Effect of Post-Irradiation Annealing on Defects in ... · - i - Foreword When I accepted this...

Thesis Title

Effect of Post-Irradiation Annealing on Defects in

Irradiated Model Alloys by Means of Positron

Annihilation Coincidence Doppler Broadening

Spectroscopy

Ir Marc Deprez Academic year 2010-2011

Promoter : prof dr Eric van Walle, KULeuven

Assessors : prof dr Jacqueline Lecomte-Beckers, ULg

prof dr Luc Dupré, UGent

Mentor : dr ir Marlies Lambrecht, SCK•CEN

Thesis submitted in partial fulfilment of the requirements for the degree of Master of Science in Nuclear Engineering

‘…we must regard it rather as an accident that the Earth contains a

preponderance of negative electrons and positive protons. It is quite possible

that for some of the stars it is the other way about, these stars being built up

mainly of positrons and negative protons. In fact, there may be half the stars of

each kind.’

Paul A. M. Dirac, out of his Nobel Prize Lecture ‘Theory of Electrons and Positrons’, Dec 12, 1933.

All property right and copyright are reserved. Any communication or reproduction of this document

and any communication or use of its contents without explicit authorization is prohibited. Any

infringement to this rule is illegal and entitles to claim damages from the infringer, without prejudice to

any other right in case of granting a patent of registration in the field of intellectual property.

© BNEN, Belgian Nuclear Education Network c/o SCK•CEN, Boerentang 200, BE-2400 Mol, Belgium

- i -

Foreword

When I accepted this subject as thesis, I had only vague ideas of positron annihilation

techniques and thermal diffusion in alloys. At the end of my experiments I was surprised to

find out that steel is actually a very lively matter. As an engineer you easily accept what you

see, try to build an empirical formula around it and make direct use of it. The scientific

approach however is much more demanding. Trying to explain the different phenomena

requires a good knowledge/experience background and sincere analytical thinking to finally

come to conclusions backed up by sufficient evidence. Long internal dialogues questioning

the different possibilities, intermingled with uncertainty and enlightened by discussions with

colleagues and experts are the part of every true scientist. I hope I can carry further this

honest and humble approach of scientific research into my further work.

I would like to thank my mentor, Marlies Lambrecht, for guiding me through the whole

process of the thesis : explaining the equipment and instruments, indicating how best to

perform the experiments, showing how to process the data, learning how to write a scientific

work. I thank her specially for her patience with me and her readiness to help me whenever

she could. Without her extensive experience in performing and interpreting the results of

positron annihilation experiments on irradiated steels, this thesis would have been

impossible.

I like to thank my promoter, Prof. Eric van Walle, for taking time to discuss the subject

with us, for some redirections in the approach and for giving me the opportunity to visit his

colleague Prof. Yasuyoshi Nagai of the Tohoku University at the Oarai Nuclear Research and

Development Center (ONRDC) in Japan.

I like to thank Prof. Yasuyoshi Nagai and Assistant Prof. Toyama Takeshi for their warm

welcome at the ONRDC, for the interesting discussions on the subject and the friendly help

with the interpretation of certain aspects of my results. Also many thanks go to the group of

PhD-students at the ONRDC and secretary Marie Okuno who took care of me when I was in

Japan.

Foreword

- ii -

A special thanks goes to Lorenzo Malerba, expert in matrix damage modeling at

SCK•CEN, who explained me well what keeps the different atoms in an alloy together and

under what conditions and how precisely these atoms move through the matrix.

I like to thank Milan Konstantinovic for this advice and help on annealing the samples

and Boris Minov for the good coordination concerning internal friction (IF) measurements

that were performed on an IF sample simultaneously annealed. I like to thank the

collaborators of the controlled zone of the LHMA1, Dirk Quirijnen, Henri Maussen, Eddy

Kox, Danny Penneman, Patrick Ouderits, for their professional and kind help and Ms. Cindy

Verachtert of the HLMA for the administrative help.

Finally, I would like to thank my friends and family, who preferred to stay anonymous.

Their love and support helped me to finish this work, although I probably failed in explaining

what it was about.

Marc Deprez

December 2010

1 Laboratory for High and Medium Activity

- iii -

Summary

Reactor pressure vessels (RPVs) are the most important components of nuclear power

plants (NPPs) as they contain practically the full inventory of radioactive material. They are

for practical and economic reasons hardly replaceable and their lifetime therefore also

determines the economic lifetime of the NPP. The lifetime of a RPV depends on its resistance

against brittle fracture or fracture toughness, which in its turn depends on the neutron

irradiation induced embrittlement of the RPV steel.

This RPV steel embrittlement is conventionally attributed to three factors : matrix

damage (vacancies and interstitial atoms), copper-rich precipitates (CRP) and phosphorus

grain boundary segregation (GBS). For Western RPVs, only the two first factors are of

importance. A way to remediate the embrittlement could be by thermal annealing of the RPV.

There is substantial practical experience on thermal annealing of VVER-4402 RPVs, albeit of

a limited circumferential surface. There is however very little material test experience on the

effect of thermal annealing, e.g. the (ductility) recovery efficiency or reembrittlement trend.

The phenomena of ductility recovery by thermal annealing can be studied by isochronal

heating of irradiated RPV steels. A start was given in this thesis. Two irradiated binary model

alloys Fe-0.3wt%Cu ; 0.1 dpa and Fe-0.3wt%Cu ; 0.2 dpa were isochronally annealed over a

temperature trajectory from 300°C to 700°C in steps of 50°C and subsequently analyzed by

means of coincidence Doppler broadening (positron) annihilation radiation (CDBAR)

spectroscopy. It is known that positrons are attracted, trapped and annihilated at both vacancy

and nano-size copper precipitate sites. Therefore, CDBAR spectroscopy is an interesting

technique for studying the evolution of matrix damage and copper precipitates with

temperature.

The limited, but lengthy experiments showed consistent results. The main finding is the

confirmation of two clear recovery stages during the annealing process, each with its rather

well defined temperature. At around 450°C, the vacancies which are believed -if surviving-

to be trapped in the copper precipitates, dissociate from these precipitates to diffuse in the

bulk and finally disappear at sinks. At around 650°C, the copper precipitates dissolve and

2 Russian type of pressurized water reactor (PWR) of 440 MWe

Summary

- iv -

diffuse into the bulk where they probably homogeneously disperse. The irradiation dose did

not seem to have a particular influence on the phenomena apart from the obvious larger

initial amount of vacancies in the second sample.

The optimum annealing temperature has always been a point of discussion. A too low

annealing temperature would lead to insufficient recovery efficiency ; a too high annealing

temperature would or could lead to creep deformation, temper embrittlement and thermal

stresses and related dimensional deformations in the RPV wall and attached piping. A study

at ORNL3 on a particular irradiated welded material showed already very good recovery

(about 80%) at 454°C, which probably corresponds to our phenomenon of vacancies

dissolving from the copper precipitates. This would mean that there is no need for dissolving

the copper precipitates (at a much higher temperature), taking also into account that the

irradiation-induced copper precipitation process is a rather fast process and that the copper

precipitates would thus be quickly reestablished during reirradiation.

3 Oak Ridge National Laboratory in the US

- v -

Samenvatting

Nucleaire reactors vormen de belangrijkste componenten van kerncentrales aangezien zij

practisch de volledige hoeveelheid aan radioactief materiaal bevatten. Om redenen van

practische en economische aard zijn ze quasi onvervangbaar wat met zich meebrengt dat de

levensduur van de reactor meteen ook de economische levensduur van de kerncentrale

bepaalt. De levensduur van een nucleaire reactor wordt bepaald door zijn weerstand tegen

brosse breuk of breuktaaiheid, welke op zijn beurt afhangt van de verbrossing van het

kuipstaal door neutronenbestraling.

Kuipstaalverbrossing wordt conventioneel toegeschreven aan drie factoren : matrix-

schade (vacatures en interstitiële atomen), koperverrijkte precipitaten en fosfor korrelgrens

segregatie. Voor westerse reactors zijn enkel de eerste twee factoren van belang. Een manier

om de verbrossing te herstellen kan bestaan uit het uitgloeien van de reactorkuip. Hierover is

reeds een belangrijke practische ervaring voorhanden via het uitgloeien van een aantal

VVER-4404 reactoren, hoewel het hier wel het uitgloeien van slechts een beperkte

cylindrische strook van de reactoren betrof. Er is echter maar een beperkte hoeveelheid test

data beschikbaar aangaande het effect van het uitgloeien van neutronen bestraalde kuipstalen,

zoals bijvoorbeeld aangaande het uitgloei-rendement (herstellen van de ductiliteit) en de

snelheid van verbrossing bij het opnieuw blootstellen aan neutronenbestraling na de

uitgloeibehandeling.

De fenomenen die gepaard gaan met het herwinnen aan ductiliteit door uitgloeien

kunnen bestudeerd worden door het isochroon uitgloeien van neutronen bestraalde kuipstaal-

monsters. Een eerste aanzet hiertoe werd gegeven in deze thesis. Twee bestraalde binaire

modellegeringen Fe-0.3wt%Cu ; 0.1 dpa en Fe-0.3wt%Cu ; 0.2 dpa werden isochroon

gegloeid over een temperatuurstraject van 300°C tot 700°C en dit in stappen van 50°C en

telkens geanalyseerd door middel van coincidence Doppler broadening (positron)

annihilation radiation (CDBAR) spectroscopie. Het is een gekend feit dat positronen

aangetrokken en ‘gevangen’ worden door en annihileren in zowel vacatures als nano-grootte

koperverrijkte precipitaten. Om deze reden is CDBAR spectroscopie een interessante

4 Russisch type drukwater reactor (PWR) met een vermogen van 440 MWe

Samenvatting

- vi -

techniek voor het bestuderen van de evolutie van zowel matrixschade als koperverrijkte

precipitaten met de temperatuur.

De experimenten (beperkt wegens tamelijk tijdrovend) geven een consistent beeld als

resultaat. De belangrijkste bevinding is de bevestiging van twee duidelijk op te merken

ductiliteitsherstelniveau’s gedurende het uitgloeiproces, elk met zijn eigen relatief goed

gedefinieerde temperatuur. Rond 450°C maken de vacatures, waarvan wordt verondersteld

dat ze -indien nog aanwezig- gevangen zitten in de koper precipitaten, zich los van deze

precipitaten en verspreiden zich via thermische diffusie om tenslotte te verdwijnen in sinks.

Rond 650°C lossen de koper precipitaten op en verspreiden zich homogeneen over het

materiaal. De bestralingsdosis bleek daarbij geen specifieke invloed te hebben, behalve dan

de initiële, hogere densiteit aan vacatures in de tweede modellegering.

De optimale uitgloeitemperatuur heeft altijd ter discussie gestaan. Een te lage

uitgloeitemperatuur leidt tot onvoldoende herstel van de ductiliteit ; een te hoge uitgloei-

temperatuur kan leiden tot kruipvervorming, thermische verbrossing (door wijziging van de

korrelstructuur) en tot dimensionele vervormingen via thermische spanningen in de reactor

kuipwand en eraan bevestigde pijpleidingen. Een studie aan het ORNL5 betreffende een

specifieke, neutronenbestraalde lasverbinding gaf aan dat reeds een zeer goede ductiliteits-

herstel (rond de 80%) kan bekomen worden op 454°C, wat wellicht overeenkomt met het

hierboven beschreven fenomeen van vacatures die zich losmaken van de koper precipitaten.

Dit zou meteen ook betekenen, gezien dit resultaat, dat het weinig waardevol is te pogen de

koper precipitaten op te lossen (op een hogere temperatuur), temeer daar het proces van

stralings-geïnduceerde koper precipitatie een relatief snel proces is en derhalve de koper

precipitaten zich snel zullen herstellen bij hernieuwde blootstelling aan neutronen-bestraling.

5 Oak Ridge National Laboratory in de USA

- vii -

List of symbols

List of chemical symbols

Al Aluminum

Ar Argon

C Carbon

Co Cobalt

Cr Chromium

Cu Copper

Fe Iron

Mn Manganese

Mo Molybdeen

Na Sodium

Ne Neon

Ni Nickel

P Phosphorus

Si Silicon

List of symbols

- viii -

List of physical symbols

A Area

A+ Positron affinity

bara bar absolute

c Speed of light, c = 2.998 x 108 m/s

E° Standard (electrode) potential

Eb Binding energy

eV Electron volt, 1 eV = 1.602 x 10-19

J

h Planck’s constant, h = 6.626 x 10-34

J∙s

m0 Electron (positron) rest mass, m0 = 0.511 MeV/c² = 5.49 x 10-4

u

mrad milliradian or 10-3

radian, 1 rad = 57.2958°

n Neutron

ν Velocity of the center of mass

p Linear momentum

pL Longitudinal component of the linear momentum

σ Statistical dispersion

V Volt

- ix -

Acronyms and abbreviations

3D-AP 3-dimensional atom probe

ACAR Angular correlation (positron) annihilation radiation

AP Atom probe

APT Atom probe tomography

ASME American Society of Mechanical Engineers

ASTM American Society for Testing and Materials

bcc Body-centered cubic

BWR Boiling water reactor

CDB Coincidence Doppler broadening

CDBAR Coincidence Doppler broadening (positron) annihilation radiation

CPR Copper-rich precipitates

CVN Charpy V-notch

DBTT Ductile-to-brittle transition temperature

DDA Digital data acquisition (card)

dpa Displacements per atom

DSP Digital signal processor

ECCS Emergency core cooling system

EDF Électricité de France

fcc Face-centered cubic

FWHM Full width at half maximum

GBS Grain boundary segregation

HAZ Heat affected zone

HCFB Hot cell feed box

HPGe High purity germanium

HRTEM High resolution transmission electron microscope

ICP-MS Induced coupled plasma mass spectroscopy


- x -

IGS Intergranular segregation

INT Interface (card)

JIS Japanese industrial standards

IF Internal friction

JRC Joint Research Centre

KIa Stress-intensity factor for crack-arrest fracture toughness

KIc Stress-intensity factor for static fracture toughness

KId Stress-intensity factor for dynamic fracture toughness

LHMA Laboratory for High and Medium Activity

NPP Nuclear power plant

ONRDC Oarai Nuclear Research and Development Center

ORNL Oak Ridge National Laboratory

PALS Positron annihilation lifetime spectroscopy

PAS Positron annihilation spectroscopy

PC Personal computer

PCI Peripheral component interface

PED Pressure equipment directive

PET Positron emission tomography

PIA Post-irradiation annealing

PLC Programmable logic controller

PTS Pressurized thermal shock

PWR Pressurized water reactor

QED Quantum electrodynamics

RPV Reactor pressure vessel

RT Room temperature

RTNDT Reference temperature for nil ductility transition

S S or shape parameter of a CDB spectrum

SANS Small angle neutron scattering

SIA Self-interstitial atom


- xi -

TM Tele-manipulator (hot cell)

TNDT Temperature for nil ductility transition

T41J Charpy index temperature at 41 J energy

USE Upper shelf fracture energy

USNRC United States Nuclear Regulatory Commission

UTS Ultimate tensile strength

VVER-440 Russian type PWR of 440 MWe

W W or wing parameter of a CDB spectrum

wt% Mass percentage

Z Atomic number

- xii -

Table of content

Foreword i

Summary iii

Samenvatting v

List of symbols vii

Acronyms and abbreviations ix

Chapter 1 : Introduction 1

1.1. The nuclear power industry 1

1.2. The reactor pressure vessel 1

1.3. Embrittlement of RPV steel 3

1.4. Thermal annealing 4

1.5. Experimental thesis work 5

Chapter 2 : Framework 6

2.1. Hardening and embrittlement 6

2.2. Causes and mechanisms of embrittlement and recovery 11

2.2.1. Matrix damage 13

2.2.2. Copper precipitation 13

2.2.3. Recovery of ductility 15

2.3. Thermal annealing as remediation for embrittlement 17

2.3.1. Practical experience 17

2.3.2. Annealing equipment 18

2.3.3. Recovery and reembrittlement : parameters 22

2.3.4. Recovery and reembrittlement : general considerations 25

2.4. CDBAR as technique for nanoscale material studies 28

2.4.1. The positron 29

2.4.2. Positron thermalization 30

2.4.3. Positron trapping 31

2.4.4. Positron annihilation 32

2.4.5. PALS 33

2.4.6. ACAR spectroscopy 35

2.4.7. CDBAR spectroscopy 36

Chapter 3 : Experiments 43

3.1. Binary alloy samples investigated 43

3.2. Annealing setup and work method 44

3.3. CDBAR setup and work method 46

3.4. Experimental results 51

3.5. Interpretation of the results 60

Chapter 4 : Conclusions 63

Table of content

- xiii -

Appendix A 65

Appendix B 66

Appendix C 67

Appendix D 68

References 70

p. 1

Chapter 1 : Introduction

1.1. The nuclear power industry

Since the demonstration of a sustained chain fission reactor in 1942, nuclear power has

emerged as a proven technology for the production of electricity throughout the world.

Because of the world’s continuously improving living standards, increased population and

the concern over the increased concentration of greenhouse gas emissions caused by fossil

fuel, there may be an increasing demand for nuclear power in the next decennia.

In 2008, 439 nuclear power plants with a capacity of about 350 GWe supplied 16% of

global electricity. Of these, about 327 nuclear power plants have been in operation for 20

years or more and these older units have proven to be most profitable because of high

reliability and partially or fully amortized capital costs [1]. Apart from the structural integrity

issue of the reactor pressure vessel (RPV) mainly due to irradiation embrittlement, there are

no significant technical, safety or economic reasons not to continue the operation of well

managed nuclear power plants over a longer period and therefore the issues of plant life

management and license extension are receiving increasing attention in many countries [2].

Of the nuclear power plants in operation, the most common type is the pressurized water

reactor (PWR) and the second most common is the boiling water reactor (BWR). Although

BWR pressure vessels are constructed of the same steels as the PWR pressure vessels, they

are larger in diameter which results in a lower irradiation exposure due to the larger water

gap between the vessel inner surface and the reactor core. Therefore, this thesis primarily

concentrates on the mechanisms of embrittlement and its possible remediation for pressure

vessels of the PWR type.

1.2. The reactor pressure vessel

The RPV is the key component in a nuclear power plant and is considered to be

irreplaceable [3]. This means that if its mechanical properties degrade sufficiently, it can be

the life-limiting factor of the nuclear power plant. The RPV houses the reactor core and


p. 2

because of this it has a direct safety significance. In case a leak would develop in the RPV at

or below the level of the reactor core and the coolant flow through the leak would be larger

than the maximum flow capable of being supplied by the emergency core cooling system(s)

(ECCS), then the reactor core could become uncovered and subsequently become damaged.

So clearly it is necessary that the nuclear power plant owner can demonstrate at all times that

disruptive failure of the RPV has a very low probability of occurrence. Failure of the RPV

could occur because of a weakness in its material or construction, or as a result of an internal

or external event outside the design basis of the nuclear power plant. Such events could be a

molten fuel/coolant explosion inside the vessel or a failure of the RPV support system at the

outside of the vessel. Within the reasonable expectation that such events have a low

probability of occurrence, main consideration must be given to the resistance to fracture of

the RPV itself.

The original 40 year term for reactor licenses was not based on technical studies of

material degradation or limitations within nuclear technology, but was based on economic

considerations. The current target for re-licensing of most of the plants in many countries in

Europe, Japan and the USA is plant operation up to 60 years [3]. Careful attention should be

given to review if nevertheless some structures or components were designed on the basis of

an expected 40 years lifetime and thus have to be replaced. Of far larger concern however is

the age related degradation of the mechanical properties of the RPV steel and in particular the

neutron irradiation induced embrittlement of the steel. This could lead to a sudden, fast

(brittle or non-ductile) fracture from a critically-sized crack in the RPV steel. On the one

hand, there is the menace that cracks could grow due to corrosion or fatigue to a critical size.

On the other hand, there is the menace that neutron irradiation could degrade the mechanical

properties of the RPV steel to such an extent that it becomes susceptible to failure. As a

consequence, an RPV structural integrity assessment program (also called surveillance

program) has to be in place during the entire operating life of the nuclear power plant [4].

Fracture toughness is the important material property considered in this assessment program.

Neutron irradiation degrades the mechanical properties of RPV steels and the extent of

the degradation is determined by the type and structure of the steel, and factors such as

neutron fluence, irradiation temperature, neutron flux and chemical composition. The most

sensitive location in the RPV is the region adjacent to the reactor core, termed the beltline


p. 3

region. Welds and their heat affected zones (HAZs) in this beltline region are particularly

important since these regions have a higher probability of having flaws.

1.3. Embrittlement of RPV steel

Embrittlement of RPV steel has been attributed to precipitates and grain boundary

segregation, both from impurity atoms (Cu, P) and from alloying elements (C, Mo, Mn, Ni,

Cr, Si, …), associated with irradiation-induced point defects (vacancies and interstitials) and

their clusters, termed matrix damage [5]. The detailed mechanism of their formation and

possible interaction is however not well understood yet. The embrittlement itself is the result

of the fact that the precipitates and other defects hinder the dislocation motion during

deformation. Embrittlement mechanisms in Western and Russian type RPV steels are

partially different in that the main factor in the former steels appears to be due to copper and

in the latter due to phosphorus [6]. The (impurity) Cu-content of older Western RPV steels

could be quite high (e.g. 0.31 wt% for 73W steel). In addition, copper was brought into the

RPV steel via the welds. Until the early 1970s, copper-coated weld wire was used to improve

the electrical contact in the welding process and to reduce corrosion during storage of the

weld wire. When it was discovered that copper and phosphorus increased sensitivity to

irradiation embrittlement, RPV manufacturers imposed strict limits on the percentage of

copper and phosphorus in the welds as well as in the plates. Recent RPV steels therefore

contain far less copper (e.g. 0.06 wt% for 16MND5 steel) [7].

Improvements that have reduced the problem of RPV embrittlement include the use of

tougher steels with lower impurity content, reduction in the neutron flux impinging on the

vessel wall by using an optimized fuel management scheme (the so-called ‘low-(neutron)

leakage core’) or shielding of the RPV wall with materials such as stainless steel or hafnium,

and elimination of beltline welds. However, embrittlement remains a potential issue for some

older vessels and, if severe, may lead to premature power plant closure or alternatively (and

hardly executed) to replacement of the reactor vessel or to RPV thermal annealing in order to

restore the vessel’s ductility. Thermal annealing can repair damage to the crystal lattice and

contribute to the dissolution of metallic precipitates, but it can however not repair cracks in

the reactor vessel.


p. 4

1.4. Thermal annealing

Post-irradiation heat treatment or thermal annealing of RPVs is not particularly a recent

subject. The first RPV annealings were done in 1967 with the US Army SM-1A reactor and

in 1984 with the Belgian BR-3 reactor [8]. These annealings were performed according to the

wet annealing technique whereby the primary coolant is used as heat transfer medium. This

technique is easy to implement, but unfortunately it can be utilized only with reactors which

have a low service temperature. The RPVs of PWRs are not designed to stand the pressure of

water at higher temperatures and the critical point of water is reached already at 374°C (pcrit =

219 bar). Due to the thus obtained limited recovery, wet annealing is normally not a practical

solution for power reactors. Following the publication in 1983 of the Westinghouse

conceptual procedure for dry thermal annealing of an embrittled reactor vessel, the Russians

(and recently the Czechs) undertook the thermal annealing of several highly irradiated

VVER-440 RPVs [9-10]. These reactor vessels showed a high rate of embrittlement mainly

ascribed to phosphorus segregation in the grain boundaries. The dry annealing technique

normally consists of using electric resistance radiant heaters placed inside the reactor.

Annealing in this way offers a good control of the heated area, as the electric heaters are

constructed in sections, separately powered and controlled. The feasibility of an alternative

method using radiant heat from a gas-gas heat exchanger placed inside the reactor and heated

by gas-fired burners placed outside the containment building was demonstrated on the (non-

activated) reactor of the uncompleted Marble Hill power plant in Indiana, US in 1996 [11].

As of today, 15 RPVs type VVER-440 have been thermally annealed. The VVER experience,

along with the results of relevant laboratory scale research with materials irradiated in test

reactors and with materials from commercial RPV surveillance programs, shows that good

recovery of all of the mechanical properties is observed when the thermal annealing

temperature is about 450°C (or 150°C above the irradiation temperature) and this for about

168 hours (1 week). In addition, the re-embrittlement rates upon subsequent reirradiation

were similar to the embrittlement rates observed prior to thermal annealing. The dominant

factors which influence the degree of recovery are the annealing temperature relative to the

irradiation or service temperature, the time at the annealing temperature, the level of impurity

and alloying elements present, and the type of product and its crystalline structure (plate,

forging, weldment, etc…).


p. 5

1.5. Experimental thesis work

This thesis can be situated in the context of thermal annealing of embrittled RPVs. As

already indicated, the mechanisms underlying the formation of point defect clusters and

solute atom precipitates in irradiated steel are not well understood yet and differently

explained in literature. The same accounts for the recovery of irradiated metal during a

thermal annealing process. This thesis aimed at contributing to a better understanding of

these formation and deformation/recovery mechanisms by study of a model alloy of RPV

steel. The importance of using model alloys lays in the restriction of the number of

influencing parameters during the study of the metal behavior, especially the number and

content (wt%) of impurity or alloying elements. The studied model alloys were a binary alloy

Fe-0.3wt% Cu irradiated at a level of 0.1 dpa and an identical binary alloy irradiated at 0.2

dpa. The alloy samples were subjected to an isochronal annealing process which consisted of

annealing the samples together for 30 minutes at 50°C intervals over a temperature trajectory

from 300°C to 700°C. After each annealing, the samples were analyzed using a positron

annihilation technique called coincidence Doppler broadening annihilation radiation

(CDBAR), explained in chapter 2, to determine the evolution of vacancies and Cu-

precipitates. The results of this study, which can be found in chapter 3, may give some

insight into the recovery process of irradiated RPV steel during thermal annealing and of

what can be expected of in-service annealing of nuclear reactors.

Chapter 2 : Framework

2.1. Hardening and embrittlement

The low-alloy ferritic RPV steels exhibit the classic ductile-to-brittle transition behavior

with decreasing temperature [7]. In the low toughness region, transgranular6 cleavage is the

most occurring failure mode, while ductile rupture (shear fracture) is the most occurring

failure mode in the high toughness region. As temperature increases from the low to high

toughness region, the cleavage fracture possibility decreases and the ductile rupture

possibility increases. Neutron irradiation tends to increase the temperature at which this

transition occurs and tends to decrease the ductile toughness. Neutron irradiation thus shifts

the ductile-to-brittle region towards the operating domain of the ferritic RPV. At all

scenario’s, sufficient margin has to be available between the RPV operating domain and the

ductile-to-brittle transition region in which unstable cleavage fracture is possible.

As the RPV steel thus becomes progressively more brittle with neutron irradiation, the

changes in strength and toughness with temperature must be known so that the RPV can be

safely operated within the (P,T)7 design & operation envelope. These changes in material

properties due to neutron irradiation are monitored by means of surveillance programs. In

most countries the layout of the surveillance programs is standardized either in national

standards or by adopting the American Standard ASTM E185-82 [12]. Surveillance specimen

capsules are located near the inside vessel wall in the beltline region, so that the specimen

radiation history duplicates to the extent practicable the irradiation conditions experienced by

the reactor vessel inner wall. The ASTM standard even recommends that the surveillance

capsules lead factors8 be in the range of one to three, so that to a certain extent predictions

can be made concerning the material property changes. The surveillance capsules are

6 Although transgranular cleavage is the predominant mode of brittle fracture in RPV steels, solute (e.g. phos-

phorus) segregation to grain boundaries can result in another type of brittle fracture known as intergranular (grain boundary) fracture. 7 (pressure, temperature)

8 The lead factor is the ratio of neutron exposure of the surveillance specimens to the highest anticipated ex-

posure at the RPV wall. This lead factor can be increased by placing the surveillance capsules somewhat nearer to the center of the reactor core (normally attached to the thermal shield)


p. 7

withdrawn according to a certain time schedule and, in general, the specimens are subjected

to tensile strength tests and Charpy V-notch (CVN) tests.

The effects of neutron irradiation on the material properties of ferritic steel can be

summarized as follows :

.an increase in the CVN ductile-to-brittle transition temperature (DBTT)

.a drop in the CVN upper shelf fracture energy (USE)

.an increase in the yield strength and the ultimate tensile strength (UTS)

Embrittlement in terms of the CVN test is qualitatively described by 2 parameters :

.the 41J temperature9 (T41J) is typically used to define the DBTT change

.the USE is typically used as a measure of the ductile fracture toughness at

higher temperatures

Fig 1 - Temperature shift in the CVN curve due to neutron irradiation [13]

9 US literature uses T30 instead of T41J referring to 30ft-lb which is equivalent with 41J


p. 8

The assessment of the RPV structural integrity is a quantitative study and is based on

material fracture toughness. The key input for the assessment is the variation of the material

fracture toughness with the temperature. Although fracture toughness is a material property

and can be measured via a standardized test10

, it is only recently that fracture toughness tests

are performed on RPV steels. One of the reasons was that the larger fracture mechanics

specimens were originally not foreseen in the surveillance capsules. The RPV integrity

assessment was then performed according to fracture toughness lower bound curves, indexed

by a reference temperature and this according to a methodology outlined in ASME codes11

,

ASTM E185 and USNRC Regulatory Guide 1.99. This methodology uses data of CVN tests

and not direct measurements of fracture toughness12

.

The ASME Code KIc curve, or reference static fracture toughness lower bound curve

provides the toughness characterization for the static crack initiation and is given by [14]

KIcASME

= 36.48 + 22.78 exp [0.036 (T - RTNDT)] (2.1)

The ASME Code KId curve, or reference dynamic fracture toughness lower bound curve

provides the toughness characterization for the dynamic crack initiation and is given by [14]

KIdASME

= 26.77 + 12.43 exp [0.0145 (T - RTNDT)] (2.2)

The ASME Code KIa curve, or reference crack-arrest fracture toughness lower bound

curve provides the toughness characterization for the crack-arrest and is given by [14]

KIaASME

= 29.45 + 13.675 exp [0.0261 (T - RTNDT)] (2.3)

The ASME Code reference toughness curves were developed by plotting all the known

data (KIc, KId, KIa) versus (T - RTNDT) and then drawing the lower bound curves, such that all

the experimentally measured fracture toughness values of RPV steels SA 533B, SA-508-2

and SA-508-3 would fall above the reference toughness curve.

10

ASTM test method E199 11

ASME Boiler and Pressure Vessel Code, Section III, Div 1-NB, art 2330 and Section XI, art G 2000 12

Fracture toughness is denoted by KI, the mode I (opening mode) stress-intensity factor and is given in Pa.m

1/2


p. 9

Fig 2 – The reference static fracture toughness curve as lower bound curve [13]

The 3 curves are indexed by RTNDT, being the Reference Temperature for Nil Ductility

Transition. This temperature is determined via Charpy impact test criteria, according to

ASTM E208 [15], as being equal or higher than TNDT, the Temperature for Nil Ductility

Transition. The latter is determined via a Pellini drop-weight test, being a test conducted at

5K temperature intervals until a break/no-break temperature is determined and set as TNDT.

To take the effect of neutron irradiation into account, the fracture toughness lower bound

curves are shifted by a temperature amount ΔRTNDT, which is set equal to ΔT41J, being the

temperature shift in the CVN diagram. This whole methodology is schematically represented

in figure 3.

Fig 3 - ASME methodology to obtain temperature shift in fracture toughness curve from CVN tests [13]


p. 10

Safety regulations on nuclear power plants (NPPs) require that there is sufficient

temperature margin when comparing the RPV operating conditions (both normal and

accidental) with the fracture toughness curves. Especially conditions of cold overpressure

form a significant challenge to the structural integrity of the RPVs and should be thoroughly

reviewed. A typical example is a pressurized thermal shock (PTS) event [16] in which rapid

cooling of the downcomer13

occurs due to the injection of ambient temperature makeup water

via the ECCS to compensate for loss of coolant via e.g. a pipe break in the primary circuit.

The temperature drop produced by rapid depressurization, coupled with the ambient

temperature of the makeup water, produces significant thermal stresses in the thick-section

steel wall of the RPV. For embrittled RPVs, these stresses could be high enough to initiate a

running crack that could propagate all the way through the vessel wall.

13

The downcomer section of a nuclear reactor is the annular region between the reactor vessel and the reactor core. The cooling water is directed downwards through this region towards the bottom of the RPV where it is deflected and runs upwards through the reactor core.


p. 11

2.2. Causes and mechanisms of embrittlement and recovery

Before discussing the fundamentals of embrittlement and recovery of RPV steels, a short

introduction of the typical RPV materials and in-service conditions may be opportune. The

pressure vessels of current Western PWRs were built out of high toughness, quenched and

tempered ferritic steels [7]. During vessel fabrication, base metals and welds underwent

several stress-relief heat treatments at a temperature of 550-610°C, which determined their

microstructure. Most of the base metals were low-alloyed NiMnMo ferritic steels, typically

A533B Class 1 and its forging equivalent A508 Class 3 (corresponding to 16 MnNiMo 05

French standard) or A508 Class 2. Some composition specifications are given in appendix A.

A533B- and A508-type steels have a tempered bainitic-austenitic structure in which the prior

austenite grains are about 30 μm in size. The weld metal has a lower carbon content than that

of the base metal. In general, welded joints have a bainitic-martensitic structure with a very

fine carbide distribution. We may already remark here that irradiation-induced embrittlement

is almost independent of the microstructure of the RPV steel, although the unirradiated

mechanical properties may be quite different.

Most of the Western RPVs operate at a temperature ranging from about 270°C to 330°C.

They are subject to neutron irradiation, with the peak located at the vessel core mid-plane,

also termed the beltline region. The fission process however produces neutrons of widely

differing energy levels, but almost all neutrons have an energy lower than 3 MeV. Usually

only neutrons of energy higher than a threshold value are considered, supposing that the other

neutrons have a negligible effect on the irradiated material. In Western countries, a threshold

value of 1 MeV is used. Neutron irradiation is mostly described in terms of flux (n∙cm-2

∙s-1

),

fluence (time-integrated flux and given in n∙cm-2

) or by its effect, being the number of

displacements per atom (dpa). Typical values for a French PWR after 32 years of full power

operation are [17] :

.flux 6∙1010

n∙cm-2

∙s-1

.fluence 6∙1019

n∙cm-2

.dose 0.1 dpa

RPVs are designed and fabricated in accordance with consensus codes (ASME, PED,

JIS, etc.) that are based on mechanical and physical properties of the steels used to construct


p. 12

the vessel. The fracture toughness of the RPV in the unirradiated condition is thus generally

excellent at and above room temperature. However, exposure to high energy neutrons can

result in embrittlement of irradiation-sensitive RPV materials. Embrittlement or the increased

sensitivity to brittle fracture can be subdivided into hardening and non-hardening

embrittlement [1]. Hardening embrittlement is in general terms caused by obstacles in the

material that hinder the diffusion motion of dislocations, whereas non-hardening

embrittlement is caused by solute atom segregation at the grain boundaries that locally

decrease grain cohesion strength and thus create preferential lines along which (intergranular)

cleavage fracture can develop.

Three causes are generally recognized as being the main contributors to irradiation-

induced embrittlement of RPV steels, namely matrix damage, copper-rich precipitates

(CRPs) and intergranular segregation (IGS) of phosphorus impurity atoms [5]. The first two

causes, being matrix damage and CRPs, are responsible for hardening embrittlement. The

third cause, being IGS of phosphorus, is responsible for non-hardening embrittlement or also

termed intergranular embrittlement. The latter phenomenon will not be further treated within

the context of this thesis because of the lesser importance for Western RPVs. IGS, also

known as thermal equilibrium segregation, is particularly significant in the temperature range

of 350°C to 600°C. The ensuing embrittlement is sometimes also called temper

embrittlement because it often occurs during the tempering process where precise control of

cooling time and temperature leads to a specific combination of strength and ductility of the

steel via modification of the crystalline structure. Experimental programs have shown that for

concentrations lower than 0.015wt% P, the shift in the Charpy temperature transition curve

with increasing phosphorus content was insignificant. Therefore, and since the beginning of

the 1970s, the phosphorus concentration in Western RPV steels and their welds have been

limited to a maximum value of 0.015wt% and later even to lower values. The Russian

VVER-440 RPVs had a higher phosphorus content in both the base metal (0.020wt%) and

the weld metal (0.023wt%) [6] and 15 RPVs of this reactor type were annealed during the

period 1987-1996 in order to recover from temper embrittlement. For the Western RPVs with

lower phosphorus content, the danger of IGS is much lower but nevertheless care has to be

taken as there is increasing evidence that phosphorus contributes to irradiation embrittlement

through the mechanism of matrix damage and also by participation in CRPs.


p. 13

2.2.1. Matrix damage

Matrix damage is produced when neutrons of sufficient energy displace lattice atoms that

result in displacement cascades which produce large numbers of point defects : vacancies or

unoccupied atom sites in the crystal structure and interstitials14

or supernumerary atoms in a

crystal structure [7]. In general, point defects can move over relatively long distances due to

thermal diffusion and disappear via vacancy-interstitial recombination or via sinks, such as

matrix interfaces, grain boundaries, free or cavity surfaces, etc. Vacancies in pure iron are

already mobile at 200K (-73°C) and the diffusion of interstitials is considered to occur even

much faster and at lower temperature. This means that not only at RPV operating

temperatures, but also at room temperature, vacancies and interstitials can only survive in an

aggregated form such as small vacancy clusters, microvoids, interstitial-type or vacancy-type

dislocation loops, etc. or when they associate with solute atoms such as in vacancy-solute

atom complexes/clusters. It is generally accepted that on average only 30% of the initial point

defects survive by this way [18]. It is also generally accepted that the change of the

mechanical properties of RPV steels under neutron irradiation, and in particular

embrittlement, is not due to the atomic displacements and subsequent rearrangements, which

is also termed primary damage, but due to the migration of the surviving point defects.

2.2.2. Copper precipitation

Copper precipitation in low-alloy ferritic steels is a more complex phenomenon. First of

all, there is the very low solubility of copper in iron15

. Nevertheless, Cu precipitates are not

present in the commonly used RPV steels when the vessel is commissioned. This is due to

the fact that Cu precipitation is suppressed by the quenching of the RPV steels after the

homogenization heat treatments at 600 to 700°C. Copper is at that moment in supersaturated

solution in the ferrite. During metallurgical studies, copper precipitation in these quenched

ferritic steels has only been observed during a high-temperature treatment above 500°C. Also

experimental studies by Nagai et al. [19] showed that Cu precipitation was completely absent

without irradiation after thermal ageing for 8 years at 300°C and 400°C. However, even

under low dose-rate neutron irradiation (4.2∙108 n∙cm

-2∙s

-1), copper precipitation is noticed.

14

Interstitials is the short term for interstitial atoms 15

The solubility limit of Cu in Fe at 300°C is, depending on the literature, quoted as 0.0002 wt% or 0.05 wt% [20]


p. 14

There has been an evolution in the characterization of the kinetics of point defect

migration, Cu precipitation and their possible interaction. There is the generally accepted

view that the relation between neutron irradiation and Cu precipitation is due to a super-

saturation of vacancies that enhances the mobility of the Cu atoms [18]. This is termed

irradiation-enhanced diffusion. Earlier, it was believed that the diffusion of vacancy clusters

and the diffusion of Cu atoms via irradiation-induced vacancies were two separate

mechanisms and that hardening embrittlement was caused by the random encounter or

encounter at a germ (e.g. impurity atom) of vacancy clusters and Cu precipitates.

New experiments at the nanostructure level and modeling of irradiation effects pointed

towards a second diffusion mechanism, termed as irradiation-induced diffusion, whereby

vacancies interact with Cu solute atoms to form vacancy-solute complexes [18]. By this way,

Cu solute atoms are dragged by the vacancies towards germs where small vacancy clusters or

microvoids are formed which are surrounded by Cu atoms or -in a more demonstrative way-

microvoids of which the inner surface is decorated with Cu atoms. The interaction of

vacancies and Cu solute atoms is in correspondence with the findings from thermodynamical

modeling where binding energies between Cu atoms and vacancies in α-Fe were found to be

strong between first- and second-nearest-neighbor sites (resp. 0.17 and 0.18 eV). Barashev

and Arokiam [21] have shown that because of this strong binding up to the second-nearest-

neighbor, the vacancy can turn a copper atom around without dissociating the Cu-vacancy

pair, which results in a positive interaction between vacancy and Cu solute atom diffusion in

α-Fe. This vacancy-solute atom interaction is also in agreement with 3D-AP studies on

irradiation-induced Cu clusters [22]. The AP images of the Cu clusters reveal a low

concentration in solute atoms and these loose cluster formations were therefore called

‘atmospheres’ or ‘clouds’. The question was raised why these Cu clusters keep a dilute

morphology and do not collapse into a tighter precipitate. The most probable answer is the

presence of a high vacancy concentration (microvoid) within the Cu precipitates.

Some important observations have to be summarized here [23]. First, Cu precipitation

starts already under very low neutron irradiation dose-rates (4.2∙108 n∙cm

-2∙s

-1). Secondly, the

Cu precipitation rate is very fast in the early stage of the irradiation and is independent of the

copper content. Thirdly, above a certain neutron fluence (2∙10-19

n∙cm-2

), the hardening rate

becomes constant and independent of the copper content. The latter observation has led to the


p. 15

general acceptance of the linear superposition model for irradiation-induced hardening. The

total irradiation-induced hardening can be seen as a linear superposition of solute atom

precipitation hardening and matrix hardening due to point defect clusters. Copper thus plays

a role in the very early stage of irradiation, but quite rapidly shows a saturation behavior.

Although not further treated within this thesis, nickel also has a deleterious impact on

irradiation-induced embrittlement [24-25]. Nickel can integrate with CRPs and also partici-

pate in nickel manganese-rich phases containing a small amount of copper (termed late bloo-

ming phases). The impact may become high for nickel contents of more than 1 to 1.2% and

also increases with the copper content [26].

2.2.3. Recovery of ductility

Hardening embrittlement can thus be interpreted as caused by the obstruction to

dislocation loop diffusional motion by microvoids (10 to 30 monovacancies) which are

locally surrounded by Cu atoms. Thermal annealing now consists of removal of these

obstructions. Earlier, it was thought that at around 450°C, complete microvoid-solute atom

clusters gradually dissolved. Recent annealing experiments by Nagai et al. [18] however

revealed that this dissolution occurs over two temperature stages. This phenomenon was also

revealed in the experiments carried out within this thesis.

There is a first stage at around 400 to 450°C whereby the microvoids thermally dissociate

in monovacancies which subsequently migrate into the bulk material and disappear at sinks.

The copper is thereby left behind as ultrafine (nm) defect-free Cu precipitates. There is a

second stage at around 600 to 650°C whereby the Cu precipitates thermally dissolve into

individual Cu atoms which diffuse into the bulk and homogeneously distribute as solute

atoms.

These two thermal recovery stages are the explanation for the fact that annealing at

450°C results in an almost complete embrittlement recovery, as will be indicated later in this

chapter, but only in a partial hardness recovery. It is indeed the matrix hardening that is being

recovered, leaving the precipitation hardening intact as the defect-free Cu precipitates do not

dissolve in the temperature range of 450-600°C. This phenomenon was also observed and


p. 16

confirmed by Vickers microhardness measurements during annealing experiments by Nagai

et al. [18].

These two thermal recovery stages can be related to stage IV and stage V of the electrical

resistivity recovery stages for irradiated material [27]. As neutron irradiation increases the

yield strength and hardness (and decreases the ductility), it also increases the electrical

resistivity of the material. By annealing at increasing temperatures, the electrical resistivity

can be recovered (decreased) and this process can be divided into 5 stages, each related to its

initiating annealing temperature and phenomenon at the atomic level :

.stage I : migration of self-interstitial atoms (SIAs)

.stage II : long-range migration of SIA clusters and SIA-impurity complexes

.stage III : (a longstanding controversy) near universal agreement that it is associated with

vacancy migration

.stage IV : thermal dissociation of vacancy clusters

.stage V : thermal dissociation of precipitations

Fig 4 - Five thermal recovery stages of the electrical resistivity

during isochronal annealing [27]


p. 17

2.3. Thermal annealing as remediation for embrittlement

Thermal annealing of RPVs as currently applied for recovering fracture toughness

properties after irradiation exposure, is a low temperature heat treatment, i.e. at < 500°C for

about one week. As already mentioned in chapter 1, there are two techniques for applying a

thermal annealing treatment, namely the wet technique using the coolant as the heat transfer

medium and thus being limited by design restrictions to an annealing temperature of around

340°C16

, and the dry technique using an air medium which allows a higher annealing

temperature in the range of 430°C - 500°C. For dry thermal annealing, two different heating

methods have been tested : radiant heat from electric resistance heaters and radiant heat from

an indirect gas-fired heat exchanger.

2.3.1. Practical experience

The largest practical experience comes from the successful dry annealing of 15 VVER-

440 RPVs in Russia, Armenia, Bulgaria, Slovakia, former East Germany and Finland during

the period 1987 till 199617

[28]. There are five VVER-440 RPVs still in operation after

annealing. The post-irradiation annealing (PIA) in the VVER-440 RPVs was focused on only

one circumferential seam weld in the reactor core zone, for which the irradiation

embrittlement due to phosphorus boundary segregation was considered too high. The base

metal was not found in need of a recovery treatment and so only a narrow beltline zone of 1

meter in height was annealed. This made it easier to prevent exceeding the acceptable

thermal stress level, which would otherwise lead to residual strains, and this was done using

sufficiently low heat-up and cool-down rates and controlling temperature gradients. In VVER

type RPVs the cylindrical part is constructed of forged rings and only one circumferential

weld is situated in the area of high neutron flux, whereas in the old Western RPVs the

cylindrical part is made of rolled plates, joined together with circumferential and axial welds.

In the latter case, the irradiation embrittlement is encountered all over the core zone and

extends to the nozzle rings, which makes it difficult to avoid excessively high thermal

stresses. This may lead to unwanted dimensional changes in the RPV and associated piping.

It is clear that the subject of thermal stresses will be a key issue in demonstrating the

feasibility of thermal annealing in different types of reactor constructions.

16

A sufficient margin has to be taken from the critical temperature of water Tcrit = 374°C 17

See for list in Appendix B


p. 18

Apart from these 15 VVER type RPVs, only three additional reactors have been

annealed worldwide. Two reactors were wet annealed and have already been

decommissioned : the US Army SM-1A annealed in 1967 and shut down in 1972, and the

Belgian BR3 annealed in 1984 and subsequently operated until 1987. The third reactor was

dry annealed within a demonstration project conducted at the canceled/uncompleted Marble

Hill plant, Indiana, in 1997 with the aim to show the engineering feasibility of performing a

dry annealing using an indirect gas-fired heat exchanger in a US-designed RPV. The Marble

Hill plant was partially completed with the vessel in place and provided a unique opportunity

to test a dry annealing treatment of a large commercial US vessel, albeit a non-irradiated one.

A second demonstration project was scheduled for early 1997 at the canceled/uncompleted

Midland plant, Michigan, for testing dry annealing using electric resistance radiant heaters,

but this project was delayed indefinitely.

The US regulatory framework for executing PIA on RPVs is fully in place since 1997. In

the mid-1990s, the Palisades power plant at Covert, Michigan was considering a dry thermal

annealing of their RPV. This plan was canceled later when fluence re-evaluation calculations

showed that the plant could just meet projected end-of-life PTS regulatory requirements.

However, in preparation for the planned annealing of the Palisades RPV, the necessary

related documents were issued by USNRC, ASME and ASTM18

. As of today, the US nuclear

industry has not yet made use of this regulatory framework for non-technical reasons.

2.3.2 Annealing equipment

Worldwide there are 4 full sets of RPV annealing equipment [29] : Russia has two sets

for dry annealing with electric resistance radiant heaters, Germany has one similar set and the

US has developed and tested a set for dry annealing by means of an indirect gas-fired heat

exchanger.

The two Russian sets are very similar and differ only slightly in length and number of

individually controlled heat zones. Russian set 1 is shown in figure 5, has 54 heater panels in

three elevations and they form 9 individually controlled groups. The heater panels are about

20 cm away from the vessel surface. The temperature is monitored by 18 thermocouples

brought in contact with the inner surface of the reactor vessel by disengaging mechanisms.

18

See for a list of the documents in Appendix C


p. 19

The temperature is also measured at several points on the outside of the vessel. These

temperature measurements provide information to the annealing control system that monitors

the temperature distribution profile and the ensuing thermal stresses in the vessel wall. The

maximum output of the heater is 800 kW, but during the heat-up phase only about 225 kW is

actually used and about 150 kW is needed to keep the vessel at the holding temperature. In

order to limit thermal stresses, the maximum heat-up rate was set at 20°C/h and the cool-

down rate at 30°C/h. The temperature variations on the inner surface of the annealing zone

had to be within 50°C. The maximum temperature on the outer surface remained about 20°C

below the temperature of the inner surface. The maximal thermal stresses, using a 1 meter

nominal annealing zone and a 20°C/h heating rate, were 200 MPa after 32 h heating. The

Russian set 2 shown in figure 6 has 13 individually controlled heating zones and has a 975

kW output. The most important difference compared to set 1 is that set 2 has heater elements

also outside the actual targeted annealing area. This means that the control of the axial

temperature profile is more flexible.

Fig 5 - Russian annealing equipment set 1 [30]


p. 20

Fig 6 - Russian annealing equipment set 2 [31]

In Germany, Siemens KWU19

has designed and constructed an electric resistance radiant

heating annealing device for a 600 MW reactor. It has, however, never been used.

In the US, the feasibility of dry thermal annealing using an indirect gas-fired heat

exchanger, as shown in figure 7, was demonstrated at the Marble Hill plant, Indiana, in 1997

[11]. Gas-fired burners provide superheated air that is blown through sealed ducts in existing

openings in the containment building, such as the equipment hatch, into a gas-gas heat

exchanger placed inside the reactor vessel. The superheated air is then discharged outside of

the containment through another duct into the atmosphere. Since the air never comes in

19

Siemens subsidiary Kraftwerk Union AG


p. 21

contact with any contaminated surface, it does not become contaminated. The gas-fired

heaters are located outside the containment building so they can be easily serviced in case of

failure. The partially completed plant had a brand new RPV in place which was a

Westinghouse design four-loop pressurized water reactor (H=7m, D=5m)20

with nozzle

supports. The annealing heat exchanger was designed for potential reuse and easy clean-up

after the annealing procedure. The demonstration was unique in that the whole vertical vessel

wall up to the nozzle ring was annealed up to a temperature of 454°C at the inner surface of

the reactor vessel for a time period of one week. The heat exchanger reached a temperature of

about 593°C during this process. Analytical stress models of the Marble Hill vessel and the

reactor coolant system were shown to be correct based on measured temperatures and strains

in the actual vessel during the annealing process. Documentation of critical vessel

dimensions both before and after the annealing procedure confirmed that all vessel interfaces

and dimensions were maintained within acceptable tolerances.

Fig 7 - Marble Hill annealing heating system [32]

20

H stands for height ; D stands for diameter


p. 22

Fig 8 - Marble Hill heat exchanger prior to its installation inside the reactor vessel [32]

2.3.3. Recovery and reembrittlement : parameters

The key elements with respect to continued operation of a RPV after annealing are the

degree of recovery and the reembrittlement trend. Ideally, both of these elements should be

measured using existing surveillance capsules containing the limiting reactor beltline

materials and this by Charpy V-notch (CVN) impact tests or alternatively, fracture toughness

tests or via deduction of the toughness property changes from tensile, hardness, indentation,

or other miniature specimen testing.


p. 23

The degree of recovery of the Charpy impact toughness properties is referenced to the

shift in the transition temperature T41J and the drop in USE, both due to neutron irradiation

[33]. A 100% recovery would indicate that the values of T41J and USE after annealing have

fully recovered their unirradiated values.

The percentage recovery of the T41J is defined as the ratio of the residual transition

temperature shift after annealing to the shift due to irradiation before annealing :

% recovery T41J =

. 100 (2.4)

The percentage recovery of the USE is defined in a similar manner :

% recovery USE =

. 100 (2.5)

The reembrittlement trend will have to be measured by testing surveillance capsules,

but for an immediate quantitative assessment a reembrittlement trend curve has to be

adopted. The Standard Guide ASTM E 509-03 [12] provides two methodologies that

encompass the possible reembrittlement trends. The conservative methodology of post-anneal

reirradiation trend curve development is schematically shown in figure 9 and is termed

lateral shift since the initial irradiation curve is merely translated to project reirradiation

behavior. An alternative trend curve approach is the vertical shift where the portion of the

initial irradiation trend, projected as reirradiation behavior is translated down vertically (as if

annealing would not have any influence) as shown in figure 10. The use of this estimated

trend curve should however be justified with actual post-anneal reirradiation data since the

vertical shift method predicts significantly lower changes in RTNDT after thermal annealing.


p. 24

Fig 9 - Lateral shift as reembrittlement trend curve approach [12]

Fig 10 - Vertical shift as reembrittlement trend curve approach [12]


p. 25

2.3.4. Recovery and reembrittlement : general considerations

It is difficult to give general data concerning recovery and reembrittlement, because of

the large number of factors that influence these processes. These factors not only include the

ones that influence embrittlement, such as neutron fluence, neutron flux spectrum, alloy

composition, radiation temperature and microstructural characteristics (grain size,

metallurgical phase, etc.) but also the annealing temperature and time. Several studies were

made on the subject of the response of irradiated RPV steels to thermal annealing. The results

of the study of Iskander et al. [33] at ORNL are here presented as an example. Figure 11

shows the response of Charpy impact energy to annealing at 343°C for 168 h of HSSI weld

73W21

Charpy specimens in the unirradiated, irradiated and irradiated/annealed conditions.

The USE has recovered 45% of the drop lost to irradiation, but the recovery of the T41J is

insignificant.

Fig 11 - Response to CVN curve of annealing at 343°C for 168h of HSSI weld 73W [33]

Figure 12 shows the response of the Charpy impact energy to annealing at 454°C for 168

h of the same material, from which can be seen that the recovery of T41J is nearly complete

(92%) whereas the USE shows a remarkable over-recovering (187%). The latter phenomenon

has been also reported by other researchers and the reasons for this phenomenon are still

21

High Cu content, low USE


p. 26

under investigation. It should also be noted that the USE recovers at a much faster rate than

the T41J.

Fig 12 - Response to CVN curve of annealing at 454°C for 168h of HSSI weld 73W [33]

From these results can already be deduced that if the annealing time is taken sufficiently

large (e.g. the usual 168 h), then the annealing temperature is the major determinant in the

degree of recovery. During preparation and planning of an in-service anneal, sufficient

attention should be given to the selection of the annealing temperature as it requires a balance

of opposing conditions. Higher annealing temperatures can produce greater recovery of

fracture toughness and thereby increase the post-anneal lifetime. On the other hand, higher

temperatures can create other undesirable property effects such as permanent creep

deformation, temper embrittlement or dimensional deformations of the vessel, nozzles and

attached piping due to excessive thermal stresses during the annealing.

A JRC AMES22

report [34] on the subject comes to the following common conclusions

about the role of the annealing temperature and the role of impurities for PWR and VVER-

40 material recovery : at an annealing temperature of 340°C only an insignificant recovery in

T41J for irradiated material is observed, being 20% on average. For temperatures in the

22

Ageing Materials European Strategy. This EU project brings together the organizations in Europe with main capabilities on RPV materials assessment and research.


p. 27

interval of 340-420°C the dependence of T41J recovery is close to linear ; in the temperature

interval 420-470°C the recovery is about 80% or more, which corresponds to a residual

embrittlement after annealing not larger than 20-30°C. For material with low level of P or Cu

annealing at temperatures in the interval of 450-470°C leads to nearly full recovery of T41J. If

phosphorus content is more than 0.02% and copper content is more than 0.2%, the residual

shift of T41J could be up to 30-40°C.


p. 28

2.4. CDBAR as technique for nanoscale material studies

Until the 1980s, no microstructural studies had been able to provide direct information

on irradiation-induced defects in RPV steels. Even transmission electron microscopy does not

have sufficient resolution to reveal such nanoscale defects. The first direct information about

the nature and structure of irradiation-induced defects was obtained when characterization

techniques with very high spatial and chemical resolutions became accessible for industrial

applications. The most commonly used techniques are high resolution transmission electron

microscopy (HRTEM), atom probe tomography (APT), small angle neutron scattering

(SANS) and positron annihilation spectroscopy (PAS).

Positron annihilation spectroscopy is a well established technique used to characterize

features containing vacancies (e.g. vacancy clusters, vacancy-solute clusters) [35]. The

method is based on trapping and subsequent annihilation of positrons at open-volume defects

whereby 2 γ-photons of 511 keV are emitted in opposite directions. There are three basic

methods of analyzing the signals : one based on the lifetime of the positron and called

positron annihilation lifetime spectroscopy (PALS) and two based on the momentum

distribution of the annihilating electron-positron pair and called angular correlation

(positron) annihilation radiation (ACAR) spectroscopy and coincidence Doppler broadening

(positron) annihilation radiation (CDBAR) spectroscopy.

Fig 13 - The scheme of positron experiments (PAS techniques) [36]


p. 29

2.4.1 The positron

Paul Dirac23

published in 1928 a paper24

[37] that suggested the possibility of an electron

having both a positive charge and a negative energy. This paper introduced the Dirac

equation, a unification of quantum mechanics, special relativity and the concept of electron

spin. The paper did not explicitly predict a new particle, but did allow for electrons having

either positive or negative energy as solutions. Paul Dirac developed his theory further in a

paper25

[38] in 1929 in which he predicted the existence of an as-yet unobserved particle that

he called an ‘anti-electron’ that would have the same mass as an electron and that would

mutually annihilate upon encounter with an electron. The particle was discovered in 1932 by

Carl D. Anderson26

when studying cosmic radiation with a cloud chamber and termed by him

as ‘positron’27

[39]. Only one year later, it was realized that positrons could be produced

through different mechanisms such as pair production or radioactive decay [40].

Positrons find their useful application mainly in three domains [41]. The best known

application is the one of positron emission tomography (PET) scanners whereby positrons

created through the decay of a radioactive tracer, are followed through the body [42]. PET

scanners can thus create detailed 3D-images of metabolic activity within the human body.

Another application is the use of positrons as charged particles in particle accelerator

experiments where by collisions at relativistic speeds subatomic particles are created and

studied by physicists [43]. The application that will be further treated in this thesis is the one

of positron annihilation spectroscopy (PAS), a technique used in materials research to study

defects within solid materials [36].

Conventional sources of positrons for PAS are artificial radioisotopes emitting β+-

radiation [36]. Energy spectra of positrons emitted by such radioisotopes are continuous,

ranging from zero to an end-point energy which is typically of the order of 0.1-1 MeV. Mean

penetration depths in solids are typically of the order of 10-100 μm, making such positrons

valuable elements to probe volume properties in matter. The most commonly used source is

the 22

Na radioisotope. A simplified decay scheme is shown in figure 14. 22

Na decays to an

23

Paul Dirac (1902-1984) received the Nobel Prize for physics in 1933 24

‘The Quantum Theory of the Electron’ 25

‘A Theory of Electrons and Protons’ 26

Carl D. Anderson (1905-1991) received the Nobel Prize for physics in 1936 27

Contraction of ‘positive electron’


p. 30

excited state of 22

Ne with a β+

branching ratio of 90%. This excited state has a mean lifetime

of 3.7 ps and de-excites to the ground state of 22

Ne by emitting a 1.274 MeV photon. This

photon can be used to register the positron’s birth. 22

Na has a conveniently long lifetime of

about 2.6 years and is commercially available as a NaCl solution. Because β+-decay is a two

particle decay (positron+neutrino), the positrons emitted by 22

Na have a broad energy

distribution extending from almost zero to 545 keV.

Fig 14 - Simplified decay scheme of 22

Na [36]

2.4.2. Positron thermalization

Positrons rapidly lose their energy when injected in matter [41]. The positrons scatter

elastically and inelastically off ion cores and conduction electrons and within picoseconds

reach energies of about 1 eV, followed by a slower approach towards thermal equilibrium at

0.025 eV. The whole thermalization process is very fast (1-10 ps) compared to the positron

lifetime in solids (typically 100-500 ps), and this despite of the huge energy transfer that

takes place during the thermalization. The fast positron thermalization process has also been

proven experimentally by ACAR measurements [44]. In the same context, a study by Jensen

et al. [45] showed that non-thermal trapping of positrons in metals is marginal and needs not

to be considered for PAS studies of open-volume defects in metals. Remark that the ratio of

the positron thermalization time to the positron lifetime is indeed a good criterion for the

usefulness of the annihilation method for studying the energy or momentum spectrum of the

annihilating electrons in the material, since if a positron is thermalized the basic contribution

to the momentum distribution of the centers of mass of the annihilating electron-positron

pairs comes from the annihilating electrons.


p. 31

The mean time a positron spends at high energy is thus negligible and therefore only the

two- and three-photon annihilation should be taken into account according to quantum

electrodynamics (QED). The ratio of two- to three-photon annihilation is approximately

371.2. This value was experimentally confirmed in metals by triple coincidence

measurements [46].

2.4.3. Positron trapping

After reaching thermal equilibrium with the host, the positron state develops as a

diffusion process. During this diffusion, the positron interacts with its surroundings and

eventually annihilates with an environmental electron. In homogeneous defect-free media, all

positrons annihilate at the same rate which is a characteristic of the given material. Due to the

Coulomb repulsion by the positive ion cores, positrons tend to preferably reside in interstitial

regions. Furthermore, positrons are ‘attracted’ by locations where they can additionally relax

and lower their energy. These potential well centers have two origins [47] :

(i) Open-volume defects (vacancies, vacancy clusters, voids, dislocations, misfit defects

at precipitate-matrix boundaries, etc.) where the potential sensed by the positron is

lowered due to the reduction in Coulomb repulsion between the positron and the ion

cores. Also the redistribution of the electrons causes a negative electrostatic potential at

this type of defect, that additionally attracts positrons. As a result, a localized positron

state at a defect can have a lower energy than the state of the delocalized or free

positron. The transition from the delocalized state to the localized one is called positron

trapping. Positron binding energies Eb to defects like e.g. monovacancies are typically

of a few eV.

(ii) Precipitates or clusters of atoms that have a larger (negative) positron affinity A+ than

the surrounding atoms. The positron affinity A+ is defined as the sum of the positron

and electron chemical potential. The positron chemical potential is defined as the

energy of the lowest positron state relative to vacuum. When two metals are in contact

with each other the Fermi levels equalize, inducing a potential with a depth defined by

the difference in positron affinity A+. The importance of A+ consists in the fact that the

difference of the positron energies in different materials with which they are in contact

is just the difference in their A+ values. Positrons try to occupy the lowest absolute


p. 32

energy level and therefore go into the material having the lowest (negative) A+ value,

e.g. prefer to go towards Cu (A+ = -4.81 eV) instead of towards Fe (A+ = -3.84 eV). A

minimum precipitate radius rc = 0.31. is needed for formation of the bound

positron state, where ΔA+ is the positron affinity difference between the matrix and the

precipitate [48]. Under favorable circumstances, nano-precipitates consisting of just a

few atoms only become trapping centers for positrons.

2.4.4. Positron annihilation

In vacuum a positron is a perfect stable particle, just as an electron. In matter, it will

annihilate with its counterpart, the electron, thereby converting their masses in energy. In the

theory of Dirac, the positron-electron annihilation can be seen as the radioactive de-excitation

of the electron [38]. This process can be described by QED and may proceed by the creation

of zero, one, two or three photons within the constraints of energy, momentum and spin

conservation. When the positron has a low kinetic energy, the most probable is annihilation

with creation of two or three γ-photons.

The kinetic energy of the annihilating positron-electron pair is typically a few electron

volts. In the frame of reference of the center of mass of the electron-positron pair, the two γ-

photons produced in the annihilation have the same energy28

,

E0 = hν0 = m0c² = 511 keV, (2.6)

and they emerge in opposite directions which means that the angle between them is π. In the

laboratory frame of reference, however, the velocity of the center of mass is v and the linear

momentum29

of the electron-positron pair is therefore

p = 2m0v (2.7)

As a result of momentum conservation during the annihilation process, the momentum of

component pz in the propagation direction z of the γ-photons results in a Doppler shift ΔEγ of

the annihilation energy of 511 keV, which amounts approximately to

ΔEγ cpz/2 (2.8)

28

With h being Planck’s constant, m0 being the electron (positron) rest mass and c being the speed of light 29

for the remaining of this thesis ‘momentum’ has to be read as ‘linear momentum’, unless otherwise stated.


p. 33

Since numerous annihilation events are measured, the energy line of the annihilation is

broadened due to the individual Doppler shifts in both ± z directions, and a complete Doppler

spectrum is obtained. This effect is utilized in CDBAR spectroscopy.

The momentum components px,y perpendicular to the propagation direction lead to

angular deviations θx,y of the collinearity of the annihilation γ-photons according to

θx,y px,y / m0c (2.9)

These equations hold for small angles only. The θx,y can be registered simultaneously in both

x and y directions by a coincidence measurement using position-sensitive detection of the γ-

quanta. This technique is ACAR spectroscopy.

Conservation of momentum during the annihilation means that the momentum of the

annihilating electron-positron pair is equal to the total momentum of the annihilation γ-

photons. As we can consider the momentum of the thermalized positron to be zero, the

momentum of the annihilation radiation is then just the momentum of the annihilating

electron. Measuring the distribution of momenta of the annihilation radiation will give the

distribution of the electron momenta at the (defect) site at which the positron settles [36].

2.4.5. PALS

Positron annihilation lifetime spectroscopy is based on the measurement of time between

implantation and annihilation of the positron [49]. This is made possible by the fact that in

the 22

Na source a γ-quantum with energy 1.274 MeV is emitted almost simultaneously with

the positron. The positron energy, which extends up to 540 keV, decreases in the sample

within a few picoseconds by non-elastic interactions to around 0.1 eV. The mean positron

penetration depth of this so-called thermalization process is of the order of 100 μm. The

thermalization time of a few picoseconds is small compared with the positron lifetime (100-

400 ps) and can be neglected as well as possible non-thermal trapping/annihilation of

positrons. Another valid argument for neglecting the positron thermalization time is the fact

that in PALS the positron lifetime is measured from the moment of the emission of the γ-

quantum of the 22

Na source which happens 3.7 ps after the emission of the positron (see fig

14). On reaching thermal energies, the positron diffuses in the material before it is trapped in

a lattice defect. The diffusion length is of the order of 100 nm.


p. 34

The positron lifetime of a single event can be measured by detecting the time difference

between the birth γ-quantum of the β+-decay in the source and one of the annihilation γ-

quanta of energy 511 keV in the sample. The activity of the source must be sufficiently low

in order to ensure that on average only one positron is present in the sample. This avoids the

intermixing of start and stop quanta originating from different annihilation events. The

scheme of a possible positron lifetime setup is shown in figure 15. PALS measurements on

irradiated steel are however severely complicated by the induced 60

Co-activity present in the

steel. 60

Co emits two coincident γ-rays30

that produce disturbing signals in a classical 2-

detector lifetime setup. Therefore, for better distinction between 22

Na-events (the positron

source) and 60

Co-events, a 3-detector setup is more appropriate. By this way, the three

coincident photons31

with known energy of the positron annihilation can be measured. The

count rate in the lifetime spectrum in this setup is however reduced by a factor of 5 [50].

When positrons are trapped in open-volume defects, such as in vacancies and their

agglomerates, the positron lifetime increases with respect to the defect-free sample. This is

due to the locally reduced electron density of the defect. Thus, a longer lifetime component,

which is a measure of the size of the open volume, appears. The intensity of the component is

directly related to the defect concentration. In principle, both items of information, i.e. the

kind and the concentration of the defect can be obtained independently by a single

measurement. This is the major advantage of PALS compared with ACAR and CDBAR

spectroscopy.

Fig 15 - Scheme of a PALS setup [49]

30

1173 keV and 1332 keV 31

The 1274 keV start-γ and the two 511 keV annihilation-γ’s


p. 35

2.4.6. ACAR spectroscopy

Angular correlation of annihilation radiation spectroscopy is based on the deviations θx,y

of the collinearity of the annihilation γ-rays due to the momentum components px,y

perpendicular to the propagation direction z [50]. Since the annihilation γ-quanta are emitted

simultaneously, θx,y can be measured in a coincidence arrangement by position-sensitive

detectors, such as multi-wire proportional chambers or Anger cameras. By this way a two-

dimensional angular correlation with the momentum distribution can be measured. The

scheme of a possible angular correlation setup is shown in figure 16. The sample-detector

distance amounts to several meters in order to be able to measure the small angles θx,y with

sufficient resolution. Because of this large distance, much stronger sources compared with

conventional PALS and CDBAR spectroscopy measurements are required. In order to

minimize the reduction of the counting rate due to the distance of several mm between

sample and source, a strong magnetic field32

of about 1T is usually applied to guide the

positrons to the sample. The typical angular resolution is in the range of 0.2x0.2 mrad². The

energy resolution of a corresponding Doppler broadening experiment would be in the range

of 0.05 to 1.3 keV (which corresponds to 0.2 to 5.1 mrad)33

. Thus, ACAR spectroscopy has a

much better momentum resolution than CDBAR spectroscopy.

Fig 16 - Scheme of an ACAR spectroscopy setup [50]

32

In tesla - unit of magnetic flux density 33

A Doppler shift of 1 keV corresponds to an angle deviation of 3.91 mrad


p. 36

2.4.7. CDBAR spectroscopy

Coincidence Doppler broadening (positron) annihilation radiation spectroscopy measures

the momentum distribution of the annihilating electron-positron pair [50]. When a positron

annihilates with an electron, two γ-photons are emitted which have each an energy of

approximately 511 keV (= m0c²) and with a total energy equal to 2m0c²-EB , with m0c² the

electron rest mass energy and EB the binding energy of the electron. Due to the Doppler shift

effect explained above, the energy of one photon is increased by cpz/2, whereas the energy of

the other photon is decreased by the same amount, with pz the component of the electron-

positron momentum along the photon emission direction. A detector will see both the red

shifted (less energy) and blue shifted (more energy) photons as they are emitted randomly in

space. This causes a broadening of the annihilation photon-energy peak by 2ΔE = cpz.

In principle, a single high-resolution energy-dispersive detector could suffice to measure

the momentum shift of the positron annihilation radiation, as in the setup shown in figure 17.

Usually liquid-nitrogen-cooled pure germanium (HPGe) crystal detectors are used due to

their high energy resolution and reasonable efficiency34

for high-energy photons. A single

detector setup however leads to relatively high background contributions in the momentum

spectrum, especially at the low-energy side of the annihilation peak mainly due to photon

scattering effects. The main background contributions come from (i) Compton scattered

1.274 MeV photons that are emitted by the 22

Na-source in coincidence with a positron (ii)

511 keV photons that have lost a small amount of energy through low-angle scattering with

the sample or the shielding (iii) incomplete charge collection within the detector (iv)

summing effects35

. The background in the momentum spectrum can be drastically reduced36

by the application of the coincidence Doppler-broadening technique, whereby both

annihilation γ-quanta are registered. The second γ-quantum, being almost collinear to the

primarily monitored quantum, can be detected in a second HPGe detector. The experimental

setup for this technique is shown in figure 18. This technique is especially important for the

measurement of the high-momentum part of the spectrum, as far as 9 keV away from the

center of the annihilation line. Positron annihilation with low-momentum valence or

34

Around 20% 35

e.g. a Compton scattered 511 keV and 1274 keV photon can be detected simultaneously and if their energy sum is around 511 keV, they cannot be distinguished from an annihilation event 36

By a factor 1000


p. 37

conduction electrons result in a small Doppler shift. Annihilations with core electrons result

in a large Doppler shift, contributing to the wings of the 511 keV annihilation line. As the

core electrons are not, or only little effected by the chemical environment, they are unique for

each type of atom. Therefore, the shape of the core electron contributions to the annihilation

spectrum forms some kind of a fingerprint of the atoms to which these core electrons belong.

By this way, CDBAR spectroscopy can be used for the chemical identification of the atoms

at the site of the annihilation.

Fig 17 - Scheme of a non-coincidence DBAR spectroscopy setup [50]

Fig 18 – Scheme of a coincidence DBAR spectroscopy setup [50]


p. 38

The result of the measurement is a two-dimensional array of counting rates, where the

dimensions represent the energy scales of the respective detectors. An example is given in

figure 19. The elliptical region obtained as the intensity profile along the diagonal from the

upper left to the lower right of the measured array, is a coincidence spectrum of both HPGe

detectors, whereby both the γ-photons have an energy of E 1= E2 511 keV. This diagonal

profile can be explained by momentum conservation during the annihilation process. The

increase in the annihilation γ-ray energy in one detector according to the Doppler shift leads

to a simultaneous reduction of the γ-ray energy in the second detector, i.e. the sum of the

annihilation γ-ray energies remains constant at 1.022 MeV. The profiles parallel to the axes at

the energy of 511 keV are also coincident spectra produced by coincidences of a 511 keV

photon with a background photon. The one-dimensional spectrum is obtained by projection

of the two-dimensional spectrum on the E1+E2 = 1.022 MeV axis. The projection includes

events in a window along the diagonal with 2m0c²-δ E1+E2 2m0c²+δ with δ the width of

the window (typically 1 keV up to 4 keV). The width of this window should be large enough

to include the contribution of the high-momentum electrons as they are shifted towards lower

energies due to their large binding energy Eb.

Fig 19 - Two-dimensional CDBAR spectrum [50]


p. 39

Fig 20 – One-dimensional CDBAR spectrum [51]

The quantitative evaluation of the effect of positron trapping in defects on the Doppler-

broadened spectrum N = f(E) can be carried out with specific line shape parameters [52].

The S parameter37

is defined as the area of the central low-momentum part of the

spectrum, AS, divided by the area below the whole curve A0

S = AS /A0 with A0 =

(2.10)

The W parameter38

is taken in the high-momentum region far from the center. It is

calculated as the area of the curve in a fixed energy interval, AW, divided by A0

W = AW /A0 with AW =

(2.11)

37

The S parameter is also called shape parameter of valence annihilation parameter 38

The W parameter is also called wing parameter of core annihilation parameter


p. 40

The interval limits are shown symmetrically around the energy of E0 = 511 keV for the

calculation of the S parameter : E0 ES. The energy limits E1W and E2W for the W parameter

must be defined in such a way as to have no correlation effects with the S parameter. The

choice of ES, E1W and E2W is to some extent arbitrary. The highest sensitivity of the S

parameter for changes in the line shape is usually obtained if the S parameter is close to 0.5.

In practice, the limits ES for the central area are chosen once, e.g. for the reference spectrum,

such that Sref = AS,ref /A0 = 0.5. These limits are then used for all other spectra that have to be

compared to each other. The limits EW for the W parameter are even more arbitrary. They

can be chosen such that Wref = AW,ref /A0 = 0.05 or such that the W parameter value is 25% of

the S parameter value.

Remark that the S and W parameters do not have an individual meaning. They can only

be used in relative terms to discuss trends in momentum distribution spectra. The S parameter

is sensitive to changes in low-momentum contributions for the annihilation peak, whereas the

W parameter is sensitive to variations in high-momentum contributions. When positrons are

trapped by vacancies, the S parameter will increase because the fraction of positrons that

annihilate with valence or conduction electrons (low-momentum) is larger than the fraction

that annihilates with core electrons. The W parameter gives then an indication of the fraction

of positrons that annihilates with core electrons (high-momentum). The W-parameter will

thus be influenced by the chemical environment at the annihilation site.

As previously mentioned, the shape of the core electron contributions to the annihilation

spectrum forms some kind of a fingerprint for the atoms to which these core electrons belong.

On figure 21, the fingerprints of a few pure elements are shown. Fe, Ni and Cu have an

electronic structure of respectively [Ar].3d6.4s

2, [Ar].3d

8.4s

2, [Ar].3d

10.4s

1. The shape of the

Doppler-shifted momentum profiles are similar, but the more 3d-electrons an atom has, the

broader the spectrum. Al and Pb have no 3d electrons, so their profiles are smaller in the

region of high momentum.


p. 41

Fig 21 - Doppler profiles of Fe, Cu, Al, Ni and Pb [51]

To emphasize the difference between the chemical elements often (half) ratio curves are

made. Figure 22 gives an overview of the ratio curves of some pure chemical elements to

pure Fe. The figure shows that the CDBAR technique can give valuable information on the

presence of precipitates in a matrix if the precipitating element can trap positrons such as is

the case with Cu. On figure 22, it can be noticed that the form of the ratio curves of Cu and

Ni to Fe are very similar. It would therefore be very difficult (if not impossible) to distinguish

between precipitates of these two elements if both would be present in the sample. Figure 22

also shows the typical fingerprint of pure Cu in the ratio curve to pure Fe : a broad peak

around 24 mrad and a small valley at around 6 mrad. The ratio curve of Ni to pure Fe shows

a smaller broad peak around 24 mrad. The Doppler-shifted momentum profiles of Cu and Ni

are lying above the Fe profile in the high momentum region. The Doppler-shifted momentum

profile of Al, Pb and Si are lying below the Fe profile in the high momentum region. Their

ration curves show a broad valley in the high momentum region and an enhancement above

one at the low momentum region.


p. 42

Fig 22 - Ratio curves of pure elements to Fe [51]

Chapter 3 : Experiments

3.1. Binary alloy samples investigated

Isochronal annealing for 30 minutes at temperatures from 300°C to 700°C in steps of

50°C and subsequent analysis by means of CDBAR spectroscopy was performed on 2 binary

model alloys which differ only by neutron irradiation dose :

.Fe-0.3wt%Cu ; 0.1 dpa

.Fe-0.3wt%Cu ; 0.2 dpa

Of each model alloy, two samples were made available with dimensions 10x10x1mm³.

The nominal composition of the samples is given in table 1 in mass percentage (wt%). The

analyses were performed by induced coupled plasma mass spectroscopy (ICP-MS) at EDF

R&D (France) for the metallic elements, while C and N were measured using the combustion

technique. The carbon concentration of the alloys is below the detection limit of the

technique (0.005 wt%).

Table 1 The chemical composition of the binary alloys obtained by ICP-MS in wt% [53]

Material Composition in wt%

Fe-0.3 wt% Cu C N Si P S Mn Ni Cu

< 0.005 < 0.001 0.012 < 0.005 < 0.005 0.010 < 0.005 0.315

The alloys were fabricated at EDF using argon-arc melting and zone refinement

methods. Argon-arc melting is a melting procedure that heats charged material by means of

an electric arc under an atmosphere of argon. Zone refining is a method of separation by

purification in which a molten zone traverses a long ingot of impure metal. The impurities

concentrate in the molten region, which is moved to one end of the ingot, and removed at the

end of the process. The resulting ingots were cold worked after tempering. A final heat

treatment for 1 hour at 1075 K in an argon atmosphere was applied in order to release the

stresses and to obtain well recrystallized material. This heat treatment was followed by a

water quenching.


p. 44

3.2. Annealing setup and work method

The 2 model alloys were isochronally annealed for 30 minutes over a temperature

trajectory from 300°C to 700°C in steps of 50°C. The annealing took place in an electric oven

placed in a hot cell such that all manipulations inside the hot cell had to be performed by

means of telemanipulators. The work procedure is described in Appendix D. A schematic

overview of the annealing setup39

is shown in fig 23. The fixed part of the equipment is the

so-called sample holder head through which air can be extracted or gas (air/He/N) can be

supplied. Furthermore there are 2 independently moveable parts, namely the vacuum/

thermal-resistant quartz tube and the electric oven. During the cooling phase, the electric

oven is moved aside and the quartz tube and its content are cooled by means of free cooling

to the air (hot cell is ventilated).

Fig 23 - Schematic overview of the annealing setup

One of the problems encountered was oxidation of the samples during the annealing

process. Although one may expect influence of the oxidation on the results through the

oxygen, especially as the positrons annihilate in the upper layer of the samples, this turned

out to be only minor. Nevertheless, it is important to avoid oxidation as much as possible to

keep the samples intact and not loose material during the different annealing steps as the

39

This setup has been explained in more detail in [54]


p. 45

samples are already thin (1mm). An obvious manner to proceed is to extract the air and thus

the oxygen. This has however a serious impact on the heat transfer needed during the

annealing process. It means a considerable longer time needed to bring the samples to the

desired set temperature as the quartz tube is brought to high vacuum (10-6

bara) and few

atoms are left to assure the heat transfer. Therefore an alternative method was used whereby

the high vacuum of 10-6

bara was lowered to 10-3

bara by means of injection of helium gas.

Helium is an inert gas so no additional oxidation has to be expected and it is also known as

an excellent heat transfer gas. However still minor oxidation was observed during the

experiments. An additional measure that was not taken during these experiments but is used

at ONRDC (Japan) is adding a piece of zirconium to the sample basket. Zirconium is more

susceptible to oxidation than iron as its standard potential is more negative (E° = -1.53 V)

than the standard potential of Fe (E° = -0.440 V) [55].


p. 46

3.3. CDBAR setup and work method

A schematic overview of the experimental setup [56] is given in fig 24. A 22

NaCl

positron source with a typical activity of 2 MBq is placed between two identical binary alloy

samples (10x10x1mm³) fixed into a sample holder. By this way, nearly all emitted positrons

enter the samples. The source is made by evaporating 22

NaCl onto a Kapton40

foil of 7 µm

thickness, which is then sealed off with an identical Kapton foil. Due to the β+-decay, the

source emits positrons with a broad energy spectrum ranging from zero up to 545 keV. The

positrons are implanted in the samples from the surface to the bulk and annihilate there with

an electron into two collinear 511 keV photons.

Fig 24 - Schematic overview of a CDBAR spectroscope setup (HV = high voltage, DSP = digital signal

processor, INT = interface card) [56]

40

Kapton is a polyimide film developed by DuPont. It is used as insulation material in aircraft and spacecraft applications. It is used preferentially by the positron community for its transparency and relative insensitivi- ty to radiation damage


p. 47

The photons are detected by two movable HPGe detectors (Canberra type GC3018)

facing each other and placed collinearly at 18 cm from the sample. The detectors have an

input count rate of 10 kHz per detector, an efficiency of 25% and an energy resolution

(FWHM)41

of typically 0.8 keV at the 122 keV line of 57

Co and 1.80 keV at the 1.332 MeV

line of 60

Co. The outputs of the detector preamplifiers (Canberra PRE AMP model 2101P)

are connected to two digital signal processors (DSP, Canberra, model 2060) units. The digital

signals of the DSP units are further processed by an interface (INT) card. Finally, the data are

transferred to the PC by an internal digital data acquisition (DDA) card42

. This card is based

on PCI technology43

and allows a fast transfer of data to the PC (up to 106 events per second).

Actually, for each measurement more than 106

coincidence counts per hour are recorded

during approximately 24 hours. The analysis software based on LabVIEW44

constructs a two-

dimensional energy spectrum (energy of detector 1 on the x-axis and energy of detector 2 on

the y-axis) as shown in fig 25.

Fig 25 - Example of the LabVIEW software output with the 2-dimensional energy spectrum [53]

41

FWHM = Full width at half maximum 42

National Instruments type PCI-IO-32HS 43

PCI = Peripheral component interface : computer bus for connecting hardware devices in a computer 44

A National Instruments product


p. 48

Analysis of the energy spectra is done via a Matlab program. In the lower left side of Fig

26, an example of a typical two-dimensional energy spectrum is shown. The elliptical region

corresponds to events where both the photons have an energy of E1 = E2 511 kev. The

horizontal and vertical band extending from the central peak are produced by coincidences of

a 511 keV photon with a background photon. To obtain the corresponding one-dimensional

spectrum a projection onto the E1+E2 = 1.022 MeV axis is made.

Fig 26 - The MatLab program converts the 2-dimensional spectrum (left) into a 1-dimensional spectrum

(right) [53]

Fig 27 shows the CDBAR spectroscope as installed in a hot cell laboratory at SCK•CEN.

On the first picture one can see the box with electronic cards and the PC. On the second

picture one can see the lead container from which the sample holder is loaded into a position

in between the two HPGe detectors, which are mounted on top of their liquid nitrogen

cryostat dewars. In the foreground of the picture, one can see the PLC and pneumatics that

control the loading/ejection of the sample holder.


p. 49

Fig 27 - CDBAR spectroscope as installed at SCK•CEN [53]


p. 50

A fraction of the positrons emitted by 22

Na can annihilate in the supporting Kapton foil

before they reach the material or they can be backscattered at the sample surface and

annihilate in the Kapton foil. This fraction depends on the material and can be as high as 15%

for high Z materials as e.g. Ni, and the Kapton contribution can thus be significant in the

Doppler spectrum [57]. Therefore a coincidence measurement was performed on a Kapton

sample and this contribution is subtracted from the experimental spectra.

The resolution of the system is an important parameter as it has been shown that the

position of the peaks in the ratio curves depends on the resolution of the system [58].

Therefore it is important to always mention the resolution of the system. The resolution of the

sum energy spectrum is given by the convolution of the resolution of the two single spectra.

In our case we have a symmetrical resolution function that can be described reasonably well

by a Gaussian function so that the resolution is given by σ = with σ1 and σ2 the

resolution of each detector at 511 keV. The two annihilation photons have Doppler shifts

which are equal but of opposite sign. Thus the observed Doppler effect is twice as large as in

a single detector spectrum. This results in an improvement of the effective resolution to σ/2.

For identical detectors this means an improvement with a factor of . In our setup we thus

have an effective resolution (FWHM) of 0.9 keV.

To check the stability of the system during normal measurements, a 57

Co source is added

to the setup. 57

Co is chosen to avoid a Compton scattering contribution to the 511 keV peak.

Indeed, the photon energies emitted by this source are smaller than 511 keV, i.e. respectively

122 keV and 136 keV. These peaks are used in combination with the 511 keV peak to

monitor the calibration and resolution every hour of the measurement.


p. 51

3.4. Experimental results

The CDBAR spectroscopy data have been processed into the following formats :

the absolute and relative values of the S and W parameter in table format

the S and W parameter versus the annealing temperature

the S parameter versus W parameter for the different annealing temperatures

the ratio curves (referenced to pure Fe) for the different annealing temperatures

The data as set out in the figures look consistent. The absolute S and W parameters were

-as described in paragraph 2.4.7.- respectively calculated as the ratio of the low momentum

region (0 < pL < 2.5∙10-3

m0c) in the CDBAR spectrum to the total region and as the ratio of

the high momentum region (15∙10-3

m0c < pL < 25∙10-3

m0c) in the CDBAR spectrum to the

total region. These momentum regions correspond respectively with a low energy region

(511 keV < Eγ < 512.3 keV) and a high energy region (518.7 keV < Eγ < 523.8 keV). These

ratio values are then subsequently normalized to the ones of pure, defect-free iron in order to

obtain the relative S and W parameter values. Both the absolute and relative S and W values

are given in table 2 and 3.

In figures 28, 32 and 36, the S and W values are plotted against their respective annea-

ling temperature. Remark here that the low momentum or S parameter, related to the presen-

ce of vacancies, shows a decline around 400°C and that the high momentum or W parameter,

related to the presence of Cu precipitates, shows a steep decline around 650°C. These two

clear and separate trends confirm the adequate choice made in determining the momentum

energy limits for the S and W parameters.

In figures 29, 33 and 37, the S values are plotted against their respective W values for the

different annealing temperatures. By this way, a -for model Fe-Cu alloys- typical, hooked

diagram is obtained, showing two more or less linear trends broken up at 550°C.

Figures 30, 31, 34 and 35 show the evolution of the momentum spectra over the

annealing temperature range by means of half ratio curves. Remark here the decline of the

low momentum (or vacancy-related) peak (pL < 5∙10-3

m0c) at around 400°C and its

subsequent stabilization. Remark also the rise and fall of the high momentum (or Cu


p. 52

precipitation related) peak (around pL = 25∙10-3

m0c) and its sudden disappearance at 650°C.

The differently colored lines represent different, applied methods of trend smoothing.

Table 2 - Absolute and relative values of S and W parameter for annealing of model alloy Fe-0.3wt%Cu ; 0.1 dpa

Table 3 - Absolute and relative values of S and W parameter for annealing of model alloy Fe-0.3wt%Cu ; 0.2 dpa

Sample D32 Anneal Temp

0.3 wt% Cu abs value rel value abs value rel value

0.1 dpa

D32-H-0 RT 0,396392 1,122315 0,013299 0,897658

D32-H-1 300°C 0,395873 1,120845 0,013797 0,931266

D32-H-2 350°C 0,394980 1,118340 0,014409 0,972596

D32-H-3 400°C 0,384406 1,088327 0,015681 1,058434

D32-H-4 450°C 0,368257 1,042622 0,018983 1,281301

D32-H-5 500°C 0,361318 1,022996 0,020941 1,413457

D32-H-6 550°C 0,359555 1,018002 0,021380 1,443102

D32-H-7 600°C 0,360263 1,020020 0,019632 1,325148

D32-H-8 650°C 0,360169 1,019740 0,014024 0,946588

D32-H-9 700°C 0,360341 1,020239 0,013944 0,941222

S-Parameter W-Parameter

Sample D42 Anneal Temp

0.3 wt% Cu abs value rel value abs value rel value

0.2 dpa

D42-H-0 RT 0,401328 1,136285 0,012895 0,870409

D42-H-1 300°C 0,400350 1,133503 0,013418 0,885725

D42-H-2 350°C 0,399453 1,130974 0,013307 0,898239

D42-H-3 400°C 0,388423 1,099747 0,014754 0,995856

D42-H-4 450°C 0,371780 1,052638 0,017671 1,192776

D42-H-5 500°C 0,363241 1,028454 0,020255 1,367179

D42-H-6 550°C 0,360723 1,021314 0,020990 1,416791

D42-H-7 600°C 0,360882 1,021760 0,019381 1,308206

D42-H-8 650°C 0,360181 1,019753 0,014075 0,950037

D42-H-9 700°C 0,359862 1,018864 0,014011 0,945711

S-Parameter W-Parameter


p. 53

Fig 28 - Evolution S and W parameter for annealing of model alloy Fe-0.3wt%Cu ; 0.1 dpa

Fig 29 - S and W plot chart for annealing of model alloy Fe-0.3wt%Cu ; 0.1 dpa

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

25°C 300°C 350°C 400°C 450°C 500°C 550°C 600°C 650°C 700°C

S an

d W

par

ame

ter

(rat

io t

o p

ure

iro

n)

Annealing temperature

Evolution S and W parameterSample D32 : Fe-0.3wt%Cu ; 0.1 dpa

S

W

RT300°C

350°C

400°C

450°C

500°C550°C

600°C

650°C700°C

Pure Cu

Pure Fe

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

1,000 1,050 1,100 1,150

W p

aram

ete

r(r

atio

to

pu

re ir

on

)

S parameter(ratio to pure iron)

S and W plotSample D32 : Fe-0.3wt% Cu ; 0.1 dpa


p. 54

[ pL in 10-3 m0c ]

Fig 30 - Ratio curves for annealing over temp range 300°C-450°C of model alloy Fe-0.3wt%Cu ; 0.1 dpa

0.5

1

1.5

2

S = 1.023 W = 1.4135

S = 1.023 W = 1.4135

500°C

0,5

1

1,5

S = 1.018 W = 1.4431

S = 1.018 W = 1.4431

550°C

0,5

1

1,5

S = 1.02 W = 1.3251

S = 1.02 W = 1.3251

600°C

0 5 10 15 20 25 30 350,5

1

1,5

S = 1.0197 W = 0.94659

S = 1.0197 W = 0.94659

650°C

0,5

1

1,5

2

S = 1.1223 W = 0.89804

S = 1.1223 W = 0.89804

25°C

0,5

1

1,5

S = 1.1208 W = 0.93127

S = 1.1208 W = 0.93127

300°C

0,5

1

1,5

S = 1.1183 W = 0.9726

S = 1.1183 W = 0.9726

350°C

0,5

1

1,5

S = 1.0883 W = 1.0584

S = 1.0883 W = 1.0584

400°C

0 5 10 15 20 25 30 350,5

1

1,5

S = 1.0426 W = 1.2813

S = 1.0426 W = 1.2813

450°C


p. 55

[ pL in 10-3 m0c ]


0.5

1

1.5

2

S = 1.023 W = 1.4135

S = 1.023 W = 1.4135

500°C

0,5

1

1,5

S = 1.018 W = 1.4431

S = 1.018 W = 1.4431

550°C

0,5

1

1,5

S = 1.02 W = 1.3251

S = 1.02 W = 1.3251

600°C

0,5

1

1,5

S = 1.0197 W = 0.94659

S = 1.0197 W = 0.94659

650°C

0 5 10 15 20 25 30 350.5

1

1.5

S = 1.0202 W = 0.94122

S = 1.0202 W = 0.94122

700°C


p. 56

Fig 32 - Evolution S and W parameter for annealing of model alloy Fe-0.3wt%Cu ; 0.2 dpa

Fig 33 - S and W plot chart for annealing of model alloy Fe-0.3wt%Cu ; 0.2 dpa

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

25°C 300°C 350°C 400°C 450°C 500°C 550°C 600°C 650°C 700°C

S an

d W

par

ame

ter

(rat

io t

o p

ure

iro

n)


Evolution S and W parameterSample D42 : Fe-0.3wt%Cu ; 0.2 dpa

S

W

RT300°C

350°C

400°C

450°C

500°C

550°C

600°C

650°C700°C

Pure Cu

Pure Fe

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

1,000 1,050 1,100 1,150

W p

aram

ete

r(r

atio

to

pu

re ir

on

)


S and W plotSample D42 : Fe-0.3wt%Cu ; 0.2 dpa


p. 57

[ pL in 10-3 m0c ]


0.5

1

1.5

2

S = 1.1363 W = 0.87041

S = 1.1363 W = 0.87041

25°C

0,5

1

1,5

S = 1.1335 W = 0.90572

S = 1.1335 W = 0.90572

300°C

0,5

1

1,5

S = 1.131 W = 0.89824

S = 1.131 W = 0.89824

350°C

0,5

1

1,5

S = 1.0997 W = 0.99586

S = 1.0997 W = 0.99586

400°C

0 5 10 15 20 25 30 350,5

1

1,5

S = 1.0526 W = 1.1928

S = 1.0526 W = 1.1928

450°C

0.5

1

1.5

2

S = 1.0285 W = 1.3672

S = 1.0285 W = 1.3672

500°C

0,5

1

1,5

S = 1.0213 W = 1.4168

S = 1.0213 W = 1.4168

550°C

0,5

1

1,5

S = 1.0218 W = 1.3082

S = 1.0218 W = 1.3082

600°C

0 5 10 15 20 25 30 350,5

1

1,5

S = 1.0198 W = 0.95004

S = 1.0198 W = 0.95004

650°C


p. 58

[ pL in 10-3 m0c ]


0.5

1

1.5

2

S = 1.0285 W = 1.3672

S = 1.0285 W = 1.3672

500°C

0,5

1

1,5

S = 1.0213 W = 1.4168

S = 1.0213 W = 1.4168

550°C

0,5

1

1,5

S = 1.0218 W = 1.3082

S = 1.0218 W = 1.3082

600°C

0.5

1

1.5

S = 1.0198 W = 0.95004

S = 1.0198 W = 0.95004

650°C

0 5 10 15 20 25 30 350.5

1

1.5

S = 1.0189 W = 0.94571

S = 1.0189 W = 0.94571

700°C


p. 59

Fig 36 - Evolution S and W parameter for annealing of model alloys Fe-0.3wt%Cu ; 0.1 and 0.2 dpa

Fig 37 - S and W plot chart for annealing of model alloys Fe-0.3wt%Cu ; 0.1 and 0.2 dpa

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

25°C 300°C350°C400°C450°C500°C550°C600°C650°C700°C

S an

d W

par

ame

ter

(rat

io t

o p

ure

iro

n)


Evolution S and W parameterSample D32 : Fe-0.3wt%Cu ; 0.1 dpa Sample D42 : Fe-0.3wt%Cu ; 0.2 dpa

S of D32

W of D32

S of D42

W of D42

RT

300°C350°C

400°C

450°C

500°C550°C

600°C

650°C700°C

400°C

450°C

500°C

Pure Cu

Pure Fe

0,800

0,900

1,000

1,100

1,200

1,300

1,400

1,500

1,600

1,000 1,050 1,100 1,150

W p

aram

ete

r(r

atio

to

pu

re ir

on

)


S and W plotSample D32 : Fe-0.3wt%Cu ; 0.1 dpa Sample D42 : Fe-0.3wt%Cu ; 0.2 dpa

D32

D42


p. 60

3.5. Interpretation of the results

The ratio curves probably provide the most information. For each annealing temperature,

two peaks are visible in the ratio spectra. There is a broad peak around 25∙10-3

m0c which

shows positron annihilations with core Cu electrons, namely the inner 3d orbital electrons of

Cu atoms. The smaller and lower peak in the low-momentum region (< 7∙10-3

m0c) shows

positron annihilations with valence electrons. The interpretation of these two peaks might be

somewhat confusing. As was stated in paragraph 2.2.1., vacancies in irradiated steel cannot

survive at room temperature unless they are in a clustered form and/or associated with solute

atoms. There was also mentioned that Cu precipitates are formed via both irradiation-

enhanced and irradiation-induced diffusion which finally leads to small vacancy clusters or

even microvoids (of 10 to 30 mono-vacancies) covered by Cu atoms. Paragraph 2.2.3. stated

that positrons can be trapped by both open-volume defects such as vacancies and by defect-

free ultrafine (nano-size) Cu precipitates. Bringing these statements together, one can

conclude that the peak in the low-momentum region must be due to positron trapping by

small vacancy clusters within Cu precipitates with subsequent positron annihilation with

valence electrons, whereas the peak in the high-momentum region must be due to positron

trapping by the copper precipitates with subsequent positron annihilation with the core 3d

orbital electrons of Cu atoms.

The evolution of the ratio curves (figures 30, 31, 34, 35) shows two distinct stages. A

first stage in the temperature range of 400°C to 500°C where the peak in the low-momentum

region almost disappears and a second stage in the temperature range of 600°C to 650°C

where the broad peak in the high-momentum suddenly completely disappears. The first stage

can be explained as the vacancies that break up and thermally dissociate from the Cu in the

Cu-vacancy clusters. The vacancies migrate as monovacancies in the bulk material and

disappear at sinks, thereby leaving the Cu behind as ultrafine (nm) defect-free Cu

precipitates. The second stage can be explained as the thermal dissociation of the Cu

precipitates that diffuse as individual Cu atoms in the bulk material and distribute evenly as

solute atoms, thereby bringing the material back to a nearly unirradiated state. Remark that at

annealing temperatures higher than the first stage, positrons are no longer trapped at

vacancies but are trapped at defect-free ultrafine Cu precipitates, being the second, competing

trapping mechanism for positrons.


p. 61

Due to the formation process (irradiation enhanced/induced diffusion), it is generally

accepted that the crystal lattice structure of the Cu precipitates is the bcc-structure of the α-Fe

host matrix. This means that at the location of the Cu precipitates there will be α-Fe crystals

of which part of the lattice atom positions are occupied by Cu atoms. It can also be expected

that the general phenomenon of coarsening or Ostwald ripening takes place whereby larger

Cu precipitates grow at the expense of smaller Cu precipitates and thus the number density of

the Cu precipitates decreases. Coarsening cannot be deducted from CDBAR results but the

question was raised whether coarsening would cause a lattice transformation from the bcc-

structure, typical for α-Fe, towards the fcc-structure, typical for Cu. Assistant Prof. Takeshi at

ONRDC (Japan) seems convinced that at 0.3wt% Cu in pure Fe the lattice structure will still

be a bcc-structure arguing that he has never experienced additional positron trapping after

coarsening which would be the case at the bcc-fcc matrix interface.

Reviewing the charts of the evolution of the S and W parameter versus the annealing

temperature (figures 28, 32), one might be tempted to interpret the steep increase in the W

parameter in the temperature range of 400°C to 550°C as due to coarsening which would

mean the enrichment in Cu of the larger Cu precipitates. It is however most probably simply

a relative increase of positrons annihilated at the Cu precipitates as the total number of

positron annihilations stays constant and the number of positron annihilations at vacancies

decreases.

The charts of the S,W-plane (figures 29, 33) show more or less two straight lines which

suggests two different mechanisms. The (S,W) point moves to the point close to that of pure

Cu along a straight line by annealing from 350°C to 500°C and then moves down along a

straight line to a point, probably representing the state of the unirradiated alloy, by annealing

from 600°C to 650°C. We see a first stage (400°C-500°C) with combined increase of the W

value and decrease of the S value which directs towards a relative increase of annihilations at

vacancies combined with a relative increase of annihilations at (ultrafine) Cu precipitates. A

second stage (600°C-650°C) shows a sharp decrease in the W value which points to an

important relative decrease of annihilations at (ultrafine) Cu precipitates. In general terms,

one can also conclude from these S,W plot charts that during the first stage the vacancies

disappear and during the second stage the Cu precipitates dissolve and diffuse into the bulk.


p. 62

There is very little to say about the difference in irradiation dose (0.1 vs 0.2 dpa) on the

results. In the merged S,W-plane (fig 37) one can discern an expected little shift to the right

or to higher S values for the higher irradiated sample, as more vacancy-related defects will

probably have been created. With increasing annealing temperatures, both curves move to

each other and finally reach the same -nearly totally recovered- state at 650°C-700°C.

In order to show the exceptional behavior of Cu in forming precipitates in irradiated

dilute Fe alloys, an S,W-plane chart (fig 38) is included from a study on isochronal annealing

of different model alloys by Nagai et al [59]. The samples were irradiated with 3 MeV

electrons at 50°C to a fluence of 2.0∙1018

e cm-2

. Only the Fe-Cu sample shows precipitation

formation. In case of the other alloys, the (S,W) points move to the point of nearly totally

recovered state along straight lines directly. This suggests that the irradiation effects

completely recover by annihilation of the vacancies and vacancy clusters without forming

any precipitate that could trap positrons.

Fig 38 - S,W-plane chart for different irradiated binary model alloys [59]

p. 63

Chapter 4 : Conclusions

The experimental results seem to confirm the theory of Prof. Nagai on formation of Cu-

precipitates by neutron irradiation and deformation by annealing in RPV steels [18], namely

that :

vacancies introduced by neutron irradiation migrate in Fe and preferentially connect

with Cu-atoms in order to form Cu-vacancy complexes.

these complexes aggregate (10 to 30 ea) to form micro-voids, which can be seen as

voids of which the inner surface is decorated with Cu atoms.

at 400°C, the vacancies in the Cu-vacancy complexes dissociate from the Cu-atoms

and leave the copper behind as ultrafine (1nm) defect-free Cu-precipitates.

at 650°C, the Cu-precipitates break up into individual atoms.

The confirmation of my experiments with this theory lies in the fact that my results show

clearly the two temperature stages, namely

the 400-450°C stage during which the vacancies disappear and the Cu-rich positron

annihilation sites move closer to pure Cu.

the 600-650°C stage during which the Cu-precipitates dissolve and thermally diffuse

into the bulk.

We know, also from a publication of Prof. Nagai that positrons can annihilate at nano-

size Cu-precipitates [60]. If the Cu-precipitates would grow larger, their crystal lattice

structure would evolve from the bcc-structure of α-Fe, in which the Cu-atoms are embedded,

to the fcc-structure of pure Cu and vacancies would be attracted to the matrix change

boundaries. The experimental results do not show an increase of positron annihilations

related to vacancies in the temperature range in which the Cu-precipitates are growing

towards pure Cu. This means that the Cu-precipitates are still nano-size and the Cu-atoms are

still within the bcc crystal lattice structure, typical for the α-iron.

The understanding at the atomic level of annealing phenomena might be helpful in the

context of larger annealing projects on RPVs. Especially the optimum annealing temperature

Chapter 4 : Conclusions

p. 64

could be determined this way. Further PALS and CDBAR research on isochronal annealing

of irradiated RPV steels could be done via other binary or ternary model alloys and later on

compared with identical results on real RPV specimens to see if all phenomena are well

understood. Recovery rates on the real RPV specimens could then be confirmed by CVN

tests. It would certainly have value to combine these experiments with simultaneous Vickers

microhardness measurements. In an even further stadium, reirradiation and reannealing

studies could be performed. By this way, a unique set of annealing data could be brought

together.

In the context of optimum annealing temperature for thermal treatment of real RPVs, it

might be noticed that the temperature range at which the vacancies disappear (in our case

400-450°C) is more adequate then the higher temperature range of dissolution of the Cu-

precipitates (in our case 600-650°C). Reasons for this is a study at ORNL showing already

very good ductility recovery (about 80%) at 454°C and the danger at higher temperatures for

creep deformation, temper embrittlement and dimensional changes due to thermal stresses.

p. 65

Appendix A

Main ferritic materials for Western RPV components [1]

p. 66

Appendix B

Annealing of VVER-440 type reactors [34]

Reactor Year Temp/time Country Shutdown

Novovoronezh 3

Armenia 1

Greifswald 1 (Nord 1)

Kola 1

Kola 2

Kozloduy 1

Kozloduy 3



Novovoronezh 3 *

Kozloduy 2

Bohunice V-1 / 2

Bohunice V-1 / 1

Loviisa 1

1987

1988

1988

1989

1989

1989

1989

1990

1990

1991

1992

1993

1993

1996

430°C/150h

450°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/150h

475°C/160h

475°C/168h

475°C/100h

Russia

Armenia

(East Germany)

Russia

Russia

Bulgaria

Bulgaria

(East Germany)

(East Germany)

Russia

Bulgaria

Slovakia

Slovakia

Finland

(2016)

1989

1995

(2018)

(2019)

2003

2006

1995

1995

(2016)

2003

2008

2006

(2027)

* reannealing

p. 67

Appendix C

US regulatory documents concerning thermal annealing of RPVs [1]

USNRC Thermal Annealing Rule : 10 CFR 50.66 : ‘Requirements for thermal annealing of

the reactor pressure vessel’, 1999

USNRC Regulatory Guide 1.162, ‘Format and Content of Report for Thermal Annealing of

Reactor Pressure Vessels’, Feb 1996

ASME Code Case N-557, ‘In-Place Dry Annealing of a PWR Nuclear Reactor Vessel’, Mar

1996

ASTM E 509-86, ‘Standard Guide for In-Service Annealing of Light-Water Cooled Nuclear

Reactor Vessels’, 2003

p. 68

Appendix D

Working procedure annealing

Materials : - samples

- small plastic bottles with cover for samples

- hot cell feed box (HCFB)

Hot cell : switch light on and put alarm timer45

on maximum

set temperature of oven to Tset + 10°C

start temperature recorder

Part 1

Put samples in small plastic bottle and bring into the HCFB

Attach the HCFB to the hot cell

Open access port to HCFB with the pneumatic switch

Open the access port with the tele-manipulators (TMs) and bring the material from the HCFB

into the hot cell

Close the access port with the TMs and secure the port via the pneumatic switch

Part 2

Transfer the samples from the plastic bottle to the steel basket (with the TMs)

Place the steel basket containing the samples on the sample holder

Move the quartz tube over the sample holder

Put the electromagnet on via the red push button (‘goed aansluiten vacuumzone’) on the

control panel in order to assure good connection of the quartz tube to the sample holder head

Check if mechanical vacuum pump is on (see if plug is in the socket)

Open the suction valve of the mechanical vacuum pump

From the moment there is vacuum, switch off the electromagnet

Switch on the turbo vacuum pump

(switch read-out vacuum pressure at 5 mbar to a lower range)

Let vacuum run up to 10-6

bara

(check if valves on helium tanks are open)

Switch the turbo vacuum pump off

Now vacuum will automatically decrease by injection of helium at 950 rpm for 2 minutes

(this is programmed in the vacuum controller)

After these 2 minutes, close the suction valve of the mechanical vacuum pump when 5 mbara

is reached. Due to the slowness of closing the suction valve, there should be around 1 mbara

left in the quartz tube then.

Part 3

If oven is at set temperature, move the oven over the quartz tube

From T = Tset – 10°C (read from (T1+T2)/2)46

, start the annealing time (for 30 minutes)

45

Gives alarm in the control room in case of air movement inside the hot cell with timer has timed out 46

The holder for the sample metal basket is equipped with 3 thermocouples, T1, T2 and T3. The two thermocouples closest to the metal basket are being used.

Appendix D

p. 69

After these 30 minutes, move the oven away from the quartz tube and switch off the oven.

Let cool down the sample holder to 100°C

Put the electromagnet back on

Break the vacuum by turning the switch button ‘beluchting’ (aeration) on the control panel

Move the quartz tube away

Bring the steel basket containing the samples from the sample holder to the hot cell table by

means of the TMs

Bring the samples from the sample basket in the plastic bottle and close the bottle with the

cover by means of the TMs

Part 4

Proceed further in reverse order of part 1 to bring the plastic bottle outside the hot cell

Part 5

Switch off the temperature recorder

Close the valves on the helium tank

p. 70

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