps, R D.; Franklin, D G.; Miller, S C Evaluation of Ex-Reactor Loading Event on High-fluence EBR-II Control-rod Thimble 5E3, ANL/EBR-068, February 1973 The duct broke during routine handling in the hot cell.

The resulting macroscopic strains are sometimes

very counterintuitive, however, especially with

re-spect to their sign

For example, formation of the less dense carbide

phases leads to macroscopic densification of the alloy

and shrinkage of volume,109 while the formation of

denser intermetallic phases (Chi, Sigma, Laves)

usu-ally leads to an increase in volume, a form of nonvoid

swelling.110,111This counterintuitive behavior is the

result of the different partial molar volumes of

criti-cal elements (C and Mo primarily) between the new

precipitates and the alloy matrix in which they form.Both the carbide and intermetallic phase evolutionappear to be accelerated and sometimes alteredunder irradiation

Other radiation-produced phases (w0, G-phase)also appear to induce changes in lattice parameter

Figure 42 Severe embrittlement and failure in three

BOR-60 reflector assembly ducts The ducts were made of

annealed X18H10T, the Russian equivalent of 321 steel.

Reproduced from Neustroev, V S.; Ostrovsky, Z E.;

Teykovtsev, A A.; Shamardin, V K.; Yakolev, V V In

Proceedings of 6th Russian Conference on Reactor

Materials Science; 11–15 September 2000, Dimitrovgrad,

Russia, in Russian The maximum swelling values (from left

to right) were 27.8, 29.8, and 14% Failure was the result of

high withdrawal loads arising from both swelling and

bending, the latter a consequence of radial dpa gradients in

the reflector.

Figure 43 Failure of 20% cold-worked D9 (Ti-modified

316) cladding during routine handling Failure occurred

where 90 dpa was attained at
460 C in FFTF,

producing
32% swelling Reproduced from Makenas,

B J.; Chastain, S A.; Gneiting, B C , ‘‘Dimensional

Changes in FFTF Austenitic Cladding and Ducts,

Westinghouse Hanford Company Report WHC-SA-0933VA,

Richland WA, 1990 Fuel was lost from the open section.

Yu V.; Dvoraishin, A M.; Krigan, V M.; Budylkin, N I.; Mironova, E G.; Garner, F.A J Nucl Mater 1998, 258–263, 1613–1617 All tubes lost pressure, either by cracking or by completely failing during removal from their canister Before breaking, the tubes were also bent by irradiation creep due

to swelling-induced interaction with the top of the canister Swelling of 6.2% was measured in the zero stress, annealed tube and 11.2% in the cold-worked zero stress tube.

F A J Nucl Mater 2010, 403, 121–125.