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.

4.03.1
4.03.2
4.03.3
4.03.4
4.03.4.1
4.03.4.2
4.03.4.3
4.03.4.4
4.03.4.5
4.03.4.5.1
4.03.5
4.03.6
4.03.7
References

Introduction
Basic Metallurgy of Ferritic–Martensitic Steels
Radiation Damage of Core Components in Fast Reactors
Development of Ferritic Steels for Fast Reactor Core
Influence of Composition and Microstructure on Properties of Ferritic Steels
Void Swelling Resistance
Irradiation Hardening in Ferritic Steels
Irradiation Creep Resistance of Ferritic Steels
Irradiation Embrittlement in Ferritic Steels

GBE to reduce embrittlement in ferritic steels
Development of Advanced ODS Ferritic Steels
Ferritic Steels for Out-of-Core Applications: Improvements in Joining
Summary

Abbreviations
bcc
CSL
DBTT
DICTRA
dpa
EBR
EBSD
fcc
FFTF
GBCD
GBE
HAADF
HAZ
HFIR
ITER
ODS steel
PAGS
PFR
PWHT
RIS
SIPA
SIPN
TEM
▽DBTT


Body-centered cubic
Coincident site lattice
Ductile to brittle transition temperature
Diffusion-controlled transformations
Displacements per atom
Experimental breeder reactor
Electron back scattered diffraction
Face-centered cubic
Fast flux test facility
Grain boundary character distribution
Grain boundary engineering
High angle annular dark field
Heat-affected zone
High flux isotope reactor
International Thermonuclear
Experimental Reactor
Oxide dispersion strengthened steel
Prior austenite grain size
Power fast reactor
Postweld heat treatment
Radiation-induced segregation
Stress-induced preferential
absorption
Stress-induced preferential nucleation
Transmission electron microscopy
Change in DBTT

97
98

101
102
103
105
106
108
110
112
114
116
119
119

4.03.1 Introduction
The widespread acceptance of nuclear energy
depends1 on the improved economics, better safety,
sustainability, proliferation resistance, and waste management. Innovative technological solutions are being
arrived at, in order to achieve the above goals. The
anticipated sustainability, rapid growth rate, and economic viability can be ensured by the judicious choice
of fast reactor technology with a closed fuel cycle
option. The fast reactor technology has attained
( a high
level of maturity in the last three decades, with
390 years of successful operation. The emerging international collaborative projects ( />INPRO/; have, therefore, chosen fast reactors as one of the important constituents
of the nuclear energy in the twenty-first century.
The nuclear community has been constantly striving for improving the economic prospects of the
technology. The short-term strategies include the
development of radiation-resistant materials and
extension of the lifetime of the components. The
achievement of materials scientists in this field is

remarkable. Three generations of materials have
been developed,2 increasing the burn-up of the fuel
from 45 dpa for 316 austenitic stainless steel to above
180 dpa for ferritic steels. Presently, efforts are in
97


98

Ferritic Steels and Advanced Ferritic–Martensitic Steels

progress to achieve a target burn-up of 250 dpa,
using advanced ferritic steels. The attempts by
nuclear technologists to enhance the thermal efficiency have posed the challenge of improving the
high temperature capability of ferritic steels. Additionally, there is an inherent disadvantage in ferritic
steels, that is, their susceptibility to undergo embrittlement, 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 coolants for fast reactors are being considered for
increasing the sustainability and thermal efficiency
respectively. Fusion technology, which is ushering
( in a new era of optimism with construction of the International Thermonuclear 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 reevaluate the materials technology of today for future
applications.
The genesis of the development of ferritic steels is,
indeed, in the thermal power industry. The development 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 resistance. 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 different steels evaluated, 9–12% Cr ferritic–martensitic
steels are the immediate future solution for fast reactor core material, with best void swelling resistance
and minimum propensity for embrittlement.

The high temperature capability of the ferritic
steels has been improved from 773 to 973 K, by
launching the next generation ferritic steels, which
are currently under evaluation for nuclear applications, namely the oxide dispersion strengthened
(ODS) ferritic steels (see Chapter 4.08, Oxide Dispersion Strengthened Steels). Conceptually, this
series of steels combines the merits of swelling resistance of the ferrite matrix and the creep resistance
offered by inert, nanometer sized, yttria dispersions

to enhance the high temperature limit of the ODS
steels to temperature beyond 823 K. The concerns of
this family of materials include optimization of the
chemistry of the host lattice, cost effective fabrication
procedure, and stability of the dispersions under irradiation, which will be discussed in this article.
The present review begins with a brief introduction to the basic metallurgy of ferritic steels, summarizing the influence of chemistry on stability of phases,
decomposition modes of austenite, different types of
steels and structure–property correlations. The main
thrust is on the development of commercial ferritic
steels for core components of fast reactors, based on
their chemistry and microstructure. Hence, the next
part of the review introduces the operating conditions
and radiation damage mechanisms of core components in fast reactors. The irradiation response of
ferritic steels with respect to swelling resistance, irradiation hardening, and irradiation creep are highlighted. The in-depth understanding of the damage
mechanisms is explained. The main concerns of ferritic steels such as the inferior high temperature irradiation creep and severe embrittlement are addressed.
The current attempts to overcome the problems are
discussed. Finally, the development of advanced
creep-resistant ferritic steels like the ODS steels, for
fission and fusion applications are presented. The
application of ferritic steels for steam generator circuits and the main concerns in the weldments of
ferritic steels are discussed briefly. The future trends
in the application of ferritic steels in fast reactor
technology 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 predecessors, the creep-resistant ferritic steels, over nearly
a century. The first of the series was the carbon and
C–Mn steels with a limited application to about



Ferritic Steels and Advanced Ferritic–Martensitic Steels

Liquid

1500
Liq

Liquid + a + g

uid

a

1400

9Cr steel

1300

Temperature (ЊC)

1200

a
a+g
g

1100


1000

900

a + g + (CrFe)7C3

800
a + g + (FeCr)3C
700

600

a + (CrFe)7C3
+ (CrFe)4C

a + (FeCr)3C
+ (CrFe)7C3
a + (CrFe)7C3

a + (FeCr)3C
0

5

10

a + (CrFe)4C
15


20

25

Chromium (%)

(a)

Ae3
Ferrite
Temperature (ЊC)

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 economically 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, vanadium, 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 austenite and ferrite stabilizing elements. A typical pseudobinary 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 temperature and the decomposition mode7–9 of high
temperature austenite (Figure 1(b)) dictate the resulting microstructure at room temperature and hence, the
type of steel. Accordingly, the 9CrMo family of steels
can either be martensitic (9Cr–1Mo (EM10) or stabilized 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 result10 in 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 studied11,12 and 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

99


Pearlite
Ws
Bs

Widmanstatten ferrite
Upper
bainite

Bainite
Lower
bainite

Ms
Martensite

(b)

Log {time}

Figure 1 (a) Pseudobinary phase diagram for a Fe–Cr–C
steel with 0.01% C. Reprinted, with permission, from
High chromium ferritic and martensitic steels for nuclear
applications, copyright ASTM International, 100 Barr Harbor
Drive, West Conshohocken, PA 19428. (b) Decomposition
modes of high-temperature austenite during cooling.

to elevated temperature for prolonged durations have
been extensively studied.5,13,14 The following trends
were established: The optimized initial alloy composition considered was 9Cr, W–2Mo ¼ 3, Si ¼ 0.5, with
C, B, V, Nb, and Ta in small amounts. Higher chromium content has two effects: it increases the hardenability leading to the formation of martensite and

also promotes the formation of d-ferrite thereby
reducing the toughness. A reduction in the chromium


100

Ferritic Steels and Advanced Ferritic–Martensitic Steels

(Mn+Ni)/135
Mod. 9Cr 1Mo

Mod. 9Cr 1Mo: base model

(Mn+Ni)/1.85
Mod. 9Cr 1Mo

4

(Mn+Ni)/2.32
Mod. 9Cr 1Mo

3

(Mn+Ni)/1.7
Mod. 9Cr 1Mo

1800

Mod. 9Cr 1Mo


1805

8

9

10

1795

625

(a)

1810

0.6 Si added 9Cr 1Mo

650

0.24 Si added 9Cr 1Mo

675

Plain 9Cr 1Mo

Melting point (K)

Ms, experimental (K)


1815

700

600
600

Experimental
Estimated

1820

Ms/K = 904 - 474 (C + 0.46(N - 0.15Nb ) - 0.046Ta)
-{17Cr + 33Mn + 21Mo + 20Ni + 39V + 5W)
-45Mn2 - 25Ni2 - 100V2 + 10Co } - 44.5Ta

0.42 Si added 9Cr 1Mo

9Cr–ferritic martensitic steels
725

1W-0.23V-0.05Ta 9Cr 1Mo

750

1790
625

650


675

700

725

750

1785

Ms, empirical estimate (K)

1

2

(b)

5

6

7

11

Steel designation

1000
9Cr–ferritic steel


99

Continuous heating
transformation (CHT) diagram

1248

Austenite

60

Ac3

40

1198
20

900

10

50% transformed
Ac1

850

15


Ferrite+
austenite+
carbide

1

1098

Ferrite + carbide
800
100

101

(c)

102

1148

5

Temperature (K)

Temperature (ЊC)

950

103


Time (s)

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.

content lowers the oxidation resistance. If W þ Mo
concentration is kept <3%, creep strength will
reduce, while higher amount promotes the formation
of d-ferrite and brittle Fe2Mo Laves phase. The addition 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 cooling during start-ups and shutdowns or during

variations in operating conditions. Steady state
operation in between start-up and shutdown or
transients would produce creep effects. Therefore
the low cycle fatigue (LCF) and creep–fatigue interaction assume15 importance in the safe life design
approach of steam generator components. The alloy
exhibited a decrease in fatigue life with increasing
temperature, thus limiting its upper limit of temperature up to about 773 K.
The joining technologies of Cr–Mo steels have

been well investigated.16,17 One of the major problems during welding of ferritic steels has been the
formation of d-ferrite, if the amount of ferrite stabilizers is high. The partial substitution of Mo with
W enables austenite stabilization and hence reduces
the tendency to form d-ferrite. In fact, there needs
to be an intricate balance between austenite and
ferrite stabilizing elements in 9–12Cr–Mo steels.


Ferritic Steels and Advanced Ferritic–Martensitic Steels

This would ensure a satisfactory solidification process
with a fully austenitic structure. Additionally, this
enables easier hot workability during primary processing and tubemaking, without losing high temperature creep resistance. The formation of d-ferrite
reduces toughness due to the notch sensitivity, promotes 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 extensively studied.18 The Type IV cracking in ferritic steel
weldments and the brittle layer formation in the dissimilar 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 3 shows
a schematic of clad and wrapper in a typical fuel
subassembly. The necessity to develop robust technology for core component materials arises from the
fact that the ‘burn-up’ (energy production from unit

Head
Adaptor

Shield pins,
pellet stack

B&C
shielding
Steel
shielding
Top
plannet

Section−MM
217 fuel pins

Fuel
Bottom
plannet
Section−E E

Coolant
entry tube

Section−HH

Section−BB
Discriminator
Section–XX


Fuel Pin

Figure 3 Schematic of a typical fuel subassembly.

101

quantity of the fuel) of the fuel depends on the
performance of the clad materials. The higher burnup of the fuel increases the ‘residence time’ of the
subassembly in the core, eventually lowering the cost.
The core component materials in fast breeder
reactors are exposed to severe environmental service
conditions. The differences in the exposure conditions of the clad and wrapper in a fast reactor core
are listed in Table 1. Under such exposure conditions, materials in the fast reactor fuel assemblies
exhibit many phenomena (Figure 4), specific to fast
reactor core: Void swelling, irradiation growth, irradiation hardening, irradiation creep, irradiation, and
helium embrittlement.
Another selection criterion, namely the compatibility of the core component materials with the coolant, the liquid sodium, has already been established.
Presently, methods are known to avoid interaction of
the clad material with the coolant.
Detailed books and reviews19,20,21,22,23 are available on all the degradation mechanisms mentioned
above, which are related to the production, diffusion, and interaction of point defects in the specific
lattice of the material. Hence, a brief introduction is
presented below (see also Chapter 1.03, RadiationInduced Effects on Microstructure; Chapter 1.11,
Primary Radiation Damage Formation; and Chapter
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
1022 n cmÀ2. Void swelling is caused by the condensation of ‘excess vacancies’ left behind in the lattice

after ‘recombination’ of point defects produced during irradiation. Void swelling is measured using the
change in volume (▽V/V) of bulk components of the
reactor or image analysis of voids observed using
transmission electron microscope (TEM).
The ‘irradiation growth’ (fluence $1020 n cmÀ2)
can increase the length of a cylindrical rod of uranium
three times and reduce its diameter by 50%, retaining
the same volume. This occurs mainly in anisotropic
crystals, introducing severe distortion in core components. It is caused by the preferential condensation of
interstitials as dislocation loops on prism planes of
type (110) of hcp structures and vacancies as loops
on the basal planes (0001), which is equivalent to
transfer of atoms from the basal planes to prism
planes, via irradiation-induced point defects.
Irradiation hardening refers to the increase in
the yield strength of the material with a


102

Ferritic Steels and Advanced Ferritic–Martensitic Steels

Table 1

Comparison of exposure conditions of clad and wrapper of fast reactor core

Criterion

Clad tube


Wrapper tube

Exposure conditions (only trends; exact
values depend on core design)

Maximum temperature: 923–973 K

Lower temperature range than clad:
823 K
Lower temperature gradient
Moderate stresses from coolant
pressure
Flowing sodium environment
Neutron environment similar

Major damage mechanisms

Selection criteria: mechanical properties

Corrosion criteria

Steeper temperature gradient
Higher stresses from fission gas
pressure
Chemical attack from fuel
Average neutron energy: 100 keV
Neutron flux: 4–7 Â 1011 n mÀ2 s
Neutron fluence: 2–4 Â 1019 n mÀ2
Void swelling
Irradiation creep at higher

temperatures
Irradiation embrittlement
Interactions with fuel and fission
products
Tensile strength
Tensile ductility
Creep strength
Creep ductility
Compatibility with sodium
Compatibility with fuel
Compatibility with fission products

Void swelling
Irradiation creep
Irradiation embrittlement
Interaction with sodium
Tensile strength
Tensile ductility

Compatibility with sodium

General common selection criteria
Good workability
International neutron irradiation experience as driver or experimental fuel subassembly
Availability

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 additional stress to unpin the immobile dislocations.
The irradiation creep, the most important parameter 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 ‘stressinduced 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 during irradiation with dislocations.
Irradiation embrittlement, another frequent
observation in ferritic steels exposed to irradiation,
refers to the increase in the ductile to brittle transition temperature (DBTT) during irradiation. 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 to
brittle failure. Such materials like ferritics release
more dislocations into the system when a source is
unlocked, causing hardening and loss of ductility.
Some of the engineering materials contain nickel,
an element which undergoes an (n,a) reaction, producing high concentration of helium. The binding energy
of helium with a vacancy being very high $2 eV,
the helium atoms stabilize the voids, enhancing their
growth rate. Incorporation of helium during irradiation into voids along the grain boundaries assists
grain boundary crack growth by linking voids causing
‘helium embrittlement.’
Of these many degradation mechanisms, the alloy
development programmes have focused mainly on

the void swelling, irradiation hardening, embrittlement, and the irradiation creep, since these are the
major life limiting factors.

4.03.4 Development of Ferritic
Steels for Fast Reactor Core
This section begins with the optimization of chemistry and initial microstructure to develop swelling and


Ferritic Steels and Advanced Ferritic–Martensitic Steels

103

3
a axis

(a)

Growth strain (10-4)

Swelling (ÑV/V )

2
Linear
swelling
regime

Transient
swelling
regime


Threshold dose

Irradiation dose (dpa)

1
0
-1

c axis

-2
-3

(b)

Neutron fluence (E>1 MeV) 1024 n m–2

Strain (%)

Stress

Irradiated

Unirradiated

Irradiated

Unirradiated

(c)


(d)

Strain

Time (h)

Absorbed energy

Unirradiated

(e)

Irradiated

Test temperature

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.

creep-resistant ferritic steels. The microstructural
instability during service exposure is briefly presented. The superior swelling performance of ferritic
steels is understood based on mechanisms of void
swelling suppression. Following this, the irradiationinduced/-enhanced segregation/precipitation causing
irradiation hardening is discussed. The irradiation
creep and embrittlement, their mechanisms and methods to combat the problems are highlighted. The R&D
efforts of today to reduce the severity of embrittlement in ferritic steels, using modeling methods, are

outlined. Finally, typical problems in the weldments

of ferritic steels, when used for out of core applications, are presented, emphasizing the advantage of
modeling in predicting the materials’ behavior.
4.03.4.1 Influence of Composition and
Microstructure on Properties of Ferritic
Steels
Rapid strides have been made the world over, in the
design and development of advanced creep-resistant
ferritic or ferritic–martensitic steels. The low alloy
steels can be used as either 100% ferrite–martensite


104

Ferritic Steels and Advanced Ferritic–Martensitic Steels

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 about $0.1%. The
9Cr steels have always been used in the 100% martensite 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

Table 2
Optimizing the constitution in the development
of ferritic steels
Element

Function

Cr

Basic alloying element, corrosion
resistance, hardenability
Solid solution strengthening
Strengthening by formation of
MX-carbonitride
Austenite stabilizer, solid solution
strengthening, carbonitride
formers
Grain boundary strengthening,
stabilization of carbide
Austenite former, inhibits d-ferrite
formation

Mo, W, Re, Co
V, Nb, Ti, Ta
C, N


B
Ni, Cu, Co

Table 3

elements led24 to copious, uniform precipitation of
Monte Carlo (MC) type of monocarbides, which
are very fine and semicoherent. Such precipitates are
very efficient in pinning the mobile dislocations, leading to improved creep behavior at higher temperatures. These carbides are stable at temperatures
higher than even 1273 K and hence, do not cause
deterioration of long-term mechanical properties
during service exposure.
The development of high creep-rupture strength
9–12% steels with various combinations of N, Mo,
W, V, Nb, Co, Cu, and Ta is based on optimizing
the constitution (Table 2.) and d-ferrite content,
increasing the stability of the martensite, dislocation
structure and maximizing the solid solution and precipitation hardening. The concentration of each element in ferritic steels has been optimized based on
an in-depth understanding of the influence of the
specific element on the behavior of the steel. The
extensive studies related to optimization of chemistry
are summarized in Table 3. Based on the strong
scientific insights, large number of commercial steels
have been developed (Table 4) in the later half of
the last century.
Most of this family of ferritic–martensitic steels
is used in the normalized and tempered condition or
fully annealed condition to achieve the desirable phase.
The type of structure that is deliberately favored in a
given steel depends on the end application.

The microstructure of the steels in normalized
and tempered conditions consists24 (Figure 5) of (a)
martensite laths containing dislocations with a Burgers
vector 1/2a0<111> with a density of approximately
1 Â 1014 mÀ2 (b) coarse M23C6 particles located at

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

Element

Beneficial

Detrimental

Optimum (wt%)

Carbon

Strength

0.1

Mo

Creep strength

Ni
Mn
Si


S scavenger
Void swelling, spheroidization

Ti, V, Nb, Ta, and W

Precipitation strengthening

Weldability
Hardenability
d-ferrite formation
Laves phase
Intermetallic Ni3P
Induced radioactivity
d-ferrite formation
Silicide formation
Undissolved carbides, low
hardness of martensite
Radiation embrittlement
Precipitation
Induced radioactivity
Ductility
Neutron poison

S, P, As, Sb, Sn, and Bi
Cu
Co
N, O
B

Delayed coarsening


1
0.1
0.5
0.2–0.4
<0.1
<0.001
<0.01
<0.01
<20 ppm
<0.002


Ferritic Steels and Advanced Ferritic–Martensitic Steels
Table 4
List25 of commercial ferritic steels, their
chemistry, and properties
Commercial
name

Chemistry

105 h creep strength
at 873 K MPaÀ1

T22
Stab. T22
HCM2S
T9
EM12

F9
T91

2.25Cr1Mo
2.25Cr1MoV
2.25Cr1MoWNb
9Cr1Mo
9Cr2MoVNb
9Cr1MoVNb
9Cr1MoVNb
(optimized)
9Cr(MoW)VNb
9CrWTiV
12Cr1MoV
12Cr1MoWV
12CrMoWVNbCu
12CrWVNbCo

35
60–80
100
35
60–80
60–80
100

T92
Eurofer
HT91
HT9

HCM12A
SAVE12

120
$120
60–80
60–80
120
180

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. Prolonged exposure at high temperature causes changes
in the initial microstructure, which has been studied
extensively. The M2X precipitates in the normalized and tempered stabilized 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: irradiation-induced precipitation, irradiationenhanced transformations, and the irradiation modified 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 irradiationinduced phases and their influence on hardening
and embrittlement is discussed later.

4.03.4.2

105

Void Swelling Resistance

Extensive experimental investigations found3 that the
ferritic steels, whose high temperature mechanical
properties are far inferior to austenitic stainless steels,
displayed excellent radiation resistance. The ferritic–
martensitic steels (9–12% Cr) have, therefore, been
chosen for clad and wrapper applications, in order
to achieve the high burn-up of the fuel. This is
based26–29 (Table 5) on their inherent low swelling
behavior. 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 ferritic
steels is as high as nearly 200 dpa in contrast to 80 dpa
for the present generation D9 austenitic stainless
steel. It is established that the void swelling depends

crucially on the structure of the matrix lattice, in
which irradiation produces the excess defects.
Extensive basic studies have identified19,30–33 the
following reasons as the origin of superior swelling
resistance in ferritic steels:
1. The relaxation volume for interstitials, that is, the
volumeof the matrix inwhich distortion is introduced
bycreating an interstitial, in bcc ferrite is larger19 than
fcc austenite. For every interstitial introduced, the
lattice distortion is high and hence the strain energy
of the lattice. Hence, the bias toward attracting or
accommodating interstitials in the bcc lattice is less.
This leaves higher density of ‘free’ interstitials in the
bcc lattice than fcc lattice. As a result, recombination probability with vacancies increases significantly
and supersaturation of vacancy reduces. Consequently, the void nucleation and swelling is less.
2. The migration energy of vacancies in bcc iron is
only 0.55 eV, against a high value in fcc austenite,
1.4 eV. Vacancies are more mobile in bcc than
fcc, increasing the recombination probabilities in
bcc ferrite. Another factor is the high binding
energy between carbon and vacancy in bcc iron
(0.85 eV), while it is only 0.36–0.41 eV in austenite.
This leads19 to enhanced point defect recombination in bcc than fcc, due to more trapping of
vacancies by carbon or nitrogen.
3. In bcc iron, it is known30 that there is a strong interaction between dislocations and interstitials solutes,
forming atmospheres of solutes around dislocations.
The formation of ‘atmospheres’ around dislocations
makes them more effective sinks for vacancies than
interstitials, resulting in suppression of void growth,



106

Ferritic Steels and Advanced Ferritic–Martensitic Steels

A

B

0.5 mm

(a)

(001)
(b)

V Ka

NbLa

V Kb

NbKa

FeKa
2.00

NbKa

V Ka


NbLa

4.00

6.00

8.00

10.00

12.00

14.00

16.00

(c)

1

2.00

4.00

6.00

8.00

10.00


12.00

14.00

16.00

(d)

Figure 5 Initial structure24 of normalized and tempered modified 9Cr–1Mo steel: (a) Monocarbides (MC) and M23C6
along lath boundaries in a carbon extraction replica of the sample and (b) Microdiffraction of fine particle marked B,
confirming the crystal structure of MC. Energy dispersive analysis of X-rays (EDAX) identifying the MC particles (B) to be rich in
(c) V and (d) Nb.

provided the following conditions are satisfied: ‘atmospheres’ 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 probabilities, 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 use

ferritic steels directly in fast reactor fuel assembly
was with respect to enhancing the high temperature
mechanical properties of the ferritic steels, especially
high temperature creep life and irradiation creep
resistance.
4.03.4.3 Irradiation Hardening in
Ferritic Steels
The initial microstructure of the steels evolves during service, due to high temperature and irradiation
for prolonged times, leading to modification of defect
structure and secondary phases. These changes
harden the steel, leading to concomitant embrittlement, which is discussed below.


Ferritic Steels and Advanced Ferritic–Martensitic Steels

It is reported that carbon content in 12% chromium steel is maintained high in order to use the
steel as martensitic steels. The high amount of carbon in 12% chromium steel leads to copious precipitation of carbides, that is, twice as much in 9Cr
steels. Both the steels have predominantly M23C6
carbides with a small fraction of monocarbides,
(1 21 3)

(a)

1 mm

MoLa

Laves phase

FeKa


(b)

CrKa
SiKa
MoKa

2.00

4.00

6.00

8.00

10.00 12.00 14.00 16.00 18.00

24

Figure 6 Effect of 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.

Table 5

107


eventually leading34 to deterioration of their resistance to brittle failure. The critical stress to propagate a crack is inversely proportional to the crack
length. If it is assumed that fracture initiates at an
M23C6 precipitate and the crack length at initiation
equals the diameter of a carbide particle then the
fracture stress will decrease with increasing precipitate size. The precipitates coarsen during irradiation
in the range of 673–773 K, thus causing a decrease
in fracture stress and an increase in DBTT even in
the absence of further hardening.
Additionally, Cr rich, bcc a0 precipitates formed35
in the higher chromium steel during thermal exposure and irradiation lead to hardening and embrittlement of the steel. The d-ferrite, into which there is a
repartitioning of chromium, is harmful, since it promotes formation of a0 . The presence of very fine
coherent particles of the w (Fe2Mo) phase has also
been reported in the T91 and HT9 steels. The w phase
was observed to form more rapidly in the 9Cr–2Mo
type of steels, both in the d-ferrite and martensite
phases. This is possibly due to the higher amount of
Mo in the EM12 type of steels. The w phase is
enriched in Fe, Si, and Ni and contains significant
amount of Mo and P. The G phase (Mn7Ni16Si17)
has been found to form very occasionally in the modified 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 M23C6
particles in 9–13% Cr martensitic steels after irradiation at 420–460  C in Dounray Fast Reactor. In
addition Cr3P needles and MP (M ¼ Fe, Cr, and Mo)
particles have also been detected in the 12 and 13Cr
steels in the range of 420–615  C. The formation of
these phases during irradiation may be understood in
terms of the strong radiation-induced segregation
(RIS) of alloying and impurity elements to point defect

sinks in the steels (see Chapter 1.18, RadiationInduced Segregation). The RIS of alloying/impurity
elements could lead36 to either enrichment or depletion near the sinks, depending on the size of the atom
and its binding energy with iron self-interstitials.

Void swelling resistance26–29 of some commercial ferritic steels

Commercial
name

Chemistry and country of origin

Reactor in which irradiation
was carried out

Burn-up achieved (dpa)

FV448
EM10
1.4914
EP450
EP450

12Cr–MoVNb, UK
9Cr–1Mo, France
12CrMoVNb, Germany
12Cr–MoVNb, Russia
–

PFR
Phenix

Phenix
BN-350
BN-600

132
142
115
45
144


108

Ferritic Steels and Advanced Ferritic–Martensitic Steels

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-interstitial. 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 elements happen to be smaller than iron.
The w, G and s phases are all enriched in Si and
Ni – elements which are known to segregate to interfaces 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 concentrations 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 improves the ductility and toughness of the 12Cr steel,
the fracture could be initiated at the M23C6 precipitates on the d-ferrite–martensite interface. The
presence of d-ferrite, extensive precipitation and
radiation-induced growth of M23C6 precipitates and
formation of the embrittling intermetallic phases in
the 12Cr–1MoVW steel in the temperature range
573–773 K are together responsible37 for 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.
4.03.4.4 Irradiation Creep Resistance of
Ferritic Steels
An essential prerequisite for maximizing the ‘irradiation creep resistance’ is to ensure38 the best
combination of thermal creep behavior and longterm 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 creepresistant steels are as follows:
 Introduce high dislocation density by either transformations or cold work to increase the strength of

the basic lattice;
 Strengthen the host lattice by either solid solution
strengtheners or defects;
 Stabilize the boundaries created by phase transformations by precipitating carbides along the
boundaries;
 Arrest dislocation glide and climb by appropriate
selection of crystal structure, solid solution, interfaces, dislocation interactions, and crystal with low
diffusivity;
 Resist sliding of grain boundaries by introducing
special type of boundaries and anchoring the
boundaries with precipitates;
 Ensure long-term stability of the microstructure,
especially against recovery and coarsening of the
fine second phase particles;
In the case of 9–12 Cr steels, the martensitic lath
structure (Figure 7) decorated with only MX which
should39 be stable over long-term service life is
the most desired structure. Thermo-Calc evaluations show39 that MX can be stabilized at the expense
of M23C6 only by reducing carbon to as low a value
as 0.02% in 9 Cr–1Mo steel. This value is too
low to ensure acceptable high temperature mechanical behavior of the steels. In the context of fast
reactor core components, the high chromium 9–12%
ferritic–martensitic steels assume relevance. Hence,
an extensive database25 for a large number of
commercial ferritic steels has been generated and
Lath
boundary

Lath
boundary


(a)

(b)
M23C6
MX

Figure 7 The schematic39 of most undesirable (a) and
desirable (b) microstructures for design of creep-resistant
steels.


Ferritic Steels and Advanced Ferritic–Martensitic Steels
stensile
-0.5Mo + 0.8 W

100

0
s

T
9Cr1Mo

P91

E911

P92


T

0

0

650 ЊC

50

T

600 ЊC

+ 0.04N + 0.2V
+ 0.08Nb

(a)

0 0 0 0

+1 W

T

105 h rupture stress (MPa)

150

109


T

stensile
0

0

5
540 ЊC
1 ϫ 1023 n cm–2

4

0

(a)
0

316SS

0

DD/D (%)

0
3
D9
2


(b)
Figure 9 The mechanisms of stress-induced
preferential absorption (a) and stress-induced preferential
nucleation (b) during irradiation creep.

HT-9
D21

1

D68
0
0

(b)

20

40

60

80

100

120

140


Hoop stress (MPa)

Figure 8 (a) Thermal creep40 of 9Cr1Mo ferritic steel.
(b) irradiation creep41 of ferritics in comparison to
austenitics. Reprinted, with permission, from J. ASTM Int.,
copyright ASTM International, 100 Barr Harbor Drive,
West Conshohocken, PA 19428

Figure 8(a) shows40 the continuous improvement
achieved by careful modification of alloying elements,
in the thermal creep behavior of successive grades
of different commercial ferritic steels.
While understanding thermal creep is essential
to narrow down the choice of ferritic steels for use
in a fast reactor, ‘benchmarking’ the steels developed
under irradiation is an essential stage before actually
using the radiation-resistant steels in the reactor.
The irradiation creep depends on the stress level,
the temperature, and the dose. Figure 8(b) shows41
the comparison of irradiation creep of ferritic steels
with competing materials like the austenitics and
nickel-based alloys.
It is clear that the point defects generated during
irradiation act against the design principles of developing creep-resistant materials, listed earlier. The
point defects accelerate the kinetics of dislocation
climb, coarsen the precipitates, and generally enhance
the diffusivity. In addition, the excess point defects
precipitate into either interstitial or vacancy loops,
but not randomly. The interaction between point
defects and stress leads to the precipitation of interstitial loops parallel to the applied stress, while vacancy


loops form in planes perpendicular to the stress. This
process (Figure 9(a)) called the stress-induced preferential nucleation (SIPN) results in additional creep
strain solely due to irradiation. The excess point
defects under temperature migrate randomly. But in
the presence of an additional factor, that is, stress, the
vacancies migrate preferentially to grain boundaries
perpendicular to the applied stress, while the interstitials toward boundaries parallel to the stress. This is
equivalent to removing material from planes parallel
to the stress to those which are perpendicular to
the applied stress, introducing additional creep strain.
This process is called the stress-induced preferential
absorption (SIPA) (Figure 9(b)).
The radiation-induced defects also evolve from
isolated point defect to loops and voids, which have
different types of influence on irradiation creep. Most
often, irradiation creep occurs19,42 simultaneously with
swelling and sometimes, swelling influences irradiation creep. At very small dose levels, swelling enhances
creep rates. Beyond a certain dose levels, the creep
component reduces and at high dose levels, creep disappears, while swelling continues. Figure 10 shows the
variation in creep coefficient at various dose levels, and
the regimes where swelling has an influence. The
dynamics of point defects during irradiation continuously evolve with change in structure of dislocation
network and loops. At small dose levels, there is a
uniform distribution of very fine voids, which act as
effective pinning centers for mobile dislocations. Thus
the creep rate increases. With increase in dose levels,
voids grow and multiply. The chance of interstitials and



110

Ferritic Steels and Advanced Ferritic–Martensitic Steels

Instantaneous
creep coefficient

Onset of
disappearance of creep

Swelling without
creep

Swelling
enhanced creep

Dose (dpa)

Figure 10 Schematic of variation of instantaneous creep
coefficient with dose, showing the interplay between
irradiation creep and void swelling.

vacancies impinging on the void surface becomes more
than their reaching dislocations. The number of interstitials reaching a dislocation reduces. Additionally, the
defect clusters, that is, the dislocation loops also
undergo ‘faulting’ contributing to the density of dislocations in the matrix. Hence, creep rate reduces, due
to two factors: increased dislocation density of the
matrix due to unfaulting of dislocation loops and
reduced availability of interstitials to dislocations.
The above process continues until complete cessation

of creep, with swelling continue to take place.
At very high temperatures, the point defect migration along the grain boundaries in preferential routes
causes the grain boundary aided creep.
This high temperature limit of ferritic–martensitic
steels restricts the application of these steels to at
best, wrappers of present generation fast reactors
based on oxide fuel. It is necessary to develop materials with better high temperature irradiation creep
properties and void swelling for clad applications.
The future scenario, which envisages the development
of metallic fuel to ensure sustainability by breeding,
could make use of ferritic steels for both clad and
wrapper. This advantage arises due to the lower value
of the anticipated clad temperatures with metallic fuels
(see Chapter 3.14, Uranium Intermetallic Fuels
(U-Al, U-Si, U-Mo)), whose choice is mainly to
ensure sustainability using high breeding ratio.
4.03.4.5 Irradiation Embrittlement in
Ferritic Steels
The stabilized ferritic steels in the normalized and
tempered condition have a tempered martensitic
structure with a preponderance of monocarbides

that impart the necessary creep strength, while the
prior austenite grain and lath boundaries are decorated with Cr rich M23C6 precipitates which increase
the thermal stability of the steel. It is reported that
thermal aging at temperatures above 773 K causes
gradual but continuous degradation in upper shelf
properties in addition to increase in the DBTT. The
nature of embrittlement varies for different components of the reactor. For removable components
such as clad, which are subjected to high temperature

and pressure, with a residence time of a few years,
creep embrittlement is the issue which decides their
design and performance, while for permanent support structures increase in hardening and loss in
fracture toughness on irradiation are major issues.
The origin of embrittlement is two-fold: segregation of tramp elements to prior austenite grain
boundaries which make the grain boundaries decohesive and evolution of carbides and intermetallic
phases. The latter causes progressive changes in the
tempered martensitic microstructure, which deteriorate the fracture properties of the steel, by introducing irradiation hardening effects.
The increase in the ductile to brittle transition
temperature, DDBTT, is known to be related to
irradiation hardening, which is generally observed to
saturate with fluence. Evidence for a possible maximum in DBTT was observed for the 12Cr steel irradiated in the range of 35–100 dpa in fast flux test facility
(FFTF). Based on observed data in a number of cases
it appears that a high fluence and/or high temperature are required before a maximum is observed. This
implies that the strength and impact properties are
a balance between the point defect production and
irradiation-induced precipitation. The precipitation
during irradiation hardens the steel and irradiation
accelerated recovery and aging soften the steel. The
latter process is more important at high fluences and/or
higher irradiation temperatures. Hence, hardening in
most of these Cr–Mo steels is more than compensated
for by the recovery and aging processes, leading to
saturation in irradiation hardening above 723 K.
For body centered cubic materials such as ferritic
martensitic steels, radiation hardening at low temperatures (<0.3TM) can lead to a large increase in the
DBTT and lowering of impact energy for radiation
dose as low as 1 dpa (displacement per atom). The
minimum operating temperature to avoid embrittlement in ferritic martensitic (F/M) steels is $473–
523 K, while the upper limit is controlled by four

different mechanisms: thermal creep, high temperature
helium embrittlement, void swelling, and compatibility


Ferritic Steels and Advanced Ferritic–Martensitic Steels

with fuel and coolant. Void swelling is not expected to
be significant in F/M steels up to damage levels of
about 200 dpa.
Extensive evaluation14,15,43–58 of the embrittlement behavior of the ferritic steels for different
chemistry is shown in Figure 11. The merit in focusing on chemistry around 9% chromium is very clear
based on the observation of minimum shift in DBTT
around this composition, under irradiation. However,
higher chromium improves corrosion resistance and
ease of reprocessing. Hence, chromium content has
to be selected balancing these requirements. It is
known44–48 that addition of phosphorous, copper,
vanadium, aluminum, and silicon would increase the
DBTT while sulfur reduces the upper shelf energy
(USE). The 12Cr steels, HT9, show a larger shift
(125 K) in DBTT as compared to modified 9Cr–1Mo
steel ($54 K). Hence, the balance is always between
nearly nil swelling resistant 12Cr steels and 9Cr steel
which is less prone to embrittlement than 12Cr steels.
Microstructural parameters, like the prior austenite grain size, lath/packet size, carbides, and their
distribution influence49,50 the embrittlement behavior. Studies on the effects of heat treatment and
microstructure on the irradiation embrittlement in
9Cr–1MoVNb and HT9 steels are summarized
below:









 Prior austenite grain size (PAGS) influences51 the
DBTT for the 9Cr–1MoVNb steel, but not in

250
JLF-4

DBTT shift (C)

200

150

12Cr–MoVW steel. This is attributed to the
precipitates in the microstructure controlling
the fracture behavior rather than the PAGS, in the
12Cr steel.
The size of martensitic lath and packet, which is
sensitive52 to austenitization temperature, can also
affect51 the fracture behavior. Examination of the
fracture surface revealed cleavage and regions of
ductile tearing along prior austenite grain and lath
packet boundaries. Subsurface microcracks and secondary surface cracks were found associated with
large boundary carbides. It was suggested that cleavage fracture initiated in HT9 by propagation of a

microcrack from a coarse carbide into the matrix.
Propagation was inhibited by the intercepted
boundaries, lath or grain and ductile tearing was
required53 to continue propagation. The amount of
tearing increased with increasing austenitization
temperature.
Tempering for the two normalization temperatures had very small effect on the DBTT, for the
two steels.
Irradiation of the two steels at 638 and 693 K
resulted37 in an increase in DBTT and a decrease
in USE for all conditions with the shift in DBTT
for the 12 Cr steel being almost twice that for
9Cr steel.
Although the 12Cr steel with the smallest
grain size had55 the lowest DBTT after 20 dpa,
the effect of tempering was different. In the case

2.25CrV
2.25Cr–1WV
10 dpa: 365 ЊC
2Cr–1.5V
R. Klueh et al.

JLF-6
12Cr–6Mn–1W
12Cr–6Mn–1V

36 dpa: 410 ЊC
A. Kohyama et al.


2.25Cr–2W

111

12Cr–2WV

JLF-3

100

2.25Cr–2W

F82H

5Cr–2WV

9Cr–2WV

7 dpa: 365 ЊC
10 dpa: 365 ЊC

50

7 dpa: 365 ЊC
R. Klueh et al.

36 dpa: 410 ЊC

Cr–1V
JLF-1

9Cr–1W

7.5Cr–2W

9Cr–2WVTa

0
0

2

4

6

8

10

12

Chromium content (wt%)
Figure 11 Variation43 of shift in ductile to brittle transition temperature (DBTT) for various Cr–Mo steels with irradiation
to different dose levels at around 673 K. The ferritic steel with 9Cr–1Mo has the least variation in DBTT.


112

Ferritic Steels and Advanced Ferritic–Martensitic Steels


of 12Cr steel, the higher tempering temperature
causes coarsening of precipitates thus accelerating
fracture.
 The saturation of shift in DBTT with fluence is
independent54 of tempering conditions for the 9Cr
steel, while for the 12Cr steel, a maximum is
observed, probably due to faster growth of precipitates during irradiation.
The generation of helium through (n,a) reaction in
elements of structural materials is known to cause
severe damage to the embrittlement behavior of core
component materials. Table 6 lists the shift in
DBTT, for 9 and 12CrMo steels, under reactor irradiation, with and without helium, which demonstrates56 the harmful effect of helium. These results
become more pertinent in the case of fusion reactors,
where the operating conditions include the generation of helium up to about $100 appm yearÀ1.
The increase in the DBTT due to irradiation is a
cause of serious concern for use of ferritic steels,
since it makes the postirradiation operations very
difficult. Several methods have been attempted57,58
to address this problem, which includes modification
of the steel through alloying additions, control of
tramp elements by using pure raw materials and
improved melting practices, and grain boundary
engineering (GBE). However, the propensity of the
problem is less if the clad thickness is low, which
normally is the case to ensure best heat transfer
properties. For low thickness components, the triaxial
stress necessary for the embrittlement does not
develop, which reduces the intensity of this otherwise
serious problem of embrittlement in ferritic steels.
An approach to reduce shift in DBTT is an

immediate concern in ferritic steels for core component applications and efforts to overcome this
problem by selection of high purity metals, adoption
of double or triple vacuum melting for steel making,
strict control of tramp and volatile elements, and
development of special processing methods, which
would improve the nature of grain boundaries
(GBE) are in progress.
Table 6

4.03.4.5.1 GBE to reduce embrittlement
in ferritic steels

GBE is an emerging field, which promises methods
to improve the performance of materials, whose degradation in service is caused by the presence of
high angle boundaries. The concept, first proposed59
by Prof. T. Watanabe in the early 1980s, envisages
improvement of properties of materials by controlling
the grain boundary character distribution (GBCD).
Many processes like diffusion, precipitation, segregation, sliding, cavitation, and corrosion are kinetically
faster along high angle grain boundaries. Hence,
it is possible to decelerate these detrimental processes by replacing the random boundaries with low
energy ones, coincident site lattice (CSL) boundaries
(denoted by the ‘sigma number,’ S, which is defined as
the reciprocal of the fraction of lattice points in the
boundaries that coincide between the two adjoining
grains on the basis of CSL model). Another prerequisite for GBE is to completely destroy the interconnectivity of random grain boundary network. The insight
in the field of GBE was achieved with the advent
of computer assisted EBSD (electron back scatter
diffraction) technique developed during the 1980s.
The embrittlement in ferritic steels is known to

be caused by segregation phenomena. The kinetics of
segregation can be controlled by suitable selection
of the nature of grain boundaries. GBE has been
applied60–63 to combat embrittlement problems in
ferritic steels. The task of carrying out GBE using
experimental methods is time consuming. Hence, it is
prudent to resort to computational methods, which
need to be validated using selected experiments.
A 3D Poisson–Voronoi grain structure, simulated
using MC technique was employed to study60
(Figure 12(a)) intergranular crack percolation using
percolation theory. The percolation threshold was
estimated to be 80%. To apply this model to specific
alloys like ferritic steel, system specific characteristics need to be incorporated61 in the model. One such
attempt is to define the propensity of the grain
boundaries for propagation of cracks based on relative values of the grain boundary energy and the

Comparison56 of embrittlement behavior of 9 and 12Cr steels, with and without helium

Irradiation conditions

Shift in DBTT (K)

Reactor

Temperature (K)

Dose (dpa)

9Cr1Mo(VNb)


12Cr1Mo(VW)

EBR II
EBR II
HFIR

663
663
673

13
26
40

50
50
200 (30 appm He)

125
$150
250 (110 appm He)


Ferritic Steels and Advanced Ferritic–Martensitic Steels

113

(a)


1
0.9
0.8
0.7

Failure probability

30
50
60
70
80

Fine
grains
(12 mm)

0.6

Coarse
grains
(25 mm)

% crackresistant
boundaries

0.5
0.4
0.3
0.2

0.1
0
200

300

400
500
600
Critical crack length (mm)

250

Two-step normalization and
tempering treatment

200
150

Conventional
N&T treatment

100

67.8 J lower
bound criteria

50
0


-45 ЊC

-30 ЊC

-80 -70 -60 -50 -40 -30 -20 -10

(c)

Fractal dimension increment

Charpy absorbed energy (J)

300

Temperature (ЊC)

0

10

20

30

0.185
0.180
0.175
0.170
0.165
0.160

0.155
0.150
0.145
0.140
0.135
0.130
0.125

0

5

800

15

10

900

20

1000

25

30

35


40

25

30

35

40

45

Average
Crack initial stage

0

(d)

700

Specimen A

100

Specimen B

0

(b)


5

10

15

20

0.185
0.180
0.175
0.170
0.165
0.160
0.155
0.150
0.145
0.140
0.135
0.130
0.125

45

Charpy impact energy (J)

Figure 12 Modeling and electron back scattered diffraction studies in grain boundary engineering of ferritic steels:
(a) Percolation of a crack in the 3D P–V model of grain structure generated60 using Monte Carlo methods. (b) Percolation
probability for two different grain size (c) experimental confirmation63 of reduction in ductile to brittle transition temperature

(DBTT) with grain size and (d) fractal analysis63 of the fracture surface revealing the tortuous path being responsible for
the improvement of DBTT in fine grain size. N and T refers to normalized and tempered.


114

Ferritic Steels and Advanced Ferritic–Martensitic Steels

energy required for propagation of cracks. These
calculations were carried out (Figure 12(b)) for two
different grain sizes. The prediction of finer grain size
being favorable to reduce embrittlement was confirmed (Figure 12(c)) experimentally. The GBCD,
that is, the distribution of various grain boundary
types has been evaluated62 in modified 9Cr–1Mo ferritic steel using EBSD technique. The experimental
observations confirmed the reduction of DBTT by
20 K with reduction in grain size. The fractal analysis
of the fracture surface demonstrated (Figure 12(d))
that the tortuous path which cracks need to follow
in fine grain sample is responsible63 for the observed
reduction in the propensity for embrittlement.
It is shown clearly that the low energy boundaries
can be introduced in engineering materials in three
different methods: preferential nucleation of low
angle boundaries around twins or controlled recovery or orientation relations during phase transformation, if some of the variants happen to result in CSL
boundaries. Significant improvements in properties
using GBE have been achieved64 in many austenitic
stainless steels, in contrast to ferritic steels. The major
challenges in the application of GBE to Cr–Mo ferritic steels arise from the following factors: lower
twinning probability, higher stacking fault energy,
and limited variants with CSL boundaries during

g ! a transformation during cooling.

4.03.5 Development of Advanced
ODS Ferritic Steels
In recent years, an attempt to increase the high temperature creep life of ferritics to 973 K and target
burn-up of the fuel to 250 dpa, has enabled a ‘revisit’
to the concept of strengthening the steel using 5 nm
particles of yttria (see Chapter 4.08, Oxide Dispersion Strengthened Steels), leading to the ODS ferritic steels. ODS ferritic steels are prospective
candidate materials for sodium cooled fast reactors
with peak burn-up of 250 dpa as well as GenIV and
fusion reactors. Earliest developments of ODS steels
can be traced to the efforts65 of Belgium in 1960s,
followed by Japan66 since 1987, and France67 in the
last decade. The ODS steels for fast and fusion reactors68,69 are in the R&D stage.
The design of ODS steels for fast and fusion
reactor applications is based on Fe–Cr–W–Ti–
Y2O3, either the martensitic 9 or 12Cr or the ferritic
12Cr steels. The dispersoids which confer the high
temperature creep life to the ferrite matrix are70

10 nm
Figure 13 Z-contrast in the high angle annular dark field
(HAADF) micrograph of dispersoids in oxide dispersion
strengthened (ODS)-9Cr–1M0 ferritic steels, which are
responsible for the superior high temperature creep behavior.

(Figure 13) in the size range of around 5 nm with a
volume fraction around 0.3%. The yttria dissolves in
it some amount of titanium, leading to the formation
of mixed, complex oxide, namely TiO2ÁY2O3.

The rationale for the choice of the matrix composition is as follows:
Chromium: Choice of 9% Cr and 0.1% C ensures
100% martensite, during normalization of the steel. It
is possible to ensure 100% martensite in 12% chromium steel by ensuring the carbon content to be
above 0.1%. Ferritic ODS steels can be obtained in
12% chromium steels by lowering the carbon content
to be less than 0.03%. Higher chromium provides the
corrosion and decarburization resistance in sodium at
973 K, with acceptable oxidation resistance.
Carbon: Addition of 0.1% carbon ensures 100%
martensite in 9% Cr steels, thus ensuring absence
of anisotropy during g ! a transformation. Higher
amount of carbon would promote precipitation of
M23C6, thus reducing the toughness. On the other
hand, M23C6 along the lath boundaries offers the
long-term microstructural stability of the lath
structure.
Nitrogen: The solubility of nitrogen in ferrite is
very low. This is useful in non-ODS ferritic steels
like T91, due to enhanced creep resistance by formation of V or Nb carbides/carbonitrides. But, in ODS
steels, Ti is used for refining yttria. Hence, nitrogen
content is restricted to 0.01%, preventing the formation of deleterious TiN compound.


Ferritic Steels and Advanced Ferritic–Martensitic Steels

Tungsten: Tungsten is a more effective solid
solution strengthener than Mo, but at the cost of
ductility. Tungsten stabilizes d-ferrite and accelerates
formation of Laves phase, both of which cause reduction in toughness. Hence, it is optimized to 2.0%.

Yttria: The most important constituent of ODS
steels is the yttria, which enhances high temperature
creep strength by pinning mobile dislocations and
delays void swelling by acting as sinks for point
defects produced during irradiation. The strength
increase is accompanied by a concomitant loss of
ductility and saturates around 0.4% yttria. Hence, it
is optimized to 0.35%.
Titanium: The major role of titanium in ODS
steels is to refine the yttria particles (20 nm after
mechanical alloying) to ultra-fine (2–3 nm) particles.
The complex Y–Ti–O particle imparts the necessary
high temperature creep strength. The beneficial
effect of titanium saturates around 0.2%. Further
increase introduces manufacturing problems of the
tubes and hence titanium is chosen as 0.2%.
Excess oxygen: Oxygen is present during processing
of the ODS steels. The oxygen present in excess of
the amount required for formation of required
amount of Y–Ti–O complex leads to increase in
tensile and creep strength. The Y–Ti–O complex
oxides requires about 0.07 þ 0.01% excess oxygen.
Argon: A strict control of argon (<0.002%) during
processing of ODS steels is essential to avoid embrittlement due to formation of argon bubbles during
irradiation.
Minor elements: Nickel and manganese are to be
reduced to maintain the A1 (a ! g on heating) temperature higher than the anticipated hot spot temperature. This enables the tempering temperature to be
as high as possible. Silicon, phosphorous, and sulfur
undergo RIS and cause embrittlement. Silicon also
accelerates formation of deleterious phases like the

Laves phase. Hence, their amounts are reduced to
0.05 and less.
The processing route used worldwide is the
powder metallurgy route of mechanically alloying
prealloyed powders of Fe–Cr–W–Ti–C þ Ti2O3,
followed by hot extrusion and rolling or hipping
with final heat treatments. Commercial ODS steels
(Table 7) have been developed demonstrating the
standardizing of fabrication technologies.
The ferritic–martensitic ODS steels have been
developed by adjusting the contents of chromium
and carbon. The ferritic ODS steels, with 12% chromium and carbon content less than 0.02%, derive71
their high temperature creep strength basically from

115

Table 7
List of few commercial ODS ferritic steels and
their chemistry
Commercial name

Chemistry

MA956
MA957
M11
M92
PM2000

Fe–20Cr–4.5Al–0.5Y2O3

Fe–14Cr–0.3Mo–Ti–0.27Y2O3
Fe–9Cr–Mo–0.37Y2O3
Fe–9Cr–Mo–0.30Y2O3
Fe–20CrAlTi–0.5Y2O3

the dispersoids. The ferrite matrix offers72 superior
resistance against oxidation and corrosion while the
major challenge appears to be the anisotropy73 of
properties. The martensitic steels based on either
12% chromium with 0.1–0.2% carbon or the 9%
chromium, derive their strength from the martensitic
matrix and the dispersoids. The 9% chromium steel
displays isotropic properties while it suffers from
inferior corrosion resistance.
The conventional joining technologies pose significant problems leading to coalescence of oxide
particles. Hence, solid state bonding techniques74
like the pressurized resistance welding have been
developed for joining the clad tube with the end
plug of a fast reactor. Postweld heat treatments
(PWHT) have also been developed to match the
strength levels of the clad and the end plug.
The in-service performance of ODS steels in
flowing sodium has been found to be satisfactory,
despite the formation of austenite layer on the surface of the clad tube due to the deposition of nickel.
The thickness of the oxide layer in 12%Cr ODS steel
was found to be only 50% of that in 9Cr ODS steels.
Reactor irradiation experiments have been performed75,76 on ODS ferritic steels. The dispersoids
were found74 to be stable up to a dose of 10 dpa in
JOYO. The mechanical properties at 573 K after
neutron irradiation were reported to be same as conventional ferritic martensitic steels. The studies on

long-term in-service behavior and postirradiation
behavior are being studied.
Presently, the main challenges in this variety of
ODS steels are the anisotropy observed in steels with
chromium more than 12%, less oxidation resistance in
steels with 9% chromium, fabrication procedure with
cost-effectiveness, uniformity of dispersoids in all
regions of the clad tube, stability of the dispersoids
under irradiation, and the joining technologies. It is
hoped that the above problems would be overcome in
the near future and the ODS ferritic steels can be used
in fast reactor core as clad and wrapper, for burn-up


116

Ferritic Steels and Advanced Ferritic–Martensitic Steels

beyond 250 dpa, at temperatures exceeding 973 K.
Additionally, ODS ferritic steels are also being considered for fusion reactor applications. The rich experience in the development of fast reactor materials
would enable launching the advanced ferritic steels
for fusion technology, in a shorter time span.

The HAZ comprises coarse prior-austenitic grain
martensite, fine prior-austenitic grain martensite
and an intercritical structure, as one traverses from
the weld fusion interface toward the unaffected base
metal. This is dictated by the peak temperatures
experienced by the base metal during weld thermal
cycle and the phase transformation characteristics

of the steel. It has been established that the localized
microstructural degradation in the intercritical
region of HAZ is mainly responsible for the premature creep-rupture strength of Cr–Mo weld joint and
can be overcome if residual stresses of the weld are
adequately relieved by PWHT.
The lower creep-rupture strength of weld joint
than the base metal is due38,77 to the different types
of cracking developed during creep exposure. Four
types of cracking have been identified (Figure 15)
in Cr–Mo steel weld joint. They have been categorized as Type I, Type II, Type III, and Type IV. The
Type I and Type II cracks originate in the weld
metal, propagate either through the weld metal itself
(Type I) or cross over in the HAZ (Type II). The
Type III cracking occurs in the coarse grain region of
HAZ and can be avoided by refining the grain size.
Type IV cracking nucleates and propagates in the
intercritical/fine grain region of HAZ. Type IV failure occurs at longer creep exposure and higher test
temperature, by coalescence of fine cavities leading
to microcracks (Figure 16(a)) and their eventual
propagation to the surface.

4.03.6 Ferritic Steels for Out-of-Core
Applications: Improvements in Joining
An ambitious target of increasing the temperature
and pressure of steam in many power plants has
provided a high impetus for the development of
steels with better high temperature properties. Very
often, the weld joints play a crucial life limiting role
in these components. One of the recurrent problems
is the frequent failure of weldments due to Type IV

cracking (see below), in weldments of ferritic steels
subjected to creep loading. Another problem encountered during service exposure of joints of dissimilar
ferritic steels is the failure due to the formation of
hard brittle zone at the heat-affected zone (HAZ).
Both these issues are discussed below.
The modified 9Cr–1Mo steel fusion weld joint
(Figure 14) consisting of base metal, deposited weld
metal, and the HAZ produces a complex heterogeneous microstructure due to thermal cycle. The base
metal and weld metal consist of a tempered martensite structure, with columnar grains in the weld metal.

Interface

Weld metal

Base metal

e

rfac

Inte
9Cr–1Mo
weld
metal

21/4Cr–1Mo
Base metal
20 mm

20 mm


Base
metal

ICZ

Weld
metal

FGH
AZ

CG
HAZ

20 mm

HAZ
CGHAZ

FGHAZ

ICZ

20 mm

20 mm

Figure 14 Schematic showing different zones in a ferritic steel weldment and optical micrographs obtained from weld
metal, base metal, interface, coarse grained heat-affected zone (CGHAZ), fine grained heat-affected zone (FGHAZ),

and intercritical zone (ICZ) of the weldment.


Ferritic Steels and Advanced Ferritic–Martensitic Steels

Weld metal

Base metal HAZ

II

HAZ

Base metal
IV

(a)

III
HAZ

Base
metal

I

Weld metal

HAZ


5 Cm

(b)

Figure 15 Locations77 of different types of failure in
weld geometry of the ferritic steels: (a) schematic
representation and (b) experimental observation in creep
tested weldment of 9Cr–1Mo steels.

Cavity
Cavity associated
with precipitate
10 mm

(a)

(b)

M23C6

[111] Z phase
Z-phase

1 mm
Figure 16 Type IV cracking in same sample as in
Figure 15. (a) cavities in the intercritical region and
(b) Z-phase77 in creep-tested 9Cr–1Mo steel. The inset
shows the microchemistry of the Z-phase.

The type IV cracking susceptibility, defined as the

reduction in creep-rupture strength of weld joint
compared to its base metal, depends on the type of
ferritic steel. 2.25Cr–1Mo steel is most susceptible to
type IV cracking; whereas the plain 9Cr–1Mo steel is
the least susceptible. At higher test temperature, the
type IV cracking susceptibility is higher in modified
9Cr–1Mo steel than the plain steel. This is related77
to the precipitation of Z-phase (Figure 16(b)), a complex Cr (V, Nb) N particle, in the modified steel. The
Z-phase grows rapidly at elevated temperatures during long term exposure, by dissolving the beneficial

117

MX types of precipitates. This promotes the recovery
of the substructure with associated decrease in
strength in the intercritical region of HAZ.
Although it is difficult to completely eliminate
Type IV cracking, several methods are being adopted
to improve type IV cracking resistance. It is more
severe in thick sections due to the imposed geometrical constraint. A design modification can be adopted to
decrease the variation in tensile stresses across the
welded section of the component or avoid joints in
critical regions having high system stresses and relocate them in the less critical region. Strength homogeneity across the weld joint can also be improved by
a suitable PWHT. An increase in width of the HAZ
can reduce the stress triaxiality such that the soft
intercritical region deforms with less constraint
with the consequence of reduced creep cavitation,
to minimize type IV cracking tendency. The width
of the HAZ can be increased both by changing preheat and heat-input during welding. Another contrasting approach to overcome type IV cracking is
to avoid or minimize the width of the HAZ, to eliminate the intercritical zone. This is being attempted by
employing advanced welding techniques such as laser

welding. The resistance against intercritical softening
can also be improved by increasing the base strength
of the steel with the addition of solid solution hardening elements such as W, Re, and Co and also by
microalloying the steel with boron. Microalloying
with boron retards the coarsening rate of M23C6 by
replacing some of its carbon. The boron content
needs to be optimized with the nitrogen content to
avoid BN formation. Addition of Cu is also found
to be beneficial. Copper is almost completely
insoluble in the iron matrix and when added in
small amounts, precipitates as nanosize particles to
impart creep resistance. A suitable adjustment of
the chemical composition of steel within the specification range also reduces the large difference in
creep strength between the softened HAZ, the base
metal, and the coarse grain HAZ of the joint. A weld
joint of modified 9Cr–1Mo steel with low carbon,
nitrogen, and niobium has been reported to possess
creep strength comparable to that of the base steel.
It is expected that a judicious combination of
changes in chemistry and process variables would
reduce the failures due to type IV cracking in weldments of ferritic steels subjected to creep loading.
Another frequent problem78–81 is the formation of
‘hard brittle zone’ during service exposure of dissimilar joints between ferritic steels, leading to failures.
The formation (Figure 17(a)) of microscopic layer of


118

Ferritic Steels and Advanced Ferritic–Martensitic Steels


hard, brittle zone along the HAZ in dissimilar weldments of steels is known to be responsible for the cold
cracking, stress corrosion cracking, and higher frequency of failures of the weldments. This is one of the
cases where modeling has enabled an in-depth
understanding of the problem, in addition to
providing an industrial solution to prevent the formation of brittle zone.
The brittle layer at the interface between
9Cr–1Mo weld and 2.25Cr–1Mo base metal is
shown77 to be a manifestation of a number of synergistic factors: (a) microstructural changes in regions
close to the heat source during welding (b) migration
of carbon during PWHT, driven by the gradient in
its activity and (c) formation (inset in Figure 17(a))

of series of fine carbides when there is a local
supersaturation of carbon. It has been possible to
use modeling methods like Finite Difference Methods to predict the carbon profile across the weld
region of 9Cr–1Mo and the base metal of 2.25Cr–
1Mo (Figure 17(b)), which were in good agreement
with the profiles obtained using electron probe
microanalysis. These calculations could be refined
using Thermo-Calc and diffusion-controlled transformations (DICTRA) to take into account the
simultaneous precipitation of carbides and diffusion
of carbon. These computational methods were instrumental in predicting the methods to prevent the
formation of hard zone in dissimilar joints of ferritic
steels. Three elements which would repel carbon
0.50
Temp: 1023 K

0.45
M1


B

C

A
M23C5

2.25Cr–1Mo

Soft zone

A

ppt zone

9Cr–1Mo

500 nm

Carbon concentration (wt%)

(b)

0.35
0.30

21/4Cr–1Mo

9Cr–1Mo


0.25
0.20
HZ

0.15

SZ

0.10
0.05

100 m

(a)

1h
15 h

0.40

0.00
-0.1

0.1

0.0
‘x’ (mm)

(b)


Cmaximum in hard zone (wt%)

0.5
Ni
Co
Cu

0.4

7
6.5
6

0.3

5.5

Z

5
4.5

0.2

4
6
5

5.0, 5.5, 5.5


0.1

(d)
0.00

(c)

0.02

0.04
0.06
d (mm)

0.08

Y

4

6
5.5

3 4.5

5

6.5

7


X

0.10

Figure 17 Modeling in preventing formation of hard zone in dissimilar joints of ferritic steels: (a) Optical micrograph78 of
formation of hard zone in the heat-affected zone between 9Cr–1Mo weld and the 2.25Cr–1Mo base metal, exposed to
1023 K for 15 h. The hard zone is marked as A and the soft zone as B. The inset shows the transmission electron micrograph of
individual carbides in the hard zone. (b) Finite difference method calculation78 predicting the diffusion profile of carbon in
the same weld geometry. (c) Variation79 of amount of carbon in the hard zone versus thickness of the ‘diffusion barrier’
introduced between the 9Cr–1Mo and the 2.25Cr–1Mo to prevent the formation of hard zone and (d) positions81 of carbon
atoms in a bcc iron lattice calculated using molecular dynamics.


Ferritic Steels and Advanced Ferritic–Martensitic Steels

atoms, that is, with the positive interaction energy
were chosen for this purpose. Figure 17(c) shows78
the comparison of three different metals, Ni, Cu,
and Co in preventing the formation of hard zone.
Experimental confirmation was obtained79 using interlayer between the two dissimilar ferritic steels. Further insight could also be arrived81 at in the diffusion
behavior (Figure 17(d)) of carbon interstitial in the
lattices of bcc iron and fcc nickel using molecular
dynamics. These calculations could demonstrate that
the activation energy for diffusion of carbon in a fcc
nickel lattice is higher than bcc iron. This sluggish
diffusion kinetics is due to the repulsive potential of
nickel toward carbon, which is the main reason for
the choice of nickel as the most effective diffusion
barrier between the two ferritic steels. Thus, an industrial solution to prevent the formation of brittle
zone in joints of dissimilar ferritic steels after service

exposure could be arrived at, based on an in-depth
understanding of the interaction between the lattice
potentials of atoms.
It has been demonstrated in the above studies that
modeling methods could be used most effectively to
reduce the experimental time required for overcoming an industrial problem. Experimental benchmarking was required only for final confirmation of
the predictions. These trends are becoming more
common in almost all problems in materials technology, in recent years, be it atomistic mechanisms or
fabrication technologies or prediction of life of components. It is hoped that this approach of knowledgebased design of materials would gradually replace the
time consuming empirical methods of today.

4.03.7 Summary
Future trends in the global fast reactor industry are
toward higher operating temperatures, higher burnup (250 GWd tÀ1), higher breeding ratios ($1.4) and
longer lifetime for reactor (60–100 years). These
goals require several developments in materials science and technology across all components of nuclear
plants, especially for core component materials.
Ferritic steels have a much better void swelling
resistance compared to currently used austenitic
stainless steels and are capable of enhancing the
burn-up of the fuel up to about $200 GWd tÀ1.
Ferritic–martensitic steels based on 9–12% Cr compositions exhibit the highest swelling resistance and
a number of commercial swelling resistant materials
have been marketed. The principles behind the

119

design of swelling resistant ferritic steels for core
components of fast reactors have been discussed.
However, their use is rendered difficult due to

their poorer creep strengths at temperatures higher
than $873 K. Improvement of higher temperature tensile and creep strengths in these alloys will enable us
to achieve higher temperatures, in addition to higher
burn-up, thus improving the economics of nuclear
power production. Presently, the reduced creep
strength of 9–12Cr ferritic steels at temperatures
above 798 K, has restricted their use to certain low
stressed components such as subassembly wrappers.
Another crucial problem is ‘embrittlement’ in ferritic
steels. The mechanisms and methods which are being
attempted to overcome embrittlement problems are
discussed.
Alloy development programmes are in progress to
explore ferritic–martensitic oxide dispersion strength
variants, for higher target burn-up of 250 dpa, with
enhanced high temperature ($973 K) capability, by
improving mechanical properties. Conventional alloy
melting routes will have to be abandoned in favor
of powder metallurgy techniques of ball-milling, hot
isostatic pressing, and hot extrusion for the synthesis
of these ODS steels. Process optimization for
the development of 9Cr-based ferritic–martensitic
steels strengthened by a fine dispersion of yttria
nanoparticles has been completed. The major concerns in this family of ferritic or ferritic–martensitic
steels are the anisotropy of properties in ferritic 12Cr
steels or oxidation resistance in 9Cr steels, fabrication
procedure, microstructural stability under irradiation, and dissolution during back-end technologies.
Materials science, engineering, and technology
have become an integral part of the aspiration of
the nuclear community to improve the economic

viability of fast reactors. One of the major concerns
in the alloy development programmes has been the
unacceptably long time taken to launch newer materials. It is expected that the current trends in
materials development, through intense international
collaborations and increased role of modeling in
materials behavior, would certainly reduce the time
and cost of alloy development programmes for future
reactors.

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