Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels

Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels Comprehensive nuclear materials 4 03 ferritic steels and advanced ferritic–martensitic steels

B Raj and M Vijayalakshmi

Indira Gandhi Centre for Atomic Research, Kalpakkam, India

ß 2012 Elsevier Ltd All rights reserved.

4.03.4.1 Influence of Composition and Microstructure on Properties of Ferritic Steels 103

4.03.6 Ferritic Steels for Out-of-Core Applications: Improvements in Joining 116

Abbreviations

bcc Body-centered cubic

CSL Coincident site lattice

DBTT Ductile to brittle transition temperature

DICTRA Diffusion-controlled transformations

dpa Displacements per atom

EBR Experimental breeder reactor

EBSD Electron back scattered diffraction

fcc Face-centered cubic

FFTF Fast flux test facility

GBCD Grain boundary character distribution

GBE Grain boundary engineering

HAADF High angle annular dark field

HAZ Heat-affected zone

HFIR High flux isotope reactor

ITER International Thermonuclear

Experimental Reactor

ODS steel Oxide dispersion strengthened steel

PAGS Prior austenite grain size

PFR Power fast reactor

PWHT Postweld heat treatment

RIS Radiation-induced segregation

SIPA Stress-induced preferential

absorption

SIPN Stress-induced preferential nucleation

TEM Transmission electron microscopy

▽DBTT Change in DBTT

4.03.1 Introduction

The widespread acceptance of nuclear energydepends1 on the improved economics, better safety,sustainability, proliferation resistance, and waste man-agement Innovative technological solutions are beingarrived at, in order to achieve the above goals Theanticipated sustainability, rapid growth rate, and eco-nomic viability can be ensured by the judicious choice

of fast reactor technology with a closed fuel cycleoption The fast reactor technology has attained(http://www.world-nuclear.org/info/inf98.html) a highlevel of maturity in the last three decades, with

390 years of successful operation The emerging national collaborative projects (http://www.iaea.org/INPRO/;http://www.gen4.org/) have, therefore, cho-sen fast reactors as one of the important constituents

inter-of the nuclear energy in the twenty-first century.The nuclear community has been constantly striv-ing for improving the economic prospects of thetechnology The short-term strategies include thedevelopment of radiation-resistant materials andextension of the lifetime of the components Theachievement of materials scientists in this field isremarkable Three generations of materials havebeen developed,2increasing the burn-up of the fuelfrom 45 dpa for 316 austenitic stainless steel to above

180 dpa for ferritic steels Presently, efforts are in

97

progress to achieve a target burn-up of 250 dpa,

using advanced ferritic steels The attempts by

nuclear technologists to enhance the thermal

effi-ciency have posed the challenge of improving the

high temperature capability of ferritic steels

Addi-tionally, there is an inherent disadvantage in ferritic

steels, that is, their susceptibility to undergo

embrit-tlement, which is more severe under irradiation

It is necessary to arrive at innovative solutions to

overcome these problems in ferritic steels In the

long time horizon, advanced metallic fuels and

cool-ants for fast reactors are being considered for

increasing the sustainability and thermal efficiency

respectively Fusion technology, which is ushering

(http://www.iter.org/proj) in a new era of

opti-mism with construction of the International

Ther-monuclear Experimental Reactor (ITER) in France,

envisages the use of radiation-resistant advanced

ferritic steels Thus, the newly emerging scenario

in nuclear energy imposes the necessity to

reevalu-ate the mreevalu-aterials technology of today for future

applications

The genesis of the development of ferritic steels is,

indeed, in the thermal power industry The

develop-ment of creep-resistant, low alloy steels for boilers

and steam generators has been one of the major

activities in the last century Today, the attempt to

develop ultra super critical steels is at an advanced

stage Extensive research of the last century is

responsible for identifying certain guidelines to

address the concerns in the ferritic steels The merit

of ferritic steels for the fast reactor industry was

established3 in the 1970s and since then, extensive

R&D has been carried out4 on the application of

ferritic steels for nuclear core component

A series of commercial ferritic alloys have been

developed, which show excellent void swelling

resis-tance The basic understanding of the superior

resistance of the ferrite lattice to void swelling, the

nature of dislocations and their interaction with

point defects generated during irradiation have been

well understood The strengthening and deformation

mechanisms of ferrite, influence of various alloying

elements, microstructural stability, and response of

the ferrite lattice to irradiation temperature and stress

have been extensively investigated The mechanism of

irradiation hardening, embrittlement and methods to

overcome the same are studied in detail Of the

dif-ferent steels evaluated, 9–12% Cr ferritic–martensitic

steels are the immediate future solution for fast

reac-tor core material, with best void swelling resistance

and minimum propensity for embrittlement

The high temperature capability of the ferriticsteels has been improved from 773 to 973 K, bylaunching the next generation ferritic steels, whichare currently under evaluation for nuclear applica-tions, namely the oxide dispersion strengthened(ODS) ferritic steels (seeChapter4.08, Oxide Dis-persion Strengthened Steels) Conceptually, thisseries of steels combines the merits of swelling resis-tance of the ferrite matrix and the creep resistanceoffered by inert, nanometer sized, yttria dispersions

to enhance the high temperature limit of the ODSsteels to temperature beyond 823 K The concerns ofthis family of materials include optimization of thechemistry of the host lattice, cost effective fabricationprocedure, and stability of the dispersions under irra-diation, which will be discussed in this article.The present review begins with a brief introduc-tion to the basic metallurgy of ferritic steels, summar-izing the influence of chemistry on stability of phases,decomposition modes of austenite, different types ofsteels and structure–property correlations The mainthrust is on the development of commercial ferriticsteels for core components of fast reactors, based ontheir chemistry and microstructure Hence, the nextpart of the review introduces the operating conditionsand radiation damage mechanisms of core compo-nents in fast reactors The irradiation response offerritic steels with respect to swelling resistance, irra-diation hardening, and irradiation creep are high-lighted The in-depth understanding of the damagemechanisms is explained The main concerns of fer-ritic steels such as the inferior high temperature irra-diation creep and severe embrittlement are addressed.The current attempts to overcome the problems arediscussed Finally, the development of advancedcreep-resistant ferritic steels like the ODS steels, forfission and fusion applications are presented Theapplication of ferritic steels for steam generator cir-cuits and the main concerns in the weldments offerritic steels are discussed briefly The future trends

in the application of ferritic steels in fast reactortechnology are finally summarized

4.03.2 Basic Metallurgy of Ferritic–Martensitic Steels

The advanced ferritic and ferritic–martensitic steels

of current interest have evolved5 from their cessors, the creep-resistant ferritic steels, over nearly

prede-a century The first of the series wprede-as the cprede-arbon prede-andC–Mn steels with a limited application to about

523 K Subsequent developments through different

levels of chromium, molybdenum have increased the

high temperature limit to 873, leading to the current

ferritic and ferritic–martensitic steels, that is, the

9–12% Cr–Mo steels In addition to being

economi-cally attractive, easy control of microstructure using

simple heat treatments is possible in this family of

steels, resulting in desired mechanical properties

The propensity to retain different forms of bcc

ferrite, that is, ferrite or martensite or a mixture at

room temperature in Cr–Mo steels, depends crucially

on the alloying elements Extent of the phase field

traversed by an alloy on heating also depends on the

amount of chromium, silicon, molybdenum,

vana-dium, and carbon in the steel The combined effect

of all the elements can be represented by the net

chromium equivalent, based on the effect of the

aus-tenite and ferrite stabilizing elements A typical

pseu-dobinary phase diagram6 is shown in Figure 1(a)

Increase in chromium equivalent by addition of ferrite

stabilizers or V or Nb would shift the Fe–9Cr alloy

into the duplex phase field at the normalizing

temperature The phase field at the normalizing

tem-perature and the decomposition mode7–9 of high

temperature austenite (Figure 1(b)) dictate the

result-ing microstructure at room temperature and hence, the

type of steel Accordingly, the 9CrMo family of steels

can either be martensitic (9Cr–1Mo (EM10) or

stabi-lized 9Cr–1MoVNb (T91)), ferritic (12Cr–1MoVW

(HT9)) or ferritic–martensitic (9Cr–2Mo–V–Nb

(EM12)) steel The stabilized variety of 9–12 CrMo

steels could result10in improved strength and delayed

grain coarsening due to the uniform distribution of fine

niobium or vanadium carbides or carbonitrides

The transformation temperatures and the kinetics

of phase transformations depend strongly on the

composition of the steels Sixteen different 9Cr steels

have been studied11,12and the results, which provide

the required thermodynamic database are shown in

Figure 2, with respect to the dependence of melting

point, Ms temperature and the continuous heating

transformation diagrams The constitution and the

kinetics of transformations dictate microstructure

and the properties

In the early stages, the oxidation resistance and

creep strength were of prime importance, since the

Cr–Mo steels were developed4 for thermal power

stations In addition to the major constituent phases

discussed above, the minor carbides which form at

temperatures less than 1100 K, dictate the long term

industrial performance of the steels Evaluation of

tensile and creep properties of Cr–Mo steels exposed

to elevated temperature for prolonged durations havebeen extensively studied.5,13,14The following trendswere established: The optimized initial alloy compo-sition considered was 9Cr, W–2Mo¼ 3, Si ¼ 0.5, with

C, B, V, Nb, and Ta in small amounts Higher mium content has two effects: it increases the hard-enability leading to the formation of martensite andalso promotes the formation of d-ferrite therebyreducing the toughness A reduction in the chromium

Bainite Widmanstatten ferrite

Ferrite Pearlite

Upper bainite Lower bainite

content lowers the oxidation resistance If Wþ Mo

concentration is kept <3%, creep strength will

reduce, while higher amount promotes the formation

of d-ferrite and brittle Fe2Mo Laves phase The

addi-tion or partial replacement of molybdenum with

tungsten and boron increased the stability of M23C6,

and slowed down the kinetics of recovery Lower

nickel introduced d-ferrite, while its increase reduces

creep strength When Si is less than 0.3%, oxidation

resistance gets lowered, while higher silicon content

led to agglomeration of carbides with an increased

amount of d-ferrite On similar lines, the composition

of all other elements could also be optimized, based

on structure–property correlation studies

The components of the steam generators are often

subjected to repeated thermal stresses as a result of

temperature gradients that occur on heating and

cool-ing durcool-ing start-ups and shutdowns or durcool-ing

variations in operating conditions Steady stateoperation in between start-up and shutdown ortransients would produce creep effects Thereforethe low cycle fatigue (LCF) and creep–fatigue inter-action assume15 importance in the safe life designapproach of steam generator components The alloyexhibited a decrease in fatigue life with increasingtemperature, thus limiting its upper limit of tempera-ture up to about 773 K

The joining technologies of Cr–Mo steels havebeen well investigated.16,17 One of the major pro-blems during welding of ferritic steels has been theformation of d-ferrite, if the amount of ferrite stabi-lizers is high The partial substitution of Mo with

W enables austenite stabilization and hence reducesthe tendency to form d-ferrite In fact, there needs

to be an intricate balance between austenite andferrite stabilizing elements in 9–12Cr–Mo steels

750

1820 1815 1810 1805 1800 1795 1790 1785

M s /K = 904- 474 (C + 0.46(N - 0.15Nb ) - 0.046Ta)

-{17Cr + 33Mn + 21Mo + 20Ni + 39V + 5W)

-45Mn 2 - 25Ni 2 - 100V 2 + 10Co }- 44.5Ta

9Cr–ferritic martensitic steels

Figure 2 Influence of chemistry on transformation temperatures (Ms and melting point) and kinetics of transformation

of g ! a þ carbide, in various ferritic steels.

This would ensure a satisfactory solidification process

with a fully austenitic structure Additionally, this

enables easier hot workability during primary

proces-sing and tubemaking, without loproces-sing high

tempera-ture creep resistance The formation of d-ferrite

reduces toughness due to the notch sensitivity,

pro-motes solidification cracking and embrittlement due

to sigma-phase precipitation and reduces the creep

ductility at elevated temperatures of operation Other

problems relate to solidification cracking, hydrogen

cracking, and reheat cracking, which have been

exten-sively studied.18The Type IV cracking in ferritic steel

weldments and the brittle layer formation in the

dis-similar welds are discussed in detail later

4.03.3 Radiation Damage of Core

Components in Fast Reactors

The core components in fast reactors include the

following: clad (cylindrical tubes which house

the fuel pellets) for the fuel and wrapper (a container

which houses fuel elements, in between which the

coolant flows) for fuel subassemblies.Figure 3shows

a schematic of clad and wrapper in a typical fuel

subassembly The necessity to develop robust

tech-nology for core component materials arises from the

fact that the ‘burn-up’ (energy production from unit

quantity of the fuel) of the fuel depends on theperformance of the clad materials The higher burn-

up of the fuel increases the ‘residence time’ of thesubassembly in the core, eventually lowering the cost.The core component materials in fast breederreactors are exposed to severe environmental serviceconditions The differences in the exposure condi-tions of the clad and wrapper in a fast reactor coreare listed in Table 1 Under such exposure condi-tions, materials in the fast reactor fuel assembliesexhibit many phenomena (Figure 4), specific to fastreactor core: Void swelling, irradiation growth, irra-diation hardening, irradiation creep, irradiation, andhelium embrittlement

Another selection criterion, namely the bility of the core component materials with the cool-ant, the liquid sodium, has already been established.Presently, methods are known to avoid interaction ofthe clad material with the coolant

compati-Detailed books and reviews19,20,21,22,23 are able on all the degradation mechanisms mentionedabove, which are related to the production, diffu-sion, and interaction of point defects in the specificlattice of the material Hence, a brief introduction ispresented below (see alsoChapter1.03, Radiation-Induced Effects on Microstructure;Chapter1.11,Primary Radiation Damage Formation; andChapter

avail-1.04, Effect of Radiation on Strength and Ductility

of Metals and Alloys)

Void swelling in a fast reactor core can change

a cube of nickel to increase (20%) its side from

1 cm to 1.06 cm, after an exposure to irradiation of

1022n cm2 Void swelling is caused by the sation of ‘excess vacancies’ left behind in the latticeafter ‘recombination’ of point defects produced dur-ing irradiation Void swelling is measured using thechange in volume (▽V/V) of bulk components of thereactor or image analysis of voids observed usingtransmission electron microscope (TEM)

conden-The ‘irradiation growth’ (fluence 1020

n cm2)can increase the length of a cylindrical rod of uraniumthree times and reduce its diameter by 50%, retainingthe same volume This occurs mainly in anisotropiccrystals, introducing severe distortion in core compo-nents It is caused by the preferential condensation ofinterstitials as dislocation loops on prism planes oftype (110) of hcp structures and vacancies as loops

on the basal planes (0001), which is equivalent totransfer of atoms from the basal planes to prismplanes,via irradiation-induced point defects.Irradiation hardening refers to the increase in

concomitant reduction in ductility, under irradiation

at temperatures<0.3Tm The large number density of

defect loops, voids, and precipitates generated during

irradiation pins the mobile dislocations and acts as an

obstacle to their further movement, requiring

addi-tional stress to unpin the immobile dislocations

The irradiation creep, the most important

param-eter for design consideration, is the augmentation of

thermal creep of the material, under irradiation This

leads to premature failure of the material and

restricts the service life The mechanisms responsible

for irradiation creep are identified as the

‘stress-induced preferential absorption’ (SIPA) and the

‘stress-induced preferential nucleation (SIPN)’ of

point defects by dislocations, which revolve around

the interaction of excess point defects generated

dur-ing irradiation with dislocations

Irradiation embrittlement, another frequent

observation in ferritic steels exposed to irradiation,

refers to the increase in the ductile to brittle

transi-tion temperature (DBTT) during irradiatransi-tion Drastic

loss in ductility at low temperatures results from a

lower sensitivity of the fracture stress, sf, due to

irradiation and less dependence on temperature

than the yield strength sy Materials with a high

value of the Hall-Petch constant are more prone tobrittle failure Such materials like ferritics releasemore dislocations into the system when a source isunlocked, causing hardening and loss of ductility.Some of the engineering materials contain nickel,

an element which undergoes an (n,a) reaction, ing high concentration of helium The binding energy

produc-of helium with a vacancy being very high 2 eV,the helium atoms stabilize the voids, enhancing theirgrowth rate Incorporation of helium during irradia-tion into voids along the grain boundaries assistsgrain boundary crack growth by linking voids causing

‘helium embrittlement.’

Of these many degradation mechanisms, the alloydevelopment programmes have focused mainly onthe void swelling, irradiation hardening, embrittle-ment, and the irradiation creep, since these are themajor life limiting factors

4.03.4 Development of Ferritic Steels for Fast Reactor Core

This section begins with the optimization of try and initial microstructure to develop swelling and

chemis-Table 1 Comparison of exposure conditions of clad and wrapper of fast reactor core

Exposure conditions (only trends; exact

values depend on core design)

Maximum temperature: 923–973 K Lower temperature range than clad:

823 K Steeper temperature gradient Lower temperature gradient Higher stresses from fission gas

Neutron fluence: 2–4  10 19 n m2

Irradiation creep at higher temperatures

Irradiation creep

Irradiation embrittlement Irradiation embrittlement Interactions with fuel and fission

products

Interaction with sodium

Selection criteria: mechanical properties Tensile strength Tensile strength

Creep strength Creep ductility

Compatibility with fuel Compatibility with fission products General common selection criteria

Good workability

International neutron irradiation experience as driver or experimental fuel subassembly

Availability

creep-resistant ferritic steels The microstructural

instability during service exposure is briefly

pre-sented The superior swelling performance of ferritic

steels is understood based on mechanisms of void

swelling suppression Following this, the

irradiation-induced/-enhanced segregation/precipitation causing

irradiation hardening is discussed The irradiation

creep and embrittlement, their mechanisms and

meth-ods to combat the problems are highlighted The R&D

efforts of today to reduce the severity of

embrittle-ment in ferritic steels, using modeling methods, are

outlined Finally, typical problems in the weldments

of ferritic steels, when used for out of core tions, are presented, emphasizing the advantage ofmodeling in predicting the materials’ behavior

Microstructure on Properties of FerriticSteels

Rapid strides have been made the world over, in thedesign and development of advanced creep-resistantferritic or ferritic–martensitic steels The low alloysteels can be used as either 100% ferrite–martensite

Linear swelling regime

-4)

-1 -2 -3

Irradiated

Unirradiated

Strain (c)

Unirradiated

Time (h) (d)

Irradiated Unirradiated

Test temperature (e)

Figure 4 Schematic representation of major damage mechanisms in the core component materials of fast reactors: (a) The different stages of void swelling, (b) irradiation growth, (c) increase in strength with a concomitant reduction in ductility during irradiation hardening, (d) increase in creep strain and reduction in creep life after irradiation caused by irradiation creep, and (e) increase in ductile to brittle transition temperature and reduction in upper shelf energy after irradiation caused

by irradiation embrittlement.

or a mixture of both It is possible to choose the

required structure by the appropriate choice of either

the chemistry or the heat treatment For example, a

completely ferrite matrix, yielding high toughness,

can be obtained in steels with chromium content

higher than 12%, with carbon reduced to less than

0.03% The same steel can be used to provide higher

strength by choosing the 100% martensite structure,

if carbon content is increased to about0.1% The

9Cr steels have always been used in the 100%

mar-tensite state Extensive studies have been carried out

on phase stabilities of these steels, with changes in

chemistry and heat treatment

The creep resistance of the plain Cr–Mo steels

has, further, been increased by the addition of

carbide stabilizers like Ti or V or Nb, leading to

the modified variety of 9–12Cr–Mo steels These

elements led24 to copious, uniform precipitation ofMonte Carlo (MC) type of monocarbides, whichare very fine and semicoherent Such precipitates arevery efficient in pinning the mobile dislocations, lead-ing to improved creep behavior at higher tem-peratures These carbides are stable at temperatureshigher than even 1273 K and hence, do not causedeterioration of long-term mechanical propertiesduring service exposure

The development of high creep-rupture strength9–12% steels with various combinations of N, Mo,

W, V, Nb, Co, Cu, and Ta is based on optimizingthe constitution (Table 2.) and d-ferrite content,increasing the stability of the martensite, dislocationstructure and maximizing the solid solution and pre-cipitation hardening The concentration of each ele-ment in ferritic steels has been optimized based on

an in-depth understanding of the influence of thespecific element on the behavior of the steel Theextensive studies related to optimization of chemistryare summarized in Table 3 Based on the strongscientific insights, large number of commercial steelshave been developed (Table 4) in the later half ofthe last century

Most of this family of ferritic–martensitic steels

is used in the normalized and tempered condition orfully annealed condition to achieve the desirable phase.The type of structure that is deliberately favored in agiven steel depends on the end application

The microstructure of the steels in normalizedand tempered conditions consists24(Figure 5) of (a)martensite laths containing dislocations with a Burgersvector 1/2a0<111> with a density of approximately

1 1014

m2 (b) coarse M23C6 particles located at

Table 2 Optimizing the constitution in the development

of ferritic steels

resistance, hardenability

Mo, W, Re, Co Solid solution strengthening

V, Nb, Ti, Ta Strengthening by formation of

MX-carbonitride

C, N Austenite stabilizer, solid solution

strengthening, carbonitride formers

stabilization of carbide

Ni, Cu, Co Austenite former, inhibits d-ferrite

formation

Table 3 Beneficial and harmful effects of different elements during design of creep-resistant ferritic steels

Hardenability

Laves phase

prior austenite and ferrite grain boundaries with finer

precipitates within the laths and at martensite lath and

subgrain boundaries M2X precipitates rich in Cr are

isomorphous with (CrMoWV)2CN

The initial microstructure of the normalized and

tempered steels described above does not remain

stable during service in a nuclear reactor

Pro-longed exposure at high temperature causes changes

in the initial microstructure, which has been studied

extensively The M2X precipitates in the

normal-ized and tempered stabilnormal-ized 9Cr–1Mo steels are

gradually replaced (Figure 6) by MX, intermetallic,

and Laves phases during prolonged aging at high

temperature

The high temperature and the irradiation over

prolonged time of exposure introduce microstructural

instabilities These instabilities are caused mainly by

the point defects caused by irradiation and complex

coupling of these defects with atoms in the host

lattice, their diffusion or segregation and finally the

precipitation There is a recovery of the defect

structure since the irradiation-induced vacancies

alter the dislocation dynamics There are three types

of processes with respect to evolution of secondary

phases: induced precipitation,

irradiation-enhanced transformations, and the irradiation

modi-fied phases It is seen that the evolution of these phases

depends on the composition and structure of the steel

and the irradiation parameters like the temperature,

dose rate, and the dose Evolution of

irradiation-induced phases and their influence on hardening

and embrittlement is discussed later

Extensive experimental investigations found3that theferritic steels, whose high temperature mechanicalproperties are far inferior to austenitic stainless steels,displayed excellent radiation resistance The ferritic–martensitic steels (9–12% Cr) have, therefore, beenchosen for clad and wrapper applications, in order

to achieve the high burn-up of the fuel This isbased26–29(Table 5) on their inherent low swellingbehavior The 9Cr–1Mo steel, modified 9Cr–1Mo(Grade 91), 9Cr–2Mo, and 12Cr–1MoVW (HT9)have low swelling rates at doses as high as 200 dpa.For example, HT9 shows 1% swelling at 693 K for

200 dpa The threshold dose for swelling in ferriticsteels is as high as nearly 200 dpa in contrast to 80 dpafor the present generation D9 austenitic stainlesssteel It is established that the void swelling dependscrucially on the structure of the matrix lattice, inwhich irradiation produces the excess defects.Extensive basic studies have identified19,30–33 thefollowing reasons as the origin of superior swellingresistance in ferritic steels:

1 The relaxation volume for interstitials, that is, thevolume of the matrix inwhich distortion is introducedbycreating an interstitial, in bcc ferrite is larger19thanfcc austenite For every interstitial introduced, thelattice distortion is high and hence the strain energy

of the lattice Hence, the bias toward attracting oraccommodating interstitials in the bcc lattice is less.This leaves higher density of ‘free’ interstitials in thebcc lattice than fcc lattice As a result, recombina-tion probability with vacancies increases significantlyand supersaturation of vacancy reduces Conse-quently, the void nucleation and swelling is less

2 The migration energy of vacancies in bcc iron isonly 0.55 eV, against a high value in fcc austenite,1.4 eV Vacancies are more mobile in bcc thanfcc, increasing the recombination probabilities inbcc ferrite Another factor is the high bindingenergy between carbon and vacancy in bcc iron(0.85 eV), while it is only 0.36–0.41 eV in austenite.This leads19to enhanced point defect recombina-tion in bcc than fcc, due to more trapping ofvacancies by carbon or nitrogen

3 In bcc iron, it is known30that there is a strong action between dislocations and interstitials solutes,forming atmospheres of solutes around dislocations.The formation of ‘atmospheres’ around dislocationsmakes them more effective sinks for vacancies thaninterstitials, resulting in suppression of void growth,

inter-Table 4 List25 of commercial ferritic steels, their

chemistry, and properties

provided the following conditions are satisfied:

‘atmo-spheres’ comprise of either oversized substitutional

atoms or interstitials, dislocations have high binding

energy with solutes, and concentration of solute

atoms at the core of the dislocation exceeds a critical

value On the other hand, if ‘atmosphere’ is made up of

undersized atoms like Si or P, the voids can grow The

‘atmosphere’ of interstitials reduces the dislocation

bias for additional capture and inhibits dislocation

climb, thus converting them to saturable sinks Such

a scenario would increase the recombination

prob-abilities, suppressing the void growth

These fundamental differences in the behavior of

solutes and point defects in bcc lattice make ferritic

steel far superior to austenitic steels, with respect to

radiation damage

The challenging task for materials scientists to useferritic steels directly in fast reactor fuel assemblywas with respect to enhancing the high temperaturemechanical properties of the ferritic steels, especiallyhigh temperature creep life and irradiation creepresistance

Ferritic SteelsThe initial microstructure of the steels evolves dur-ing service, due to high temperature and irradiationfor prolonged times, leading to modification of defectstructure and secondary phases These changesharden the steel, leading to concomitant embrittle-ment, which is discussed below

A

B

(001) (b)

2.00

(c)

NbKa FeKa

V Kb

V Ka

NbLa

2.00 (d)

It is reported that carbon content in 12%

chro-mium steel is maintained high in order to use the

steel as martensitic steels The high amount of

car-bon in 12% chromium steel leads to copious

precip-itation of carbides, that is, twice as much in 9Cr

steels Both the steels have predominantly M23C6

carbides with a small fraction of monocarbides,

eventually leading34to deterioration of their tance to brittle failure The critical stress to propa-gate a crack is inversely proportional to the cracklength If it is assumed that fracture initiates at anM23C6precipitate and the crack length at initiationequals the diameter of a carbide particle then thefracture stress will decrease with increasing precipi-tate size The precipitates coarsen during irradiation

resis-in the range of 673–773 K, thus causresis-ing a decrease

in fracture stress and an increase in DBTT even inthe absence of further hardening

Additionally, Cr rich, bcc a0precipitates formed35

in the higher chromium steel during thermal sure and irradiation lead to hardening and embrittle-ment of the steel The d-ferrite, into which there is arepartitioning of chromium, is harmful, since it pro-motes formation of a0 The presence of very finecoherent particles of the w (Fe2Mo) phase has alsobeen reported in the T91 and HT9 steels The w phasewas observed to form more rapidly in the 9Cr–2Motype of steels, both in the d-ferrite and martensitephases This is possibly due to the higher amount of

expo-Mo in the EM12 type of steels The w phase isenriched in Fe, Si, and Ni and contains significantamount of Mo and P The G phase (Mn7Ni16Si17)has been found to form very occasionally in the mod-ified 9Cr–1Mo and HT9 (12% Cr) variety of steels.The s phase (Fe–Cr phase, enriched in Si, Ni,and P) has been observed to form around the M23C6particles in 9–13% Cr martensitic steels after irradia-tion at 420–460C in Dounray Fast Reactor Inaddition Cr3P needles and MP (M¼ Fe, Cr, and Mo)particles have also been detected in the 12 and 13Crsteels in the range of 420–615C The formation ofthese phases during irradiation may be understood interms of the strong radiation-induced segregation(RIS) of alloying and impurity elements to point defectsinks in the steels (see Chapter 1.18, Radiation-Induced Segregation) The RIS of alloying/impurityelements could lead36to either enrichment or deple-tion near the sinks, depending on the size of the atomand its binding energy with iron self-interstitials.Table 5 Void swelling resistance26–29of some commercial ferritic steels

Commercial

name

Chemistry and country of origin Reactor in which irradiation

was carried out

Burn-up achieved (dpa)

(b)

SiKa

Figure 6 Effect24of prolonged exposure (823 K per

10 000 h) of modified 9Cr–1Mo steel Transmission electron

micrograph showing (a) formation of detrimental Fe2Mo

Laves intermetallic phase around the M23C6 The insets

show the microdiffraction pattern and magnified view of the

nucleation of Laves phase (b) EDAX spectrum confirming

the enrichment of iron and molybdenum.

Generally, a large number of alloying elements, W, Nb,

Mo, Ta, V, or Ti are dissolved into the matrix of ferritic

steels, some of them being larger than the iron atom

This could lead to the expansion of the unit cell of

ferrite, making an element say, chromium undersized,

with a positive binding energy with iron

self-intersti-tial Such a situation would lead to enrichment of

chromium near the sink-like grain boundary The

reverse could happen if the size of the alloying

ele-ments happen to be smaller than iron

The w, G ands phases are all enriched in Si and

Ni – elements which are known to segregate to

in-terfaces during irradiation With the exception of

G phase, all the other phases and the a0 phase are

rich in Cr In those ferritic steels, where Cr is

depleted near voids and at other interfaces which

act as point defect sinks, it follows that in steels

containing higher than 11 or 12% Cr, the chromium

enrichment within the matrix may lead to local

con-centrations exceeding those (14%) at which a0

forms thermally Further, enrichment of Cr may also

result from the partial dissolution of chromium rich

precipitates such as M23C6 during irradiation In

addition, RIS of phosphorus can also lead to the

formation of phosphides in some of the steels The

irradiation-induced point defect clusters and loops

may also facilitate and enhance nucleation of these

phases Although the relatively soft d-ferrite

im-proves the ductility and toughness of the 12Cr steel,

the fracture could be initiated at the M23C6

pre-cipitates on the d-ferrite–martensite interface The

presence of d-ferrite, extensive precipitation and

radiation-induced growth of M23C6precipitates and

formation of the embrittling intermetallic phases in

the 12Cr–1MoVW steel in the temperature range

573–773 K are together responsible37for the relative

change in impact behavior of 9Cr–1MoVNb and

12Cr–1MoVW between 323 and 673 K

Irradiation-induced microstructural changes are

the factors that govern the creep and embrittlement

behavior, which therefore, has to be minimized using

appropriate chemistry and structure

Ferritic Steels

An essential prerequisite for maximizing the

‘irra-diation creep resistance’ is to ensure38 the best

combination of thermal creep behavior and

long-term microstructural stability at high temperature

Hence, the present section would discuss irradiation

creep in the same sequence as mentioned above

The design principles of development of resistant steels are as follows:

creep- Introduce high dislocation density by either formations or cold work to increase the strength ofthe basic lattice;

trans- Strengthen the host lattice by either solid solutionstrengtheners or defects;

 Stabilize the boundaries created by phase mations by precipitating carbides along theboundaries;

transfor- Arrest dislocation glide and climb by appropriateselection of crystal structure, solid solution, inter-faces, dislocation interactions, and crystal with lowdiffusivity;

 Resist sliding of grain boundaries by introducingspecial type of boundaries and anchoring theboundaries with precipitates;

 Ensure long-term stability of the microstructure,especially against recovery and coarsening of thefine second phase particles;

In the case of 9–12 Cr steels, the martensitic lathstructure (Figure 7) decorated with only MX whichshould39 be stable over long-term service life isthe most desired structure Thermo-Calc evalua-tions show39that MX can be stabilized at the expense

of M23C6only by reducing carbon to as low a value

as 0.02% in 9 Cr–1Mo steel This value is toolow to ensure acceptable high temperature mechani-cal behavior of the steels In the context of fastreactor core components, the high chromium 9–12%ferritic–martensitic steels assume relevance Hence,

an extensive database25 for a large number ofcommercial ferritic steels has been generated and

Lath boundary

Lath boundary