Using Microscopy to Help with the Understanding of ... · Degradation Mechanisms Observed in...

Journal of Materials Science and Engineering B 7 (9-10) (2017) 187-220 doi: 10.17265/2161-6221/2017.9-10.002

Using Microscopy to Help with the Understanding of

Degradation Mechanisms Observed in Materials of

Pressurized Water Reactors

Laurent Legras1, Alexandre Volgin1,2, Bertrand Radiguet2, Philippe Pareige2, Cedric Pokor1, Brigitte Décamps3,

Thierry Couvant1, Nicolas Huin1,4 and Romain Soulas1,5

1. EDF R&D-EDF Lab Les Renardières, Moret sur Loing cedex 77818, France

2. GPM Rouen, Avenue de l'Université, Saint-Étienne-du-Rouvray 76800, France

3. CSNSM IN2P3, Université de Paris Sud, Orsay 91400, France

4. Institut PPrime Université de Poitiers, ISAE-ENSMA, UPR 3346 BP 30179 F86962 Futuroscope Chasseneuil, France

5. INP Grenoble, SIMAP, Saint-Martin-d'Hères 38402, France

Abstract: Even if temperature, pressure and chemistry of the cooling water are not very high and aggressive, materials used in PWRs (Pressurized Water Reactors) are exposed to different degradation mechanisms. One of the main goals of the research programs in this field is to develop physical model of the mechanisms down to the atomic scale. Such approach needs a clear description and understanding of the degradation mechanisms at the same small scale. This paper illustrates the benefits of different microscopies and of their last improvements up to the promising possibilities of monochromated and aberrations corrected TEM/STEM. A specific focus is placed on four different degradation mechanisms observed in austenitic stainless steel: irradiation ageing, corrosion fatigue, stress corrosion cracking and corrosion.

Key words: PWR, TEM, APT, 3D SEM, irradiation ageing, corrosion fatigue, stress corrosion cracking, corrosion.

Even if temperature, pressure and chemistry of the 1 cooling water are not very high and aggressive,

materials used in PWRs (Pressurized Water Reactors)

are exposed to different degradation mechanisms.

Some of these materials (internal components,

vessel...) are also exposed to irradiation, leading to

complex degradation such as IASCC (Irradiation

Assisted Stress Corrosion Cracking). One of the main

goals of the research programs in this field is to

develop physical model of the mechanisms down to

the atomic scale. Such approach needs a clear

description and understanding of the degradation

mechanisms at the same small scale.

Corresponding author: Laurent Legras, Dr., research fields:

material science, transmission electron microscopy, metallurgy, degradation mechanisms observed in materials of pressurized water reactors: corrosion, stress corrosion cracking, irradiation assisted stress corrosion cracking and fatigue corrosion.

This paper illustrates the benefits of different

microscopies and of their last improvements

(TEM—Transmission Electron Microscopy,

APT—Atom Probe Tomography, dual beam

microscopes, EBSD—Electron Back Scattered

Diffraction...) up to the very new promising

possibilities of monochromated and aberrations

corrected TEM. A specific focus will be placed on

four different degradation mechanisms observed in

austenitic stainless steel: irradiation ageing, corrosion

fatigue, stress corrosion cracking and corrosion.

(1) Irradiation ageing

The internal structures of the PWR vessel are made

in austenitic stainless steel. Irradiation continuously

induces point defects which can recombine together,

and be eliminated on sinks or aggregated. Interstitials

aggregate forming disk shape i.e. dislocations loops or

D DAVID PUBLISHING

Using Microscopy to Help with the Understanding of Degradation Mechanisms Observed in Materials of Pressurized Water Reactors

188

segregate on grain boundaries. Vacancies could lead

to the formation of cavities or bubbles when

combining with gas atoms such as H or He. In

addition the coupling between point defect fluxes

towards sinks and solute atom diffusion can lead to

solute aggregation at sinks. Kinetics studies and

quantification of these phenomena are performed

using TEM and APT and are used as key data for the

understanding of mechanisms and for modelling

lifetime extension of actual PWRs. Ni and Si

segregations and Cr depletion on dislocation loops

were recently showed owing to new developments of

APT leading to better understanding of their

hardening effect.

(2) SCC (stress corrosion cracking)

Strain gradients and preferential sites of

intergranular SCC initiation have recently appeared to

be correlated. Combining complex strain on specimen

pre-mapped and the possibility of extracting TEM thin

foils on a precise location using dual beam offer a

promising possibility to understand fracture

mechanisms ahead of the crack tips. Current TEM

studies on such specimen are trying to understand the

stability of oxides films at the cracks tips as a function

of crystallographic orientation and strain. The very

first results exhibited that diffusion of Cr and Fe is

very fast at grain boundaries and accelerated by strain.

What is more the higher the strain was, the higher the

oxide penetration was, and thus the more SCC

degradation was favored.

(3) Corrosion fatigue

Specific areas of primary circuit in PWRs can be

affected by thermal fatigue. Recent TEM studies on

fatigue crack tips clearly showed that the effect of

oxidation is combined with the fatigue loading and

thus it has to be taken into account. 3D images with

dual beam coupled with TEM observations clearly

showed that oxide penetration along shear bands is

about 1.5 higher than in the matrix. Another result is

that oxide thickness measured after cyclic loading is

more than ten times higher than after a 2,000 h

monocyclic loading in SCC test.

(4) Corrosion

SCC corrosion fatigue and corrosion behaviour are

controlled by the properties of the oxide films. The

well known duplex structure of the oxides is still

suffering from uncertainties such as the very first step

of their growth and their crystallographic structure.

These two questions are currently under investigation

up to the atomic scale on very short exposure time.

HR(S)TEM (High Resolution (Scanning)

Transmission Electron Microscopy) and ASTAR

software clearly showed that the inner chromium rich

oxide is a (FeCr)3O4 spinel oxide growing with a

cube/cube epitaxial crystallographic relationship with

the base metal. HREELS (High Resolution Electron

Energy Loss Spectroscopy) also clearly showed the

possibility of quantifying the Fe2+/Fe3+ ratio in the

oxide and thus to get new informations on the passive

property of the Cr rich film layer depending if this

spinel is normal, intermediate or inversed ones.

1. General Introduction

Materials used in PWR are exposed to different

degradation mechanisms even if temperature, pressure

and chemistry of the cooling water are not very high

and aggressive. Some of the core components (internal

components, vessel...) are also exposed to irradiation

leading to complex ageing mechanisms such as

IGASCC (Irradiation Assisted Stress Corrosion

Cracking). This degradation involves a coupling

between irradiation ageing, external stress and

corrosion. SCC (Stress Corrosion Cracking) seems to

be generic and has a large impact on safety because it

occurs in the second containment barrier. Thermal

fatigue (and the possibility of a coupling effect with

PWR environment) is considered as a potential

degradation mechanism inside specific areas (surge

line, steam generator feedwater lines, hot and cold

water mixing areas…). Corrosion related degradation

of PWR materials has a significant impact on plant

availability and maintenance costs.


189

One of the main goals of the research programs

launched in these fields is to develop physical

modeling of the mechanism involved. Such approach

needs a clear description and understanding of the

ageing mechanism at the same scale that the modeling

one and thus various characterization techniques: (1)

Irradiation ageing modeling needs to know the

macroscopic evolution of irradiation defects but also

their exact chemistry and nature at sub nanometer

scale. (2) Modeling of SCC and of thermal fatigue

needs a good knowledge of the strain heterogeneity as

a function of microstructure gradient from the

component down to the grain scale but also a fine

description of the local composition, thickness and

structure at nanometer scale. (3) Properties of oxide

films are a crucial point which needs macroscopic

kinetic measurement but also a good knowledge of

their chemistry, structure and growth mechanism from

the very beginning of the oxidation process and thus

down to the atomic scale.

Since all these degradation mechanisms are

affecting numerous materials (stainless steels, nickel

based alloys, zirconium alloys...) we will illustrate the

benefit of some characterization techniques up to the

very promising possibilities of monochromated and

aberrations corrected HR(S)TEM in some undergoing

research programs on degradations mechanism

affecting austenitic stainless steels.

2. Irradiation Ageing of 316 L Austenitic Stainless Steel

2.1 State of the Art

The internal structures of PWR located close to the

reactor core are used to support the fuel assemblies, to

maintain the alignment between assemblies and the

control rods and to channel the primary water (Fig. 1).

In general these internal structures consist of baffle

plates in Solution Annealed 304 stainless steel and

baffle bolts in Cold Worked 316 stainless steel. These

components undergo a large neutron flux (up to 100

displacements per atom—dpa) at temperatures

between 280 and 380 °C. As a result, the material

exhibits a substantial increase in yield stress (Fig. 2),

reduction in ductility, swelling loss of corrosion

resistance which may be damaging to the proper

operation of the reactor [1-7]. For instance cracks

were observed in bolts (Fig. 3) [8-11]. These cracks

usually attributed to IASCC (Irradiation Assisted

Stress Corrosion Cracking), can be seen as a

consequence of the evolution of plasticity in these

materials loaded in a corrosive medium, together with

possible evolutions of grain boundary chemistry (due

to RIS—Radiation Induced Segregation) [12].

Fig. 1 (a) Schematic drawing of a PWR vessel and (b) schematic drawing of internal structures.

(a) (b)


190

Fig. 2 Hardening as a function of neutron irradiation dose on a cold worked 316 stainless steel [2].

50 μm

(a) (b)

Fig. 3 (a) Example of a crack observed of an internal bold; (b) SEM image showing the intergranular propagation of the crack.

Since evolutions of the macroscopic properties are

related to the microstructure ones, understanding

microstructure evolution under irradiation is essential

to predict lifetime of internals. A large number of

studies were focused on the characterization of the

microstructure under irradiation. On one hand, it was

shown that RIS at GBs (grain boundaries) can

contribute to the loss of corrosion resistance and could

be implicated in the IASCC [13-16]. Indeed, under

irradiation, a permanent flux of PD (point defects) is

produced towards sinks like GB. As coupling is

created between PD and solutes fluxes, it results in a

Cr depletion and an enrichment in Ni at GB (Fig. 4a).

On the other hand, TEM (transmission electron

microscopy) observations (Figs. 4b and 4c) have

shown that radiation-hardening could be caused by the

formation of a high number density of Frank loops

and PD clusters [17-21]. However, to ensure that the

yield stress evolution is only due to the Frank loops

formation, a chemical characterization of the

microstructure was performed. γ' precipitation [20, 22]

or Ni-Si enrichment at Frank loops [23] observations

using conventional TEM are reported in the literature.

This was completed more recently by A. Etienne

owing to the development of the LATAP (Laser

Assisted Tomographic Atom Probe) at the GPM


191

(Groupe de Physique des Matériaux) in the University

of Rouen (Fig. 4d [24]). The LATAP allows the

analysis of such brittle materials which were very

difficult with conventional TAP (using electric pulse).

After neutron irradiation, she observed heterogeneous

distribution of solute atoms, and the presence of two

kinds of features, Si atmospheres and Ni-Si enriched

clusters were shown in Ref. [24]. The comparison

between LATAP analysis and TEM investigation of Fe

irradiated samples suggested that Ni-Si rich clusters

Fig. 4 Examples of images and chemical analysis performed on irradiated 316 L. (a) Cavities and γ' precipitation [2].

Fig. 4 Examples of images and chemical analysis performed on irradiated 316 L. (b) Frank loops [2].


192

Fig. 4 Examples of images and chemical analysis performed on irradiated 316 L. (c) EDXS line scan analysis using FEG STEM on a grain boundary of an irradiated 316 L [16].

Fe

Cr

Ni

60nm

Fig. 4 Examples of images and chemical analysis performed on irradiated 316 L. (d) 3D reconstruction of a volume of 316 L irradiated by neutron showing Ni-Si clusters [24].


193

could be formed by induced degradation on

dislocation loops [25, 26]. But at that time, this

formation mechanism still needed to be more detailed.

The aim of this part is to illustrate the benefit of

coupling TEM observations and APT analysis rather

than explaining the clustering mechanism observed

under irradiation. This is done through presenting

some results obtained during the PhD work of A.

Volgin [27] and presented at TMS 2012 conference.

This PhD work is performed under the European

project PERFORM 60.

2.2 Results & Discussion

Specimens were extracted from a 5% cold worked

316 SS annealed 10 minutes at 1,100 °C and from a

ternary hot rolled FeCrNi alloy annealed 1 h at

1,100 °C. After annealing both samples were water

quenched. The bulk compositions of these alloys are

given in Table 1. Then 5 MeV Ni++ irradiations were

performed at Orsay JANNuS platform at 450 °C. The

temperature shift with respect to PWR conditions is

necessary for compensating flux effects [28-32]. The

cumulative dose at the depth where samples are

prepared is 5 dpa. TEM observations were done at

EDF R&D on a TECNAI G2 20F. Irradiated atom

probe samples were prepared from 3 millimeter disks

using dual beam at GPM in the University of Rouen.

Chemical characterization at the atomic scale was

performed using the Cameca LAWATAP (Wide

Angle LATAP) at GPM. The principle of this

instrument is described in Refs. [33, 34]. Analyses

were performed at 50 K, using femtosecond laser

pulses with a wavelength equal to 330 nm and a pulse

energy resulting in the same effect as that of an

electric pulse equal to 20%-25% of the standing

voltage.

The microstructure of the irradiated specimens

consists in interstitial Frank and perfect loops (Fig. 5a),

in vacancy clusters in the form of small voids (Fig. 5b)

and Stacking Fault Tetrahedra (Fig. 5c). These

observations were also done to measure the number

density and the size of these features and thus, to get

an estimation of their quantitative hardening effect.

Finally, from conventional bright and dark field TEM

observations, no clear evidence of precipitation is

observed (no chemical analyses using TEM were

performed). To clarify this point, LAWATAP

analyses were performed on irradiated samples. These

analyses clearly evidence Ni, Si segregation on

dislocation loops whereas Cr depletion is observed at

the same times (Fig. 6). Some Ni, Si enriched clusters

that are also pointed out. Because of their small size

and density, all these objects would be very difficult

to be detected using EDX or EELS analyses in TEM.

It is certainly more difficult, if not impossible, to

quantify precisely the amount of the Ni, Si enrichment

and Cr depletion using such chemical analysis in TEM.

Thus, the exact nature of the objects on which

segregation and depletion are observed is done using a

comparison between TEM images obtained from one

side and from APT analysis on the other side. Even if

the link can be made comparing the size, density,

morphology of these objects, further HAADF

HRSTEM images and EELS analyses using probe

corrected and monochromated TEM will be performed

to try to fill this gap. In addition, the same APT thin

needles will be analyzed by conventional TEM

(crystal defects) HAADF HRSTEM and APT.

The authors acknowledge Orsay JANNuS Team for

performing irradiations.

Table 1 Bulk compositions of ternary model alloy and 316 austenitic stainless steel. Concentrations are in wt. %. Balance is iron.

Alloy\Element (%wt) Cr Ni Si C N Mo Cu P S

316 L 17.60 12.00 0.59 0.028 0.08 2.39 0.066 0.01 0.01

FeNiCr 1818.27 12.40 - 0.008 - - - - 0.003


194

Fig. 5 Example of defects observed in irradiated 316 L and FeNiCr alloy. (a) Example of 1 from 4 Frank loop families.

Fig. 5 Example of defects observed in irradiated 316 L and FeNiCr alloy. (b) Cavities: Underfocused bright field TEM image.

Fig. 5 Example of defects observed in irradiated 316 L and FeNiCr alloy. (c) Example of Stacking Fault Tetrahedra.

20 nm


195

Fig. 6 LAWATAP analyses preformed on irradiated samples. (a) FeCrNi synthetic alloy.

Fig. 6 LAWATAP analyses preformed on irradiated samples. (b) 316 L.

3. SCC of 316 L Austenitic Stainless Steel


Predicting the time to initiation of SCC in austenitic

materials exposed to the primary circuit of PWRs is of

major interest for optimizing the maintenance strategy

of reactors and evaluating the best way to improve

manufacturing, mitigation and repair of components.

An engineering model was developed to predict

initiation of IGSCC in nickel based Alloy 600

exposed to primary water [35]. This empirical model

does not describe physical mechanisms, although it

includes variables that most often have a physical

meaning (temperature, stress, carbide precipitation).

The time to initiation was correlated with a stress

index, a temperature index and a material index

(assumed to depend on the intra and intergranular

precipitation [36, 37]) representative of the intrinsic

susceptibility to SCC initiation. Such an engineering

model still has to be calibrated for stainless steels.

Moreover, the existing model suffers from a lack of

accuracy. The stress index does not consider the

deleterious effect of reverse loading on the

susceptibility to cracking, and the material index is not

Ni

83.6 nm

Ni Cr27 nm

NiNi

83.6 nm

Ni Cr27 nm

Ni segregation and Cr depletion on dislocation loops

Ni/Si segregation and Cr depletion on dislocation loops


196

parameterized by the limiting microstructural factors

driving corrosion (local material chemistry, transport

of oxygen and ions within the oxide and the metal).

Due to a lack of understanding, the calibration of the

model is not accurate.

Field experience highlighted the deleterious effect

of strain hardening on SCC (Fig. 7) in nominal

primary water. Pressurizer heaters [38, 39], made of

stainless steels, and steam generator divider plates

[40-42], made of rolled Alloy 600 were both affected

by SCC. In each case, the level of deformation and

heterogeneity at a macroscale, captured by micro

hardness or neutron diffraction, was believed to play a

major role in the occurrence, location, and kinetics of

SCC [43, 44]. At a lower scale of strain localization,

scanning and transmission electron microcopy

observations suggested possible deleterious local

interaction between oxidation and mechanical fields

[45]. Fig. 8 shows precursors to IGSCC aligned with

emerging shear bands at the surface of 304 L exposed

to a primary environment, while Fig. 9 exhibits the

presence of shear bands intersecting a grain boundary

where SCC propagates.

Understanding and quantification of interactions

between localized deformation in materials and EAC

(environmental-assisted cracking) could play an

important role in maintaining the integrity of LWR

components. Thus, a detailed understanding of strain

Fig. 7 IGSCC initiation at the surface of a heater of pressurizer.

Fig. 8 SEM image of IGSCC initiation sites lined with slip bands intersecting the grain boundary on 304 L in a nominal hydrogenated primary PWR environment at 360 °C [33].

30 µm


197

Fig. 9 TEM image of intergranular propagation assisted by localized deformation. Crack growth test in a PWR environment on a 40% cold rolled 304 L [33].

Fig. 10 Back face of the cross specimen.

localization during plastic deformation and of the

underlying mechanisms is important for the

manufacture, design, and repair of materials exposed

to the environment of the PWRs primary circuit.

3.2 Results and Discussion

The deleterious effect of cross tests on

susceptibility of 304 L to IGSCC in PWR

environments (1,000 ppm B, 2 ppm Li, 20 cc H2) at

360 °C was shown in previous studies [46]. This result

has been interpreted by the authors to be a

consequence of the strain localization induced by

orthogonal strain paths as defined by Schmitt [47].

Cross tests are used to promote strain localization. The

cross specimen (Fig. 10) was specially designed. Only

one face is perfectly polished for material

characterization, and local strain evaluation.

Grids which were deposed at LPMTM2 (Paris XIII

University, Galileo Institute) by electronic micro

lithography, at the center of the specimens before

testing to allow quantification of superficial strain

localization after pre-strain hardening in air and

followed by deformation in the primary PWR

environment at 360 °C. Gridline images were

collected before and after each step of deformation

2 Laboratoire des Propriétés Mécaniques et Thermodynamiques des Matériaux (Villetaneuse, France).

Prior strain (air) axis X

Axis Y

for testing in primary water

Shear Bands


198

using SEM (scanning electron microscopy).

Comparison of the grid before and after (Fig. 11)

deformation revealed the material superficial

displacement field [48].

Tools were developed (Matlab) in order to do cross

correlations between microstructural parameters such

as grain orientation and grain boundary as collected by

EBSD and mechanical fields based on DIC of strain

and strain path. Using this method, distribution can be

evaluated and compared and the local deformation

path known. Thin foils with well known

crystallographic orientation, strain level, and strain

path were extracted with Dual Beam. Then, the effect

of the local strain and path influence on the oxide

growth and on cracks initiation and growth were

studied by TEM. Figs. 12 and 13 show that the oxides

Fig. 11 (a) Grid before deformation. Initial mesh size = 5 µm.

Fig. 11 (b) Grid after deformation in ambient air. Initial mesh size = 5 µm.


199

STEM HAADF Image obtained on a cross-section 316 L site 1

EDX line scan profile 1



Fig. 12 STEM HAADF Image and EDX line scan profile cross the oxides layers (at surface and in intergranular crack) on a cross section of 316 L site 1 (Prior-Exx = 0, second Eyy = 0.03, t = 720 h, pH320 °C = 7.2).

1,5µm

EDX profile 1 EDX profile 2

EDX profile 3


200

ZLP image O filtered image Cr filtered image Fe filtered image Ni filtered image

Fig. 13 EFTEM images obtained from the intergranular crack observed on Figure (cross section of 316 L site 1 (Prior-Exx = 0, second Eyy = 0.03, t = 720 h, pH320 °C = 7.2)).

formed at the surface of specimen are the same than

those formed inside the cracks. Some other

conclusions can be drawn from such observations: the

oxide thickness is clearly affected by strain and oxides

are growing faster along activated dislocation plans

leading to a very rough metal oxide interface at the

nanometer scale (Fig. 13). It is also obvious that the

diffusions of species such as Cr and Fe are very fast at

grain boundaries and clearly accelerated by strain.

Table 2 summarizes some of the measurements made

on this specimen using image analysis on the Cr

energy filtered images to evaluate more precisely the

effect of strain on oxide growth and SCC cracks

growth. It shows that the more pronounced the strain,

the higher the penetration of oxide and thus the more

SCC degradation is favored. All these data are now

used to develop a physical modeling of IGSCC on

stainless steels alloys.

4. Fatigue Corrosion of 304 L Austenitic Stainless Steel


Despite the fact that many mechanical data on the

response of 304 stainless steel to cyclic loading are

available, the current fatigue codification suffers from

a lack of understanding of the environmental effects

(more precisely for high number of cycles and for low

amplitude of loading) [49-51]. The strain rate effect at

the early stage of fatigue life in PWR primary

environment of a 304 L stainless steel was recently

studied [57]. Some of the results are plotted in Fig. 14.

It shows a clear deleterious effect of environment on

fatigue time life. Whereas the microstructural

evolution of 316 L under cyclic loading in air was

quite well described by Refs. [52-60], very few data

were available on 304 L until the PhD Work of A.

Garcia [61]. She mapped the microstructure evolution

after failure (Fig. 15). In her work, orientation

mapping facilities ACT [62] and ASTAR [63] were

very helpful for finding without ambiguity that highly

strained small precipitates of α’-martensite nucleated

at the intersection of Corduroy [64, 65] systems.

Based on this microstructural description of fatigue

mechanism in air and in water, A. Garcia concluded

that the environment is not expected to affect the

deformation mechanisms in the bulk of the specimens.

In addition, she showed using TEM cross section

observations, that at 300 °C in primary water, oxides

are formed up to the cracks tips. Thus, there are clear

indications that environment may promote the

initiation of fatigue cracks in 304 L stainless steel.

From the mid-nineties, the development of “Dual

Beam” systems, (SEM—Scanning Electron

Microscope, equipped with a Focused Ion Beam-FIB),

gave the possibility to make cross-section, TEM

preparation, 3D images, 3D EBSD and chemical

mapping from an area of interest with a precision

directly linked to SEM and FIB resolution [66]. This

offers new helpful tools for applying a quantitative

microscopic approach when studying effects of

environment on fatigue crack initiation and

propagation. Such multi scale observations of oxide

morphology and microstructure are carried out in the

PhD work of N. Huin [67] and will be briefly illustrated


201

Table 1 Analysis of 3 cracks initiated at the surface of the 316 L specimen, 720 h, pH320 °C = 7.2).

Crack Flank Local strain in the

sample

Crack length (nm)

Oxide at the crack tip (nm)

Affected grain boundary length (nm)

Analyzed length (nm)

Cr-rich inner oxide penetration (nm)

Prior Exx Second Eyy Average Median Min. Max. SD

1 Left 0.05 0.05

899 59 170 786 88 89 48 119 16

Right 0.08 0.10 619 76 83 24 125 28

2 Left 0.00 0.02

548 136 136 304 67 66 50 89 9

Right 0.00 0.02 322 66 64 56 82 7

3 Left 0.1 0.01

317 110 125 123 43 43 33 52 5

Right 0.1 0.01 143 36 38 16 47 7

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Material Environment Temperature Strain rate

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As-received PWR

300 °C

0.4

0.004

AIR 0.4

Fig. 14 Environmental effect on 304 L fatigue life at 300 °C [49].

in the following.


Initial surface of materials plays an important role

in fatigue crack initiation. Preferential cracks

initiations are localized where local strains induced by

cycling are high. PSBs (Persistent Slip Bands) create a

surface roughness composed of intrusion and

extrusion (Fig. 16) which are precursors of material

damaging [68]. Crystallographic orientation of each

grain plays an important role in the process of

emerging PSB. Indeed, during a fatigue test, some

grains underwent low plastic strains while others were

strongly strained with mainly only one slip system

activated [69]. The Schmid factor (giving the

orientation of a gliding plane and its

direction/macroscopic applied stress) can be used to

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. 15

Mai

ns m

icro

stru

ctur

e fe

atur

es o

bser

ved

afte

r ru

ptu

re o

n 3

04 L

tes

ted

by

olig

ocyc

lic

load

ing

[61,

64]

.

Using Microscopy to Help with the Understanding of degradation mechanisms Observed in Materials of Pressurized Water Reactors

203

using the <110>{111} slip system, was performed and

results are presented in Fig. 17. These results show

that 90% of grains where cracks initiated have a

Schmid factor higher than 0.41.

A very clear environment effect is seen in Fig. 18.

At 300 °C oxides penetration and outgrowth are larger

in primary water than in air showing that the

environment is affecting crack initiation. Fig. 19

shows a cross-section of the oxide layer formed on the

surface of a sample tested at 300 °C in PWR primary

environment. Shear bands emerge at the surface of the

specimen. Some of them are preferentially oxidized as

highlighted with the red squares. Area 1 in Fig. 19

shows a crack initiation site at the emergence point of

a shear band. The crack is aligned with the

deformation substructure and presents a preferential

oxidation indicating that it seems to propagate along

the sub laying shear bands. Nevertheless, these

observations are not sufficient to determine whether

cracks initiated by mechanical failure and were

subsequently filled by oxide, or if they were affected

by oxidation. In Area 2 (Fig. 19), preferential

oxidation along shear bands before the crack started to

open can be observed suggesting that preferential

oxidation occurred before its opening. Fig. 19 also

shows that not all emerging shear bands are oxidized.

Fig. 16 Slip occurring at the surface, 316 L, 20 °C, ∆εp/2 = 2×10-3 [70].

Fig. 17 Distribution of Schmid factors of untracked grains (red) and grains presenting TG cracks, 316 L, 20 °C, ∆εp/2 = 2×10-3 [70].


204

Fig. 18 Fatigue crack initiation in a 304 L SS, 300 °C , air environment, Δεt/dt = 0.004%/s, Δεt/2 = 0.5%.(STEM image obtained in the dual beam), (a): in air.

Fig. 18 Fatigue crack initiation in a 304 L SS, 300 °C , air environment, Δεt/dt = 0.004%/s, Δεt/2 = 0.5%.(STEM image obtained in the dual beam), (b): in PWR environment.

Pt coating

304L foil

Oxide

PSB

PSB

Pt coating

304L foil

Oxide


205

(a)

(b)

Fig. 19 FIB cross section of a sample tested at Δεt/dt = 0.004%/s, Δεt/2 = 0.5%, 300 °C in PWR primary environment observed at 5 kV, Secondary electron images with electron beam (a) ion beam (b) [70, 71].

3D reconstructions of oxide layers were performed

on two 316 L samples tested in PWR Primary

environment at 300 °C and at different strain rates

(Δεt/dt = 0.004%/s and Δεt/dt = 0.4%/s). The duplex

structure of the oxides is clearly visible in Fig. 20

(yellow part corresponds to the Cr-rich layer and red

part to the Fe-rich layer).

Fig. 20 is showing some emerging shear bands

more oxidized than the matrix. The distribution of

oxide thickness using 3D data analysis confirmed this

point (Fig. 21): oxide penetration in matrix mean

depth is 420 nm whereas it reaches up to 600 nm (1.5

times higher) along shear bands. At a lower strain rate,

the oxides have the same duplex structure but with a

more important Cr-rich oxide penetration. The mean

intragranular oxide thickness is about 770 nm and

reaches about 1,150 nm along shear bands with a

corresponding ratio remaining about 1.5. It should

also be noted that in case of monotonic loading during

about 2,000 hours [73], oxide thickness is about 10

times lower, suggesting that cyclic loadings also

highly increase the corrosion processes.

The reconstructed topographic contrast of Cr-rich

layer/Metal interface (Fig. 22) also shows preferential

oxidation along shear bands as well as areas oxidized

in “round” shape with a mean diameter of 410 nm

(Fig. 23). The measured distribution and mean size of

these features are in good agreement with the

dislocation cells that were measured by Garcia [61]. In

addition, this round shape structure was not present

at Δεt/dt = 0.4%/s, range of solicitation amplitude

where Garcia did not see any dislocations cells.

This suggests that oxidation processes are not

only accelerated along emerging shear bands but also on

Area 1

Area 2

Fe rich oxide

Cr rich oxide

Shear bands


206

Fig. 20 3D oxide mapping using dual beam, Δεt/dt = 0.004%/s, Δεt/2 = 0.5%/s, 300 °C in PWR Primary environment (16 µm × 9 µm).

0 200 400 600 800 1000 1200 1400 16000

0.005

0.01

0.015

0.02

0.025

0.03

0.035

0.04

Fre

quen

cy

Oxide pentration (nm)

Mean value = 425.0198

Median value = 422.349

SD = 122.8401

Fig. 21 Distribution of Cr-rich oxide penetration, Δεt/2 = 0.5%, 300 °C in PWR Primary environment. (a) Δεt/dt = 0.4%/s, analyzed volume: 1.5 µm × 20 µm, Nf = 3,309 cycles.

: Cr rich oxide layer

: Fe rich oxide layer


207

0 200 400 600 800 1000 1200 1400 16000

0.005

0.01

0.015

0.02

0.025

0.03

0.035

0.04

Fre

quen

cy

Oxide pentration (nm)


Median value = 763.989SD = 277.1833

Fig. 21 Distribution of Cr-rich oxide penetration, Δεt/2 = 0.5%, 300 °C in PWR Primary environment; (b) Δεt/dt = 0.004%/s, analyzed volume 16 µm × 9 µm, Nf = 1,328 cycles.

Fig. 22 Reconstructed topographic contrast based on 3D oxide reconstruction, Δεt/dt = 0.004%/s, Δεt/2 = 0.5%, 300 °C in PWR Primary environment (16 µm × 9 µm), whiter colors correspond to deeper oxide penetrations.

Area 1

Area 2


208

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.90

0.05

0.1

0.15

0.2

0.25

0.3

0.35

0.4

Fre

quen

cy

Round object (µm)


Median value = 0.38889

SD = 0.13715

Fig. 23 Oxidized dislocations cells distribution, Δεt/dt = 0.004%/s, Δεt/2 = 0.5%, 300 °C in PWR Primary environment (16 µm × 9 µm).

dislocations cells formed during cyclic loading.

The variation of distance between oxidized shear

bands was also measured using images performed

with the back scattered electron detector in Z-contrast

mode enhancing chemical contrasts (Fig. 24). Results

on 200 assumed shear bands randomly selected at the

surface of the 2 samples are presented in Figs. 25a and

25b. These results clearly demonstrate that distances

between oxidized shear bands increase when strain

rate decreases. Thus at low strain rate, a deeper oxide

penetration along shear bands is coupled with a higher

distance between slip bands. This may induce a

difference of strain localization at the metal/oxide

interface potentially leading to an earlier initiation

stage at low strain rate. The relation between oxide

morphology/structure and strain rate is still

investigated as well as the very first step of

environment interaction and crack initiation.

5. Corrosion of 316 L Austenitic Stainless Steel


As seen when describing SCC and corrosion fatigue

that are involved microstructural evolution, the

properties of the oxides films play a very important

role as well as in other coupled corrosion/mechanic

degradation mechanisms (flow assisted corrosion,

tribocorrosion...). Many factors affecting corrosion

rate and cracks growth were studied by several authors

[74-83]. Nevertheless, the contribution of the oxide

layers themselves remains crucial and is still suffering

from a lack of understanding. The microstructure and

chemistry of the oxide layers formed on mild steels in

contact with high temperature water was described in

the 60 s by Potter & Mann [84]. They reported that a

double oxide layer composed of an outer layer of large

Fe3O4 crystallites and an inner layer of small Fe3O4

crystallites is formed based on optical and mass

spectroscopy investigations. In the 1980s, working on

stainless steels, Robertson [85] mentioned that the

inner layer was Cr-enriched and consisted of a

chromite spinel rather than chromine Cr2O3. These

observations were detailed twenty year later by

Ziemniak [86] showing that the inner layer was

composed of a fcc spinel with a stochiometry close to

(Ni0.2Fe0.8)(Fe0.3Cr0.7)2O4 and that the outer layer was


209

Fig. 24 BSE image, Δεt/dt = 0.004 %/s, Δεt/2 = 0.5%, 300 °C in PWR Primary environment.

2 4 6 8 10 120

0.02

0.04

0.06

0.08

0.1

0.12

0.14

0.16

0.18

0.2

Mean value = 3.5226

Median value = 3.56

SD = 1.1383

Distance between oxidized shear bands (µm)

Fig. 25 Distance between oxidized shear bands (mm) 300 °C in PWR Primary environment, 200 shear bands. (a) Δεt/dt = 0.4 %/s, Δεt/2 = 0.5%.

Fre

que

ncy


210

2 4 6 8 10 120

0.02

0.04

0.06

0.08

0.1

0.12

0.14

0.16

0.18

0.2

Mean value = 6.5982

Median value = 6.55

SD = 1.8622

Fig. 25 Distance between oxidized shear bands (mm) 300 °C in PWR Primary environment, 200 shear bands. (b) Δεt/dt = 0.004 %/s, Δεt/2 = 0.5%.

almost pure magnetite Fe3O4 and more recently,

Terachi [87] showed that they were AB2O4 spinel (A

and B = Fe, Cr and Ni). The effects of chromium and

minor elements, cold-work, water chemistry and heat

treatments rather than the oxidation mechanism itself

were also studied [84, 85]. All the open literature

agreed that oxide layer formed in PWR primary

environment has duplex structure while its chemistry

and crystal structure remain controversial [87-101].

Even if it is well established that the outer layer is

composed of magnetite Fe3O4, the inner layer

composition is reported to be either a fcc phase

enriched in chromium [76, 77, 90-101], or an

hexagonal phase (Cr2O3) [87, 92, 93], or even a

mixture of both phases [92]. What’s more the early

stages of oxidation process (1 min to 5 hours) without

applying stress was still not investigated since the

already published results describe oxides formed after

long exposure times (up to 1,000 hours). This was

performed by Machet [99] on Inconel alloys (thought

to behave similarly than austenitic stainless steels). He

showed that oxidation mechanisms are operative at the

very beginning of exposure to the coolant. Even if

Ziemniak et al. [88] briefly suggested the existence of

an epitaxial relationship, their remark was entirely

qualitative and no detailed information on the

potential epitaxial relationship between oxide and

metal is available. These two points need to be

clarified if a better understanding of the properties

(diffusion rate of species, passive and mechanical

properties...) of these layers is to be reached. The

results presented in this part are obtained in the PhD

work of R. Soulas [102]. This PhD work is performed

under the European project PERFORM 60.


5.2.1 Crystallographic Structure of the Oxide

Nanolayers

Similarly for the Inconel case [101], the oxidation

kinetics of 316 L austenitic stainless steel is very fast

within the first ten minutes of exposure leading that

oxide layers formed in that condition are in the range

of 10 nm (Fig. 26), making it almost impossible to

study their local structures using electron diffraction

in TEM [103]. Nevertheless, the analysis with

ASTAR software of HRTEM images acquired on a Cs

Fre

que

ncy


211

Fig. 26 HRTEM image on a sample with a grain [111] oxidized during 5 h. Images made with Titan 80-300 kV probe corrected and available in CIM-PACA platform in Marseille (T. Neisus).

objective lens corrected TEM allows us to fill this gap

[63, 104]. Indeed post analysis with ASTAR software

replacing electron diffraction pattern by Fourier

transformed ones, give access to a precise evaluation

of the phases contained in nanometric layers. Owing

to this technique, R. Soulas studied the crystalline

structure of such oxide layers as a function of 316 L

grains crystallographic orientation, and their evolution

from an amorphous state to a monocrystalline layer

which has epitaxial relationship with the metallic

substrate (Figs. 26 and 27 [105-107]). Combining

these structural information with EFTEM images and

EELS chemical analysis (Fig. 28), Soulas showed

without any ambiguity that the inner oxide layer

formed in primary environment on 306 L is not a

chromite spinel (FeCr2)O4 but rather a (Fe1.5Cr1.5)O4

spinel. ASTAR analysis of HRTEM images clearly

showed that (Fe1.5Cr1.5)O4 oxide is oriented close to

the same crystallographic direction than sub laying

316 L grain (i.e. cube-cube orientation) whatever the

initial orientation of the surface, and that a 10°

misorientation was measured in the worst case. One

limit should be pointed out from ASTAR use. It was

impossible to make the difference between the outer

Outer oxide layer

Inner oxide layer


212

Fig. 27 ASTAR orientation map obtained on the orange square plotted on Fig. 26.

Fig. 28 (a) EFTEM images of O, Fe, Cr and Ni recorded from the sample oxidized 5 hours. (b) Cr/Fe and Ni/(Cr+Fe) variations in metal, nickel enriched layer, inner and outer oxide layer.

iron spinel and the inner chromium rich spinel because

of the small difference of their lattice parameter (less

than 1%).

Applying this mythology to the oxide layers formed

at different exposure times in primary water, Soulas et

al. [105] was able to describe the first step sequences

of the oxide growth on 316 L down to the atomic

scale.

5.2.2 Epitaxial Relationship between Oxide and

Metal; Misfit Dislocations Network at Metal & Oxide

Interface

By selecting reflections in the FFT (Fast Fourier

Transform) patterns of HRTEM images (Fig. 29a),

IFFT (inverse fast Fourier transform) images in

known directions can be reconstructed (Fig. 29b). The

IFFT images corresponding to (002) and (00-2)

reflections displayed in Fig. 29b, clearly show that the

interfacial misfit is accommodated by dislocations

every 1.26 nm. Knowing orientation and lattice

parameter of the inner layer and of the metal,

calculations showed that a deformation of 16% has to

be applied to accommodate if no dislocations are

created. By creating dislocations every 7 planes, the

deformation is reduced to an elastic deformation less


213

Fig. 29 HRTEM analysis of the misfit dislocations network lying at metal/oxide. (a) Magnified zone at the metal/oxide interface presented in (Fig.), insert corresponding with a [002] and [00-2] reflections circled.

Fig. 29 HRTEM analysis of the misfit dislocations network lying at metal/oxide. (b) Inverse FFT on selected reflections; (a)

Magnified zone showing perfect edge dislocations network (periodicity: 7 planes) with the Burger’s vector [202].

(a)

(b)


214

710 715 720 725 730 735 7400.0

2.0x105

4.0x105

6.0x105

8.0x105

1.0x106

1.2x106

Inte

nsit

y /u

.a.

Energy loss /eV

Metal Inner oxide layer Outer oxide layer

(a)

Fig. 30 EELS and HREELS spectrum on L3/L2 Iron edge obtained on the oxide layers (a) Measure of the energy shift based on internal standard (energy spread 0.7 eV, 316 L oxidation time oxidation time 2 mn).

710 720

0.0

5.0x103

1.0x104

1.5x104

2.0x104

2.5x104

3.0x104

Inte

nsit

y /u

.a.

Energy loss /eV

Metal Inner oxide layer Outer oxide layer

(b)

Fig. 30 EELS and HREELS spectrum on L3/L2 Iron edge obtained on the oxide layers (b) Details of the Fe L3/L2 edge using a monochromated beam (energy spread 0.2 eV, 316 L oxidation time 2 mn).

than 3%. Theoretically, a perfect edge dislocation

induces a displacement of 2

2a i.e. of 2.54 Å. With a

10° misorientation, this value decreases to 2.50 Å

which appears to be enough for accommodating the

difference between metal and inner oxide parameters.

5.3 Study of the Oxydation Degree

As previously mentioned, the inner oxide layer on

316 L is transforming from nanocrystallites spinel

mixture to a monocrystalline chromium rich spinel

layer during the first ten minutes of exposure to

Fe3+ Fe2+


215

primary environment. EELS spectrometry was used

for studying the evolution of metallic compounds.

Some of the results obtained on a Cs condenser

corrected and monochromated TITAN at EDF R&D

are presented in Fig. 30. A few eV shift of the L3/L2

edge toward the higher energy losses are observed

when getting from metal to the oxide (Fig. 30a). When

using a monochromated beam a clear splitting of the

L3/L2 bands is observed in the oxide indicating that

iron has two oxidation state in the oxide Fe2+ and Fe3+

(Fig. 30b) [107]. Quantifying this ratio and its

variation according to corrosion time and oxide

thickness, will give us the possibility to know if the

spinel oxide is normal, intermediate or inversed and

thus to get some new information on the passive

property of the oxide films.

6. General Conclusion

Development of microcopy techniques such

LAWATAP, Dual Beam (low kV ion, 3D facilities),

aberration corrected TEM and STEM with or without

a monochromated beam now gives access to valuable

quantitative information for improving the

understanding of materials degradations observed in

PWR reactors. It was seen that the observed

segregation on dislocations loops is an important key

for improving the understanding and modeling of

hardening mechanisms in irradiated stainless steel and

their contribution to IASCC susceptibility. 3D images

using dual beam and analysis of oxide topography as a

function of cyclic solicitation parameters clearly

underlined the coupling effect on primary

environment under cyclic loading in 304 L.

Deformation mapping along strain path coupled with

TEM studies on specimen with perfectly known strain

path, strain amplitude and crystallographic orientation

with dual beam microscope gave access to a precise

evaluation of the interaction of mechanical solicitation

on oxide and cracks growth. A clear effect of strain

and its speed, and also of the nature of mechanical

loading (tensile or cyclic loading) on corrosion

behavior was shown. The oxidation mechanism can

now be better understood thank to the combination of

all these advanced techniques.

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Presented at the Pasturel., 25th conference of SFME, Strasbourg.

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