Comprehensive nuclear materials 5 09 material performance in lead and lead bismuth alloy

5.09

Material Performance in Lead and Lead-bismuth Alloy

K. Kikuchi
Ibaraki University, Ibaraki, Japan

ß 2012 Elsevier Ltd. All rights reserved.

5.09.1

Recent Lead-Alloy Activity

207

5.09.2
5.09.2.1
5.09.3
5.09.4
5.09.5
5.09.6
5.09.7
5.09.8
References

Utilization of LA
The Conceptual Models of ADS and MYRRHA
Ferritic–Martensitic Steels
Surface Treatment to F/M and Austenitic Steels
Oxide Dispersion-Strengthened Steel
Austenitic Stainless Steels

Precipitation Formation
Outlook

209
209
210
213
214
215
216
217
218

Abbreviations
ADS
AFM
BEM
DBTT
EB
EDX
F/M steel
GESA
GIF
ICP
LA
LBE
LFR
LINAC
MA
MEGAPIE

MFM
MYRRHA
ODS
OECD/NEA

SEM
WDX

Accelerator-driven nuclear
transmutation system
Atomic force microscopy
Backscattered electron microscope
Ductile-to-brittle transition
temperature
Electron beam
Energy-dispersed X-ray analyzer
Ferritic–martensitic steel
Gepulste Elektronenstrahlanlage
Generation IV International Forum
Inductive-coupled plasma atomic
emission spectrometer
Lead alloy
Lead–bismuth eutectics
Liquid-metal-cooled fast reactor
Linear accelerator
Minor actinides
MEGA-watt Pilot Experiment
Magnetic force microscopy
Multipurpose hybrid research
reactor for high-tech applications

Oxide dispersion-strengthened steel
The Organisation for Economic
Co-operation and Development/
The Nuclear Energy Agency
Scanning electron microscopy
Wave-dispersed X-ray analyzer

5.09.1 Recent Lead-Alloy Activity
A brief justification for the utilization of lead or lead
bismuth for use as a coolant in nuclear energy systems was given in 2001 by Sekimoto.1 When the
possibility of the utilization of nuclear energy was
discovered, it was expected to be a primary energy
source in the future. Fast reactors can utilize the
entire energy content of natural uranium. The selection of a coolant was an important item for designing
fast reactors. The neutron slowing-down caused by
the coolant should be minimized. This is first made
possible by decreasing the average atomic density of
the coolant in the reactor core, and second by
employing a nuclide with a large mass number as
the coolant, whose neutron moderating power is
low. A liquid metal is considered the best coolant
for using the second method. Initially, liquid mercury
was employed but it was not successful in either the
United States or Russia. Since then, several liquid
metals were considered, including lead alloys (LA),
and finally, sodium was selected. However, public
concern about the safety of sodium has increased
following sodium leakage incidents, so the development and deployment of fast reactors on more than a
prototype scale has not occurred.
In the last 10 years, the study of the utilization of

LA including lead–bismuth eutectics (LBE) has been
ongoing for application to nuclear waste transmutation
systems and lead–bismuth cooled nuclear reactors.

207


208

Material Performance in Lead and Lead-bismuth Alloy

LBE is a candidate material for a spallation target and a
reactor coolant. In the accelerator-driven nuclear
transmutation system (ADS), LBE is a candidate for
both the subcritical-reactor coolant and the spallation
neutron source target. In addition, the lead or lead–
bismuth-cooled fast reactor (LFR) is one of the four
reactor types investigated in Generation IV systems
proposed by the Generation IV International Forum
(GIF). A LBE-cooled Long-Life Safety Simple Small
Portable Proliferation-Resistant Reactor has also been
proposed.2
As a result of the investigations on LA, comprehensive literature has been published. The Working
Group on LBE of the OECD/NEA Nuclear Science
Committee3 published a handbook and review
reports on LA technology. The material properties
of lead and lead–bismuth are discussed in detail in
Chapter 2.14, Properties of Liquid Metal Coolants.
As part of the development of advanced nuclear
systems, including ADS proposed for high-level

radioactive waste transmutation and Generation IV
reactors, heavy liquid metals such as lead or LBE
were investigated as reactor core coolant and spallation targets. Heavy liquid metals were also being
envisaged as target materials for high-power neutron
spallation sources. The objective of the handbook is
to collate and publish properties and experimental
results on lead and LBE in a consistent format in
order to provide designers with a single source of
qualified properties and data and to guide subsequent
development efforts. The handbook covers liquid
lead and LBE properties, material compatibility and
testing issues, key aspects of the thermal-hydraulic
and system technologies, existing test facilities, and
open issues and perspectives.
Zhang and Li4 reviewed the studies on fundamental issues in LBE corrosion. They included phase
diagrams, thermodynamics, physical properties, corrosion mechanisms, oxygen control, experimental
results, and corrosion results. Some recommendations were proposed for future studies: precipitation
and deposition of corrosion products; oxygen transport; oxide formation and kinetics in LA; coolant
hydrodynamic effects; steel composition, microstructure, and surface effects; and corrosion models. These
are key areas for future research.
Fazio et al.5 characterized corrosion property for
ferritic–martensitic (F/M) steels and austenitic steels
in stagnant LA on the basis of the results of corrosion
tests. This report briefly summarized the current
status on LA activities. At a temperature below
450  C, adequate oxygen activities in the liquid

metal steels form an oxide layer that behaves as a
corrosion barrier. In the temperature range above
500  C, corrosion protection because of the oxide

scales seems to fail. A mixed corrosion mechanism
has been observed, where both oxide scale formation
and dissolution of the steel elements occurred. However, in this high-temperature range, it has been
demonstrated that the corrosion resistance of structural materials can be enhanced by coating the steel
with FeAl alloys. Experiments performed in flowing
LA (mostly LBE) confirm that the corrosion mechanism of the steels depends on the oxygen content in
LA. At relatively low oxygen concentration, the corrosion mechanism changes from oxidation to dissolution of the steel elements. The experimental activity
also extends up to temperatures of 750  C for oxide
dispersion-strengthened (ODS) alloys and their
welded variants in Pb. The use of materials at higher
temperatures will also require investigation of creep
rupture.
MEGAPIE was the MEGA-watt Pilot Experiment
done at Paul Scherrer Institut (PSI) in 2006 for
developing a LBE spallation target. The MEGAPIE
project was started as an essential step toward demonstrating the feasibility of coupling a high power accelerator, a spallation target, and a subcritical core
assembly. The project was expected to furnish important results regarding safe treatment of components
that had come into contact with lead–bismuth.6 The
design data was obtained and the operational mode
was confirmed.7 Corrosion rates were estimated
experimentally at 400  C for a LBE flow rate of
1 m sÀ1 and 2.2 m sÀ1 where the oxygen content in
the LBE was <10À7 wt%. No protective oxide layer
was produced on the steel surface. This oxygen content has been considered representative of the
MEGAPIE conditions, as no oxygen control and
monitoring system is anticipated to be used in the
target. The estimated corrosion rates, 40–86 mm
yearÀ1, indicate that in the given testing conditions,
the corrosion resistance of the steel does not represent a critical issue, especially since LBE temperature
is expected to be lower (320  C). The goals of the

experiment were fully accomplished8: 4 months of
reliable and essentially uninterrupted operation
(beam trips and short beam interruptions permitted)
at a power level as high as the accelerator was able to
deliver (about 0.75 MW) excellent performance of
the target and the dedicated ancillary systems, the
proof of functionality of advanced proton beam safety
devices, and, last but not least, a superb neutronic
efficiency delivering about 80% more neutrons for


Material Performance in Lead and Lead-bismuth Alloy

material usage in design studies. The material temperature at contact with LBE is slightly <500  C in the
spallation reaction area and <550  C in the fuel core
area under normal conditions.
Figure 1 shows the ADS concept. A superconducting linear accelerator (LINAC) is connected
with a subcritical fast reactor. A high-energy proton
beam is injected into the core of the reactor. Spallation reactions produce a number of neutrons from
the lead–bismuth nuclei, which are then used to
transmute minor actinides (MA). The interface
between the beam duct and lead bismuth is called
the beam window. For example, a tank type reactor
with 800 MW thermal power and LBE-coolant and
spallation target was proposed.11–13 The proton beam
energy was set at 1.5 GeV. The beam current varied
between 10 and 20 mA according to criticality swings.
In the steady-state condition, as the beam window
material generates heat by spallation reactions and is
cooled by flowing LBE. A temperature difference is

established between the LBE, the material in contact
with the LBE, and the material on the other side of
the window, with the temperatures being 400, 450,
and 500  C, respectively. As the MA core cladding
material is gamma heated and the fuel adds to the
radiation heat, temperatures reach, for example, a
maximum of 500, 550, and 600  C. The maximum
average velocity in the particular flow channel of
LBE is 1.8 and 2.0 m sÀ1, at the window and in the
MA core region, respectively.
Figure 2 shows the conceptual model of MYRRHA consisting of an inner vessel, guard vessel,

the users compared to the previously operated leadcannelloni target. Verification of performance will be
scheduled in the postirradiation experiment.

5.09.2 Utilization of LA
5.09.2.1 The Conceptual Models of ADS
and MYRRHA
Recent activity on materials research and development in LA, especially LBE, aims at realizing ADS,
MEGAPIE, LFR, and MYRRHA (multipurpose
hybrid research reactor for high-tech applications).9,10
It is valuable to know each specific environment for

Liq.He
RF
ADS

Injector
RFQ

DTL

Superconducting LINAC

Beam duct

P
Beam
window

Beam
window
MA
PbBi (Am,Cm)

n
MA

Subcritical reactor

PbBi

209

Spallation reaction

Figure 1 The conceptual model of accelerator-driven
nuclear transmutation system with beam window.

11

10

1. Inner vessel
2. Guard vessel
3. Cooling tubes

4

4. Cover
5. Diaphragm
6. Spallation loop
7. Subcritical core

5
8
12

9

6

11
8

7

9
10

7

11

3
13

2
1

12

8. Primary pumps
9. Primary heat exchangers
10. Emergency heat exchangers
11. In-vessel fuel transfer machine
12. In-vessel fuel storage
13. Coolant conditioning system

Figure 2 The conceptual model of subcritical reactor in multipurpose hybrid research reactor for high-tech applications.
Courtesy of J Bosch. ADS Candidate Materials Compatibility with Liquid Metal in a Neutron Irradiation Environment, Doctoral
Thesis, ISBN 978-90-8578-241-4, 2008; 7.


210

Material Performance in Lead and Lead-bismuth Alloy

cooling tubes, spallation loops, primary heat exchangers, and so on, but without a beam window.15 In this
system, a high-energy proton beam with an energy of

600 MeV is injected directly into the free surface
of the lead–bismuth in the subcritical reactor core.
The MYRRHA project aims to serve as a basis for
the European experimental ADS. In the first stage,
the project focuses mainly on demonstrating the ADS
concept, safety research of subcritical systems, and on
nuclear waste transmutation studies. Subsequently,
MYRRHA will be used as a fast spectrum irradiation
facility dedicated to research on structural materials,
nuclear fuel, liquid metal technology, and associated
aspects on the one hand and as a radioisotope production facility on the other. The system consists of a
proton accelerator that supplies a 600 MeV Â 3–4 mA
proton beam to a LBE spallation target, delivering the
primary neutrons, which in turn couples to a LBEcooled subcritical fast core. The structural materials for
MYRRHA need to withstand temperatures ranging
between 200 and 550  C (normal operating temperature between 300 and 450  C) under high spallation
neutron flux and contact with liquid LBE. It is clear
that the candidate materials need to fulfill challenging
requirements such as high thermal conductivity, high
heat resistance, low thermal expansion, low ductileto-brittle transition temperature (DBTT) shift, sufficient strength at elevated temperatures with limited
loss of ductility and toughness, low swelling rate, high
creep resistance, and good corrosion resistance.14
Studies of LA for developing ADS are also reported
from the points of view of conceptual ideas16,17 and
related facility.18

5.09.3 Ferritic–Martensitic Steels
One method of using materials such as F/M and
austenitic stainless steels in LA is to keep an oxide
layer on the surface of the base metal in contact with

LA by controlling the oxygen concentration in the
LA.19–21 Too little oxygen in LA will lead to dissolution of the protective iron oxide. Excess oxygen solution in the LA will lead to the production of lead
oxide that could plug the cooling tubes. Theory predicts that an adequate oxygen concentration in LA
exists between, for example, 10À6 and 10À4 wt% in
the temperature region of 400 and 700  C. An alternative method is to add anticorrosion elements such
as Al to the surface, which leads to a protective oxide
that guards base metals, as mentioned in the section
on surface treatment.

The oxide scale is not a simple structure but
consists of duplex layers: magnetite Fe3O4 near the
LA side and spinel (FeCr)3O4 near the base metal.
The original surface exits at the interface between
the magnetite and spinel but not at the front surface
of the magnetite near the LA. An early question
was how the oxide layers on the surface of the base
metal grew.
Martinelli et al.22–24 reported a global study on the
oxidation process of Fe–9Cr–1Mo martensitic steel
(T91) in static LBE. The isotope tracer oxygen-18 was
employed in the corrosion test. Also, the mass balance
of Fe and Cr was investigated theoretically. They
explained the Fe–Cr spinel growth rate mechanism as
follows: The oxidation reaction can occur because of
the presence of nano-channels. Nano-channel formation is achieved by the dissociative/perforative growth
in the magnetite. The nano-channel allows a fast diffusion of oxygen to the T91/spinel interface. Oxygen
cannot diffuse in the oxide lattice because its rate is
insufficient for Fe–Cr spinel formation, but is instead
transported via short cut diffusion paths. Even if oxygen
diffusion in grain boundaries could be possible, oxygen

would likely diffuse inside nano-channels. The nanochannels are, in some cases, called lead nano-channels
because of the results of the LBE oxidation tests. Liquid
metal does not penetrate evenly in the oxide scales;
only lead penetrations are observed. Nevertheless, in
pure bismuth oxidation tests, bismuth penetrations are
also observed in the scales. On the other hand, the iron
diffusion from T91 to the magnetite/Pb–Bi interface
leads to vacancy formation at the T91/Fe–Cr spinel
interface. Because of the presence of chromium atoms,
these vacancies can accumulate to form nano-cavities
at the T91/Fe–Cr spinel interface. This accumulation
is quasi complete; very few cavities are annihilated
on the T91/oxide interface. The Fe–Cr spinel grows
inside the nanocavity until it is completely filled.
At that moment, the oxygen can no longer reach the
T91 alloy and the oxidation reaction interrupts itself.
The formed Fe–Cr spinel thickness then becomes
equal to the consumed T91 thickness because of
this self-regulation process, as shown in Figure 3.
In this process, the limiting step of the Fe–Cr spinel
growth rate is thus the ‘iron diffusion’ across the oxide
scale.
A key issue in maintaining structural integrity is to
maintain high performance of the welded materials.
The corrosion properties between the base metal
and the weldment were investigated.25 The materials
tested were F/M steel F82H26 and the electron beam
(EB) welding of F82H. The chemical composition

Material Performance in Lead and Lead-bismuth Alloy

Oxygen

Nano-channel

211

Nano-channel

LBE

LBE
Oxide

Iron

Oxide

Original metal
surface
Nano cavity

Newly formed oxide

Figure 3 Self-regulation of the Fe–Cr spinel growth.

of F82H is 8Cr–2W–0.2VTa–bal/Fe (wt%). Oxygen
concentration was controlled to (2–4) Â 10À5 mass %.
Welded materials were prepared with a bead-onplate weldment with a 15 mm depth of melting.

F82H steel was welded after preheating at 300  C,
heat-treated at 300  C for 2 h, and then annealed at
750  C for 2 h for stress relief. Figure 4 shows optical
microscope observation of cross-section for F82H
specimens and an impinging-flow simulation around
the specimen. It was observed that the welded metal
of F82H revealed a coarse martensitic structure in
comparison with the fine microstructure in the nonwelded region because of melting and resolidification
in the welding process. The corrosion depth in F82H
was limited near the surface of the material. A failure
of the outside layer in the duplex corrosion layers was
observed. The heat-affected zone showed that the
martensitic structure became fine because of the
rapid heating and cooling during the welding process.
Regardless of the difference in microstructures, the
corrosion layer showed no apparent difference. The
growth of the corrosion depth, defined by the layers
of magnetite and spinel, followed a parabolic law,
where diffusion controls the process. The result of
the flow simulation of LBE impingement indicated
that the velocity varied from 0 to 1 m sÀ1 near the
specimen surface. At higher temperatures, for example, above 500  C, the internal oxide layer or diffusion zone was clearly identified. Furukawa et al.27
observed three layers, consisting of the duplex layers
mentioned earlier and a diffusion zone in the base
metal beneath the spinel layer in the static LBE
test at 500 and 550  C under the oxygen control
to 10À6 wt% for high Cr steel (10.54Cr–1.75W–
MnMoV) with heat treatment: 1070  C, 100 min aircooled; 770  C, 440 min air-cooled.
Tan and Allen tested high Cr steel material in the
DELTA loop, at Los Alamos National Laboratory

(LANL). The material tested was HCM12A, procured from Sumitomo Metal Industries, Ltd., with

composition provided by the supplier: 10.83Cr–
1.89W–1.02Cu–0.64Mn–0.39Ni–0.30Mo–0.27Si–0.19V–
0.11C–0.063N–0.054Nb–0.016P–0.002S–0.001Al–3.1 Â
10À5 B, and balance/Fe (wt%).28 The chemical composition and heat treatment of this material are slightly
different from those used in the experiment by
Furukawa and Muller. HCM12A is one of the thirdgeneration 12Cr ferritic steels with tempered martensite,29 which was originally developed for heavy section
components such as headers and steam pipes for use
at temperatures up to 620  C and pressures up to
34 MPa30 with good resistance to thermal shocks.31
The HCM12A was received after being annealed at
1050  C and tempered at 770  C.28 They compared the
oxide layer to the porous magnetite layer on the supercritical water exposed sample at 600  C, 667 h. Temperatures at both conditions were different. It was
found that detachment of most of the magnetite nonprotective layer occurred on the LBE-exposed sample
at 530  C–600 h earlier in time than models developed
by Zhang and Li. From a technical experimental point
of view, it is the issue how to detect the original surface
of base metal in order to evaluate the oxide thickness.
A thin yttrium coating layer will help to detect it in the
LBE corrosion test.
At temperatures above 600  C, the oxide layer grew
thinner with increasing temperature, which suggests
that around this temperature, a change occurred in
the mechanism of oxidation. At 570  C, FeO-wustite
is formed. Compared with magnetite, wustite has a
lower standard free energy of formation, which ensures
its stable existence at low oxygen potential. In fact, the
layer was formed in the region between magnetite
and base metal. Also in this temperature range that is

beyond the point of oxidation mechanism change, dissolution attack was observed at several points, and the
number of such points increased with prolongation of
run duration. The observations would suggest lowering
the maximum processing temperature in LBE applications from the point of view of the static LBE test.27


212

Material Performance in Lead and Lead-bismuth Alloy

Welded zone
Heat-affected
zone

Tip

1 mm

(a)

1 mm

(b)

Oxide layer

20 µm

Oxide layer
20 µm

(d)

(c)

Oxide layer
Oxide layer

Pb–Bi
(e)

(f)
STAR

Specimen
PROSTAR 3.10
Velocity magnitude
M/S
Local MX = 1.089
Local MN = 0.7602E-03
*Presentation grid*
1.089
1.012
0.9338
0.8560
0.7783
0.7005
0.6228
0.5450
0.4673

0.3895
0.3118
0.2340
0.1563
0.7851E−01
0.7603E−03

LBE flow
Y

x
z

(g)
Figure 4 Optical microscope observation of cross-section for F82H specimens and an impinging-flow simulation around
the specimen. (a) Macro structure, including welded zone and heat-affected zone, (b) macro structure at the specimen
end where lead–bismuth eutectics impinges from the right hand side indicated with an arrow, (c) micro structure of welded
zone tested at 450  C for 1000 h, (d) micro structure of tip region tested at 450  C for 1000 h, (e) cross-section of tip
region tested at 450  C for 3000 h, (f) cross-section of tip region tested at 500  C for 1000 h, and (g) simulated flow profile
of lead–bismuth eutectics around the specimen.


Material Performance in Lead and Lead-bismuth Alloy

Hosemann et al. attempted nano-scale characterization of HT-9 (11.95Cr–1Mo–0.6Mn–0.57Ni–
0.5W–0.4Si–0.33V–bal/Fe (wt%)) by using atomic
force microscopy (AFM), using a function of magnetic force microscopy (MFM) and C-AFM. C-AFM
is a contact mode electrical characterization technique that involves applying a voltage typically
between the conductive AFM tip and the sample
while monitoring variations in the local electrical

properties in a range of picoamperes to microamperes. The HT-9 tube was tested at 550  C in flowing
LBE under 10À6 wt% oxygen for 3000 h.32 It was
found that the oxide consists of at least four different
layers with different grain structures and therefore
conductivity/magnetic properties. The outer layers
seem to be Fe3O4 and have good conductivity, while
the inner layer is Cr enriched and has lower conductivity or is insulating. This is in agreement with the
literature where Cr additions lower the conductivity
of Fe3O4. The outer layer can be divided into two
distinct areas based on a change in grain structure.
The inner oxide layers adopt the grain structure from
the bulk steel. High pore density within these layers
suggests that these are fast diffusion paths allowing Fe
diffusion outward and O diffusion inward. The LBE
corrosion experiment in the DELTA Loop on T91,
HT-9, and EP823 conducted for 600 h at 535  C
showed multilayer oxides on the tested materials.
The wave-dispersed X-ray analyzer (WDX) measurements on the cross-sections revealed two Cr and Fe
containing oxide layers and no Fe3O4 layer. It appears
that the main difference between observed oxide
layers is the Fe content and the microstructure.
Nano-indentation tests across the oxide layers
were performed.33 The results showed lower values
of E-modulus in these oxide layers than that of the
bulk steel layers and higher hardness values for the
oxides than that of the bulk steel. The inner oxide
layer is softer than the outer oxide layer. This might
be due to the fact that the inner oxide layer has higher
porosity than the outer layer.
Yamaki and Kikuchi34 conducted a mechanical

test of oxide scales. The beam window at the boundary of the high-energy proton beam and reactor core,
as shown in Figure 1, is loaded by thermal stress
and buckling load in the deep LBE of the reactor.35
The specimen was a ring made from the F/M steel
pipe, HCM12A. The inner surface of the pipe had
been exposed to flowing LBE during the loop operation at 400–500  C for 5500 h under an oxygen concentration in the range from 1 Â 10À5 to 5 Â 10À5 wt%.
Apparently, the oxide layer had a duplex structure.
Possibly they were outside the magnetite and inner

213

side spinel. Figure 5 shows the results of the ring
compression test. The HCM12A ring was compressed by 50% and unloaded. Near position A,
cracking occurred because of excess strain to the
spinel layer rather than the Fe3O4 layer. This was
caused by the fact that the Young’s modulus of
Fe–Cr spinel layer was lower than that of Fe3O4
layer by 10% with the same hardness in both layers.
Near position E, the oxide layer was spalled off
from the boundary between the base metal and the
spinel.
F/M steels can guard the base metal by forming a
spinel oxide layer. The formation mechanism is controlled by the iron diffusion rating. The magnetite
oxide layer is not protected against contact with
LBE. Under the tensile stresses, excess strains will
spall off the oxide layer from the interface between
the base metal and the spinel layer. Over 570  C, the
oxide formation mechanism is changed by the formation of wustite.

5.09.4 Surface Treatment to F/M

and Austenitic Steels
Muller et al. demonstrated that the effect of Al-alloying
into the surface of the base metal was the F/M
steel OPTIFER IVc (10Cr–0.58Mn–0.56C–0.40W–
0.28V–bal/Fe (wt%)) using electron pulse treatment,
GESA (Gepulste Elektronenstrahlanlage).36,37 There
was no corrosion attack visible in any part of the
alloyed portion after 1500 h exposure to liquid lead
at 550  C with 8 Â 10À6 at.% oxygen. The alumina
layer that must have formed at the surface during
oxidation in lead might be very thin and could not be
detected. Only the unalloyed part of the surface was
covered with thick oxide scales. The results also
suggested that the Fe–Cr spinel layer ends at the
original specimen surface. This surface treatment
had a similar result when applied to an austenitic
steel base metal 1.4970(16.5Cr–13.8Ni–1.91Mn–
0.81Si–MoTi–bal/Fe (wt%)).
Weisenburger et al. examined T91 tubes with
modified FeCrAlY coatings in LBE. These coatings
are often used for turbine blade protection.38 The
coating had an average thickness of 30 mm after application by a plasma spray method and was remelted
using the pulsed large area GESA EB to gain a dense
coating layer and to improve the bonding between
the coating and the bulk material. They intended to
simulate a cladding material’s environment by using a
pressurized tube type specimen. The results showed


214

Material Performance in Lead and Lead-bismuth Alloy

Load
Load

Delamination

Initial
diameter
(48 mm)

Final
diameter
(24 mm)

Specimen
Flat plate

(a)

d width

Crack
Base metal

d interval

(b)
10 mm

A

B C

I

F
H G

D

27.5 mm

E

(e)

Oxide scales

10 mm

(c)
Base metal

A

No oxide scale

B
C

100 mm

D

(f)
E

(d)
Figure 5 The ring compression test. (a) HCM12A ring model before compression, (b) the ring model after compression
by 50%, (c) the ring after unloading, (d) simulation of maximum principal strain distribution induced at loading, (e) the
cross-section near position A, and (f) the cross-section near position E.

that the tangential wall stress of about 112.5 MPa
induced by an internal tube pressure of 15 MPa
increased the Fe diffusion and led to enhanced magnetite scale growth. Coated specimens, however, have
no magnetite layer. This is another advantage of the
coating. Energy-dispersed X-ray analyzer (EDX) line
scans of the cross-section of the coated T91 tube
specimen after 2000 h exposure to LBE at 600  C
show that the top oxide scale must consist mainly of
alumina followed by a thin layer enriched in Cr.
These layers protect the steel not only from LBE
attack but also from oxygen diffusion into the coating
and bulk material. The coating process needs some
improvement to avoid coating regions with aluminum concentrations below 4 wt%; otherwise, the
oxide layer will grow in the same manner as the
original material. Steels with 8–15 wt% Al alloyed
into the surface suffer no corrosion attack for all
experimental temperatures and exposure times.39
Technical concerns about surface treatment are the

effect of cyclic loading on the low cycle fatigue
endurance in air and LBE. Low cycle fatigue tests
were conducted in LBE containing 10À6 wt% dissolved oxygen with T91 steel at 550  C. T91 was

employed in two modifications, one in the as-received
state and the other after alloying FeCrAlY into the
surface by pulsed EB treatment (GESA process).
Tests were carried out with symmetrical cycling
(R ¼ 1) with a frequency of 0.5 Hz and a total elongation Det/2 between 0.3% and 2%. No fatigue effects
from LBE could be detected. Results in air and LBE
showed similar behavior. Additionally, no difference
was observed between surface treated and nontreated
T91 specimens.14
A melting process of coating materials enhances
bonding between the coating and the bulk materials.
Heat deposition because of a pulsed EB exposure
successfully demonstrated that remelted alumina or
FeCrAlY coating was effective in protecting the base
metal property from LBE attack.

5.09.5 Oxide DispersionStrengthened Steel
Takaya et al.40 investigated the corrosion resistance of
ODS steels with 0–3.5 wt% Al and 13.7–17.3 wt% Cr,
at 550 and 650  C for up to 3000 h in stagnant LBE


Material Performance in Lead and Lead-bismuth Alloy

containing 10À6 and 10À8 wt% oxygen. The ODS
steels were manufactured by hot extrusion of mechanically alloyed powders at 1150  C, and consolidated

bars were annealed by 60 min of heat treatment at
1150  C, followed by air cooling. Chemical compositions of ODS materials are (13.7–17.3) Cr–(1.9–3.5)
Al–(0.34–0.36)Y2O3–TiSiMn–bal/Fe (wt%). Protective Al oxide scales formed on the surfaces of the
ODS steels with 3.5 wt% Al and 14–17 wt% Cr, and
no dissolution attack was seen in any of the cases.
Addition of Al is very effective in improving the
corrosion resistance of ODS steels in LBE. On the
other hand, the ODS steel with 16 wt% Cr and no Al
showed no corrosion resistance, except in the case of
exposure to LBE with 10À6 wt% oxygen at 650  C.
Thus, the corrosion resistance of ODS steels in LBE
may not be improved solely by increasing Cr
concentration.
There is additional data reported by Hosemann
et al. on ODS alloys in LBE. Specimens were exposed
to flowing LBE in the DELTA Loop at LANL
at 535  C for 200 and 600 h. The oxygen content
in the LBE was about 10À6 wt%. The detailed
manufacturing process was not disclosed. Conclusively,
PM2000, which has a chemical composition of 20Cr–
5.5Al–0.5Y2O3–0.5Ti–bal/Fe (wt%), showed a very
dense, thin, and protective oxide layer because of its
higher Al content. The compositions of the oxide layers
found on the Al alloyed materials change with depth.
Elements are oxidized based on the amount of oxygen
available for oxidation and the free energy of the oxide.
It appears that at least 5.5 wt% Al in the alloy is
necessary to form a protective Al-enriched oxide.41
The oxide scale has a duplex structure below 500  C.
Over 500  C, a diffusion zone in the base metal is

apparently observed. The oxide layer appears to consist
of three layers, that is, the duplex layers plus the diffusion zones. The oxide scale does become unstable. The
outer magnetite layer is prone to be spalled off in the
flowing LA. In such an environment, the Al coating is
found to be effective in enhancing corrosion resistance.
Remelting processes, for example, by GESA EB exposure, make a good Al layer. The disadvantage of the
coating method is the disintegration or cracking due to
an uncontrolled process. This cracking could be attributed to the local reduction of Al content. The ODS
alloy is developed for cladding materials. Materials
development is progressing in the direction of highCr Al-ODS alloys. The recommended Al composition
in ODS alloys varied from 3.5 to 5.5. An adequate
amount of Al will balance the corrosion resistance and
mechanical strength. Excess Al will reduce mechanical

215

strength. The reason why the Al enrichment in Febase steel improves corrosion resistance in LA will be
determined in future investigations.
ODS steel aims at enhancing the strength of
material applicable to the cladding materials of a
fast reactor. The addition of Al to ODS improves
corrosion resistance in LBE at the fuel cladding
temperatures. On the other hand, the excess addition of aluminum reduces the strength of materials.
An optimization is needed to balance the two factors at around 5%.

5.09.6 Austenitic Stainless Steels
Austenitic stainless steels are candidate materials for
the spallation target window in ADS. In MEGAPIE,
however, F/M steel, T91, was used for the beam
window in flowing LA, and this was acceptable for a

limited duration (4 months). The lifetime of the beam
window of the T91 liquid Pb–Bi container in the
MEGAPIE target was summarized based on the present knowledge of LBE corrosion, embrittlement, and
radiation effects in the relevant condition.42 It was
suggested that the lower bound of the lifetime of the
T91 beam window was determined when the steel
became brittle at the lowest operation temperature,
230  C, with a safety margin of 30%. Evaluation
using the DBTT data and fracture toughness values
of T91 specimens tested in LBE, a dose limit of about
6 dpa, corresponding to 2.4Ah proton charge to be
received by the target in about 20 weeks in the normal operation condition, was set.
In the ADS design, for example, the beam window
material will produce about 1000 appm (3He þ 4He)
a year by 1.5 GeV proton beam bombardment in the
reactor core for austenitic stainless steel, Japanese
Primary Candidate Alloy (JPCA), and F/M steel,
F82H.43 The helium production of 1000 appm He
suggests that the DBTT will increase by 400–
500  C.44 This increase will set the design temperature at the beam window at 450–500  C. The use of
a F/M steel may lead to a brittle fracture, and those
materials should be avoided in operations for
extended times. Therefore, austenitic steel is the
candidate material. The production of hydrogen and
helium in JPCA was slightly larger, 3–4%, than
that of F82H because of the addition of nickel and
boron. JPCA, in which the chemical composition
is 0.50Si–1.77Mn–0.027P–0.005S–15.60Ni–14.22
Cr–2.28Mo–0.24Ti–0.0031B–0.0039N–bal/Fe (wt%),
was developed to reduce the helium embrittlement

216

Material Performance in Lead and Lead-bismuth Alloy

of austenitic steel for first wall and blanket structural components in fusion reactors.45 The optimized
JPCA material is manufactured by vacuum induction
melting, vacuum arc melting, and solution-annealing
at 1100  C for 1 h. The TiC precipitates within the
matrix and on the grain boundaries serve as trapping
centers for the helium produced during neutron irradiation. However, dissolution of the MC precipitates
initiates the onset of helium embrittlement as well as
high swelling during high fluence neutron irradiation. The improved stability of the MC precipitates,
which formed in the matrix during irradiation, prevents loss of ductility at 500  C and below.
The corrosion properties of an austenitic stainless
steel at low temperature demonstrated good endurance for material usage in LBE during a short time,
approximately at 300 and 470  C for 3000 h for 1.4970
austenitic stainless steel,46 and at 420  C for 2000 h
for 1.4970 austenitic stainless steel and 316L at an
oxygen concentration of 10À6 wt%.47 No dissolution
was seen in the aforementioned results. A thin oxide
scale may protect the material from attack in LBE.
As demonstrated in Figure 4, a corrosion test
under impinging flow was also conducted on JPCA
and its EB welded bar at an oxygen concentration of
2–4 Â 10À5 wt%.48 The EB welded metal of JPCA
exhibited a dendritic structure 1 mm in width, but a
heat-affected zone was not visible. Scanning electron
microscopy (SEM) observation showed no corrosion

layer for the specimens tested at 450  C and 1000 h.

1600
1400
1200
1000
800
600
400
200
1200
1000
800
600
400
200
2000

But at 3000 h, a thin corrosion layer could be
observed at 1–2 mm in depth. For the weld joint, the
depth of the corrosion layer as well as corrosion
morphology showed the same results as with the
parent material. The results of X-ray diffraction analyses showed how the oxide layer developed at 450
and 500  C. Figure 6 shows X-ray diffraction analyses of the JPCA specimens under the conditions of
1000 h at 450  C (top, JPCA-1), 1000 h at 500  C
(middle, JPCA-2), and 3000 h at 450  C (bottom,
JPCA-3). Oxidation of the JPCA at 450  C progressed in the same manner as at 500  C.

5.09.7 Precipitation Formation
The dissolution of Ni, Cr, and Fe from structural

materials into LA was studied. Saturated solubility
of Ni in LA is calculated to be a couple of wt% at
450–500  C.20 Corrosion–erosion tests have been conducted at the JLBL-1 facility of the Japan Atomic
Energy Agency (JAEA). The main circulating loop
was made of SS316 austenitic stainless steel, and consisted of the specimens at high and low temperatures,
filters, a surge tank, a cooler, an electromagnetic flow
meter, a surface-level meter, thermocouples, and a
drain tank.49 The loop was operated at a maximum
temperature of 450  C with a temperature difference
of 50–100  C and average flow velocity of 1 m sÀ1. The
oxygen concentration was estimated to be 10À7 wt%,
Cr0.19, Fe0.70, Ni0.11
Bi
M3O4

JPCA-1

JPCA-2

JPCA-3

1500
1000
500
0
10.0

20.0

30.0

40.0

50.0

60.0

70.0

80.0

90.0

100.0

Figure 6 X-ray diffraction analyses of JPCA specimens under the condition of 1000 h at 450  C (top, JPCA-1), 1000 h
at 500  C (middle, JPCA-2), and 3000 h at 450  C (bottom, JPCA-3).


Material Performance in Lead and Lead-bismuth Alloy

according to measurements using an oxygen probe.
The testing specimen tube is a cold-drawn seamless
type SS316, which was produced as a tubing form with
13.8 mm outer diameter, 2 mm thickness, and 40 cm
length. The tube was solution-heat treated at 1080  C
for 1.5 min and then cooled rapidly. Figure 7 shows
both the EDX analyses of the low-temperature specimen after corrosion–erosion testing for the 3000 h
and the SEM image of an unused specimen. The surface of unused specimen was characterized by the creviced structure. This feature resulted from the acid
washing in the manufacturing process during material

preparation. It was found that precipitated materials
existed with the solidified LBE on the tube. The precipitation consists mainly of Fe and Cr as measured
using energy dispersive X-ray analyses and apparently
looks crystalline. The quantitative analyses by means of
a focused-ion beam, X-rays, and Transmission Electron
Microscope (TEM) showed that the weight concentration ratio was roughly Fe:Cr = 9:1, for example. Nickel
was not found in the crystals or in the solidified LBE.
The precipitations occur in the lead–bismuth including
impurities dissolved from SS316 at the high temperature portion of the test.
Zhang and Li4,50 calculated the corrosion/precipitation rate for iron using a kinetic corrosion model
and the temperature profile for the JLBL-1 loop.
They reported that the observed deposition zones in
the JLBL-1 loop could be exactly predicted using the
nonisothermal and multimodular corrosion model.
The predicted corrosion rate is about 0.05–0.08 mm
per 3000 h if the diffusion coefficient is selected as

SEM

BEM

Fe

Cr

217

3.9 Â 10À9 m2 sÀ1. This agrees well with the experimental results of 0.03–0.1 mm.
Ni-rich precipitation was found in the JLBL-1
loop after a total operation time of 9000 h was

achieved. Figure 8 shows Ni-rich precipitates in an
SEM (low magnitude) and laser microscope (high
magnitude) images on the surface of solidified
LBE.51 The solubility of Ni is higher than that of
Fe and Cr, around a couple of wt% in the temperature range of 350–450  C. For the measurement of
Ni in LBE, an inductive-coupled plasma atomic
emission spectrometer (ICP, ULTIMA2) was used
for analyses. It was found that the Ni concentration
was below 0.1 wt%. In addition, Ni-rich precipitates
were found not only at the high temperature part but
also at the low temperature part, and on the surface of
residual LBE as well. This was not the case for Fe–Cr
precipitates; they were only found at the low temperature part. The driving force for Fe–Cr precipitates
was concluded to be a difference of the saturation
concentration at different temperatures. It can be
assumed that the Ni-rich precipitates formed on the
surface of the residual LBE during a cooling period,
although the precipitation rate is not known for the
establishment of such a Ni-rich structure.

5.09.8 Outlook
For the use of LA as coolant and spallation target,
it is important that the compilations and databases
of material properties are extended to include

Unused specimen
surface
50 µm

20 µm

500 µm
Ni

Pb

Bi

Figure 7 Energy-dispersed X-ray analyzer analyses of
low-temperature specimen after corrosion–erosion test at
JLBL-1 and scanning electron microscopy image of unused
specimen.

Figure 8 Ni-rich precipitates of scanning electron
microscopy (low magnitude) and laser microscope (high
magnitude) on the surface of solidified lead–bismuth eutectics.
High magnification image is taken by laser microscope.
Reproduced from Kikuchi, K.; Saito, S.; Hamaguchi, D.;
Tezuka, M. J. Nucl. Mater. 2010, 398, 104–108.


218

Material Performance in Lead and Lead-bismuth Alloy

the mechanical properties of structural materials
in LA, such as the effect of temperature and strain
rate,52 fracture toughness,25,53 weldment,54 and liquid
metal embrittlement.55 They are essential to the
design work for the concepts of the ADS systems.
Erosion–corrosion of materials in flowing LA

should be considered along with details of the flow
profile. It was recognized that magnetite in the oxide
layer is eliminated in the steady-state flowing LA.
There is evidence of erosion detected in the expanded
flow channel of the JLBL-1 test specimen. The eroded
part of the specimen coincides with regions in flowing
LA, such as secondary flow or vortex flow.
In the last decade, research projects for ADS, LFR,
MEGAPIE, and MYRRHA were launched. To design
the conceptual model of the real spallation target, the
performance of neutronics, reactor physics, and thermal hydraulics has been studied along with the performances of materials in lead–bismuth or lead. For
the materials bombarded by high-energy particles,
such as proton beams, the surface coating or surface
treatment, defect formation mechanism including not
only corrosion but also the synergetic effect of irradiation field, must be understood. For the materials
applicable to the cladding of fuel rod, Fe–Al alloys
seem to be effective in the use of LA, but optimum Al
concentration must be determined.

11.
12.
13.

14.

15.

16.

17.

18.
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