4.04

Radiation Effects in Nickel-Based Alloys

R. M. Boothby
National Nuclear Laboratory, Harwell, Oxfordshire, UK

ß 2012 Elsevier Ltd. All rights reserved.

4.04.1

Introduction

123

4.04.2
4.04.2.1
4.04.2.2
4.04.2.3
4.04.3
4.04.4
4.04.4.1
4.04.4.2
4.04.5
4.04.5.1
4.04.5.2
4.04.6
References

Void Swelling
Compositional Dependence of Void Swelling

Void-Swelling Models
Swelling Behavior of Neutron-Irradiated Nimonic PE16
Irradiation Creep
Microstructural Stability
Dislocation Structures
Precipitate Stability
Irradiation Embrittlement
Fast Neutron Irradiation Experiments
Helium Implantation Experiments
Concluding Remarks

124
124
129
133
136
138
138
139
140
140
145
147
148

Abbreviations
AGR
DFR
EBR-II
EDX

HVEM
N/2
NRT
OA
PFR
PS
SIPA
ST
STA
TEM
UTS
VEC

Advanced gas-cooled reactor
Dounreay Fast Reactor
Experimental Breeder Reactor-II
Energy dispersive X-ray
High-voltage electron microscope
dpa calculated according to half-Nelson
model
dpa calculated according to Norget,
Robinson, and Torrens model
Overaged
Prototype Fast Reactor
Proof stress
Stress-induced preferred absorption
Solution treated
Solution treated and aged
Transmission electron microscope
Ultimate tensile strength

Variable energy cyclotron

4.04.1 Introduction
Research into the effects of irradiation on nickelbased alloys peaked during the fast reactor development programs carried out in the 1970s and 1980s.
Interest in these materials focused on their high resistance to radiation-induced void swelling compared to
austenitic steels, though a perceived susceptibility to

irradiation embrittlement limited their application to
some extent. Nevertheless, the Nimonic alloy PE16
was successfully used for fuel element cladding and
subassembly wrappers in the United Kingdom, and
Inconel 706 was utilized for cladding in France. Both
of these materials are precipitation hardened and
consequently have high creep strength, and much
research and development of alternative alloys was
directed toward maintaining swelling resistance and
creep strength while aiming to alleviate, or at least
understand, irradiation embrittlement effects. There
has been some revival of interest in nickel-based
alloys for nuclear applications, and various aspects of
radiation damage in such materials have recently been
reviewed by Rowcliffe et al.1 in the context of Generation IV reactors, and by Angeliu et al.2 in consideration of their use for the pressure vessel of the
Prometheus space reactor. Nickel-based alloys are
also candidate structural materials for molten salt
reactors, for which resistance to corrosion by molten
fluoride salts and high-temperature creep strength
are prime requirements, though intergranular attack
by the fission product tellurium and irradiation
embrittlement due to helium production are potentially limiting factors for this application.3
This chapter focuses on the void swelling behavior, irradiation creep, microstructural stability, and

irradiation embrittlement of precipitation-hardened
nickel-based alloys. Fundamental to all of these
effects are the basic processes of damage production
123


124

Radiation Effects in Nickel-Based Alloys

by the creation of vacancies and interstitial atoms
in displacement cascades, and the ways in which
these point defects migrate and interact with, causing
the redistribution of, solute atoms. Detailed discussions of damage processes and radiation-induced
segregation are beyond the scope of this chapter
but these topics will be introduced where necessary,
particularly in relation to void swelling models.
More detailed reviews are given in Chapter 1.01,
Fundamental Properties of Defects in Metals;
Chapter 1.03, Radiation-Induced Effects on
Microstructure; Chapter 1.11, Primary Radiation
Damage Formation; Chapter 1.12, Atomic-Level
Level Dislocation Dynamics in Irradiated Metals
and Chapter 1.18, Radiation-Induced Segregation.
Typical compositions of nickel-based alloys and
some precipitation-hardened steels, which are considered in this chapter, are shown in Table 1. Alloy
compositions are generally given in weight percent
throughout this chapter unless stated otherwise.
Precipitation-hardened alloys may be utilized in a
number of different heat-treated conditions, which

are generally abbreviated here as ST (solution treated), STA (solution treated and aged), and OA (overaged). Further information on the material properties
of nickel alloys is given in Chapter 2.08, Nickel
Alloys: Properties and Characteristics.
Neutron fluences are generally given for
E > 0.1 MeV unless indicated otherwise. Atomic displacement doses (dpa) are generally given in NRT
(Fe) units, although the half-Nelson (N/2) model was
Table 1

widely used particularly in the United Kingdom in
the 1970s6. The exact relationship between these
units will vary depending on the neutron spectrum
(which may differ, not only from one reactor to
another, but also depending on location within a
reactor), but approximate conversion factors for fast
reactor core irradiations are
1026 n mÀ2 ðE > 0:1MeVÞ
¼ 5dpa NRTðFeÞ ¼ 6:25dpa ðN=2Þ

4.04.2 Void Swelling
4.04.2.1 Compositional Dependence of
Void Swelling
Nimonic PE16 was first identified as a low-swelling
alloy in the early 1970s. Void swelling data derived
from density measurements on fuel pin cladding materials from the Dounreay Fast Reactor (DFR) were
reported by Bramman et al.7 and were complemented
by electron microscope examinations described by
Cawthorne et al.8 Swelling in STA PE16 was found to
be lower than in heat-treated austenitic steels and
comparable to cold-worked steels. Comparison of
data for PE16 and FV548 (a Nb-stabilized austenitic

steel) irradiated under identical conditions in DFR to
a peak neutron fluence of $6 Â 1026 n mÀ2 indicated
that the lower swelling of PE16 was due to smaller
void concentrations at irradiation temperatures up
to $550  C and reduced void sizes at higher

Nominal compositions (wt%) of commercial and developmental nickel-based alloys

Alloy

Ni

Cr

Mo

Ti

Al

Nb

Mn

Si

C

Nimonic PE16
Inconel 706

Inconel 718
Inconel 600
Inconel 625
Incoloy 800
Hastelloy X
D21a
D25a
D66a
D66b
D68a
D68b
PE16 matrix
Incoloy DS
Alloy 7817
Alloy 7818

43
41.5
52.5
75
61
34
48
25
30
45
40
45
34
36

39
40
40

16.5
16
19
16
22
20.5
21
8.4
10.5
12
11
12
12.5
20
18
15
15

1.1
–
3.0
–
9.0
–
9.0
1.0

3.7
3.0
2.0
–
–
4.0
–
3.2
3.0

1.2
1.8
0.9
0.3
0.3
0.4
–
3.3
1.8
2.5
3.0
1.8
1.6
–
0.04
2.0
0.3

1.2
0.2

0.5
0.2
0.3
0.4
–
1.7
1.3
2.5
1.5
0.4
0.25
–
0.02
0.9
–

–
2.9
5.2
–
3.5
–
–
–
–
–
–
3.6
2.8
–

–
–
3.0

0.1
0.2
0.2
0.2
0.2
0.9
0.5
1.0
1.0
–
0.2
0.3
0.2
0.1
1.0
0.2
0.2

0.2
0.2
0.2
0.2
0.2
0.5
0.5
1.0

1.0
0.5
0.5
0.4
0.4
0.2
2.0
0.5
0.5

0.05
0.03
0.04
0.08
0.05
0.07
0.10
0.04
0.04
0.03
0.04
0.03
0.02
0.07
0.08
0.02
0.02

a


Composition indicated by Yang et al.4
Composition indicated by Toloczko et al.5

b

Other

0.3Cu
0.5Cu
0.5W, 2.0Co

Fe
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.
Bal.



Radiation Effects in Nickel-Based Alloys

energy Ed ¼ 40 eV). In addition to precipitationhardened alloys, including PE16 and Inconel 706,
this experiment included nonhardenable high-Ni
alloys, such as Inconel 600 and Hastelloy X, a range
of commercial steels, and Fe–Cr–Ni ternary alloys
containing 15% Cr and 15–35% Ni. The alloys were
preimplanted with 15 appm helium prior to ion bombardment, and the irradiation temperature was chosen
as being close to the peak swelling temperature for ionirradiated austenitic steels. The extent of void swelling
was determined by electron microscope examinations
in low-swelling alloys, but was estimated from stepheight measurements (comparing the surfaces of irradiated and nonirradiated regions) in high-swelling
materials. As illustrated in Figure 1, the results showed
negligible swelling (<0.1%) in PE16, Inconel 706,
Hastelloy X, and the Fe–15Cr–35Ni ternary alloy,
low swelling (<1%) in other high-Ni alloys, but high
swelling (generally >20%) in austenitic steels. In commercial alloys containing $18% Cr, minimum
swelling occurred at Ni contents of about 40–45%.
Although void diameters generally appeared to be
smaller in the Ni-based alloys than in austenitic steels,
the main factor accounting for reduced swelling was a
much lower void concentration. In the ternary alloys,
reducing the Ni content from 35% to 30% resulted in

60
Commercial alloys
Fe–15Cr–Ni alloys
50
1


Commercial alloys:
1 Type 304
2 Type 321
3 Type 316
4 Type 318
5 12R72HV
6 A286
7 Incoloy 800
8 Inconel 706
9 Nimonic PE16
10 Hastelloy X
11 Inconel 625
12 Inconel 600
13 Inconel 702

2

40
Swelling (%)

temperatures. At around the same time, Hudson et al.9
compared the swelling behavior of PE16, type 316
steel, and pure nickel, using 20 MeV C2+ ion irradiations in the Harwell VEC (variable energy cyclotron).
The materials were implanted with 10 appm (atomic
parts per million) of helium prior to ion bombardment
to peak displacement doses >200 dpa (N/2) at 525  C.
Void swelling in 316 steel and nickel exceeded 10% at
the highest doses examined, compared to $0.5% in
PE16. Void nucleation appeared to occur earlier
in nickel (at $0.1 dpa) than in PE16 or type 316

($2 dpa), but the peak void concentration was higher
by a factor of about 10 in the austenitic steel than in
nickel or PE16.
Hudson et al.9 originally attributed the swelling
resistance of PE16 to the presence of the coherent,
ordered face-centered cubic, Ni3(Al,Ti) g0 precipitates, which were thought either to trap vacancies and
interstitials at their surface, thereby enhancing pointdefect recombination, or to inhibit the climb of dislocations, thereby preventing them from acting as
preferential sinks for interstitial atoms. In support of
the first of these two suggested mechanisms, Bullough
and Perrin10 argued that the surface of a coherent
precipitate would be a more effective trapping site
than an incoherent one where the identity of the point
defects would immediately be lost (and where, as a
consequence, void nucleation was likely to occur). The
efficiency of point defect trapping would be expected
to be greater the higher the total surface area of the g0
precipitates, that is, to be inversely proportional to
the precipitate size at constant volume fraction. On
the other hand, the second mechanism proposed by
Hudson et al. should be most effective at an intermediate particle size where dislocation pinning is strongest.
Support for the latter process was provided by Williams
and Fisher11 from HVEM (high-voltage electron
microscope) irradiations of PE16 at a damage rate of
about 10À2 dpa sÀ1 at 600  C, where the swelling rate
was higher at small (3 nm) and large (70 nm) g0 particle
diameters than at intermediate sizes of about 20 nm.
However, it is now considered that any effect
that the g0 precipitates may have on the swelling
resistance of Nimonic PE16 is secondary to that of
the matrix composition. The generally low-swelling

behavior of Ni-based alloys compared to austenitic
steels was shown by Johnston et al.12 following bombardment with 5 MeV Ni2+ ions at 625  C. The damage rate in these experiments was 10À2 dpa sÀ1 and
the displacement dose was originally estimated as
140 dpa but this was subsequently revised by Bates
and Johnston13 to 116 dpa (based on displacement

125

3
4

30

5

20

10
6
0

7
8 9 10

0

10

20


30

11

40 50 60
Nickel (wt%)

12 13
70

80

90

Figure 1 Swelling versus nickel content of commercial
alloys and ternary Fe–15Cr–Ni alloys bombarded with Ni2þ
ions to a damage level of 116 dpa at 625  C. Reproduced
from Johnston, W. G.; Rosolowski, J. H.; Turkalo, A. M.;
Lauritzen, T. J. Nucl. Mater.1974, 54, 24–40.


126

Radiation Effects in Nickel-Based Alloys

an increase in overall swelling from <0.1% to $12%,
although it was noted that the 35% Ni alloy showed
a localized swelling of $5% in a region close to a
grain boundary. Additional experiments reported by
Johnston et al.12 indicated that the peak swelling temperature for PE16 irradiated with 5 MeV Ni2+ ions was

675  C, but even then, swelling at 116 dpa remained
below 0.2%.
Swelling data for a wider range of pure Fe–Cr–Ni
austenitic alloys, with Cr contents up to 30% and
Ni up to 100%, following Ni ion bombardment to
116 dpa at 675  C, were reported by Bates and
Johnston.13 These results showed a strong dependence on both Cr and Ni, with the swelling increasing with increasing levels of Cr but being minimized
at Ni contents of about 45–60%. Examination of the
dose dependence of swelling in ternary alloys containing 15% Cr and 20–45% Ni showed that the
incubation dose required for the onset of swelling
increased with increasing Ni content. Furthermore,
although high-swelling rates of the order of 1% per
dpa were attained in 20–35% Ni alloys, the swelling
rate of the 45% Ni alloy remained low even at doses
above 250 dpa.
Following their earlier C2+ ion irradiation experiments, Hudson and coworkers moved to the use of
46.5 MeV Ni6+ ions to investigate void swelling
behavior. This was considered preferable because
the recoil spectra of high-energy Ni ions provided
a better simulation of fast neutron damage, and
because carbon implantation encouraged the formation of carbides which acted as void nucleation sites.
A summary of some of the Ni ion irradiation work
carried out by the Harwell group was given by Makin
et al.14 No significant differences in the swelling
behavior of Nimonic PE16 were evident between
ST or aged conditions. Peak swelling in Ni6+ ionirradiated PE16 (preimplanted with 10 appm He)
occurred at 625  C, where a swelling of $1.5% was
recorded at 120 dpa (N/2). Void concentrations in
PE16 were reported to be lower by a factor of about
5 than in similarly irradiated type 316 and 321 austenitic steels.

A drawback of charged particle irradiation experiments for evaluating void swelling is that the evolution of other microstructural features may differ
significantly from that during neutron irradiation
(see also Chapter 1.07, Radiation Damage Using
Ion Beams). In the case of Nimonic PE16, for example, the precipitation and/or redistribution of the
g0 phase during long-term neutron exposure might
be expected to influence swelling behavior. In order

to simulate swelling in a more appropriate microstructure, Bajaj et al.15 examined the effect of 4 MeV
Ni2+ ion irradiation on PE16, which had been preconditioned by exposure to neutrons in Experimental
Breeder Reactor-II (EBR-II). Reactor-conditioned
samples had been exposed to neutron fluences in
the range of 3–6 Â 1026 n mÀ2 (E > 0.1 MeV) at temperatures from 454 to 593  C. Swelling rates during
Ni ion irradiations at 675  C were higher by a factor
of about five in reactor-conditioned material than in a
nonconditioned sample. The increased swelling rate
was attributed to changes in the matrix composition
resulting from an increased volume fraction of g0 in
the reactor-conditioned material.
Early attempts to account for the effects of matrix
composition on void swelling focused on the stability
of the austenite phase. Harries16 suggested that the
swelling behavior of austenitic steels and nickelbased alloys could be rationalized in terms of their
Ni and Cr equivalent contents (i.e., the relative
austenite and ferrite stabilizing effects of their constituent elements), with the composition of highswelling alloys then falling into the (g þ s) phase
field in the Fe–Cr–Ni ternary phase diagram. Harries
postulated that the partitioning of solute elements
into the sigma phase would have a detrimental effect
on the swelling resistance of austenite. Watkin17 took
a similar approach, but found that an improved
correlation could be obtained using the concept of

electron vacancy numbers rather than Ni and Cr
equivalents. The average electron vacancy number,
Nv, of the matrix is calculated from the atomic fractions of its constituents, with allowance being made
for the precipitation of carbides and g0 (or g00, etc.),
and has been widely used to predict the susceptibility
of nickel-based alloys to the formation of intermetallic phases.18Nv was calculated from:
Nv ¼ 0:66Ni þ 1:70Co þ 2:66Fe þ 3:66Mn
þ 4:66ðCr þ MoÞ
Watkin found that void swelling in a range of alloys
with Ni contents up to $43%, which were irradiated
in DFR to a peak dose of 30 dpa at 600  C, remained
low for Nv below about 2.5 (corresponding to low
susceptibility to s phase formation), but increased
approximately linearly at higher Nv. However, as was
clearly argued by Bates and Johnston,13 correlations
based on sigma-forming tendency could not account
for the minimum in swelling observed at about 45%
Ni, since higher Ni contents should continue to be
beneficial.


Radiation Effects in Nickel-Based Alloys

A better understanding of the swelling behavior of
Fe- and Ni-based alloys resulted from a series of fast
neutron irradiation experiments which were carried
out in EBR-II in the early 1980s. Irradiation temperatures in these experiments ranged from about
400 to 650  C. Initial data for a range of commercial
alloys, including ferritic and austenitic steels, as well
as nickel-based alloys, were reported by Bates and

Powell19 and Powell et al.,20 with higher dose data (up
to a peak fluence (E > 0.1 MeV) of $25 Â 1026 n mÀ2,
corresponding to $125 dpa) being reported by
Gelles21 and Garner and Gelles.22 Swelling data for
Fe–Cr–Ni ternary alloys, irradiated in EBR-II to a
peak fluence of 22 Â 1026 n mÀ2 ($110 dpa), were
presented by Garner and Brager.23 The extent of
void swelling in these experiments was determined
by density change measurements. In general, alloys
with nickel contents in the range of 40–50% exhibited the lowest swelling. Swelling in commercial
nickel-based alloys was generally lower in ST than
in aged conditions, this being attributed to the beneficial (though temporary) effect of minor elements
remaining in solution and being able to interact with
point defects19; subsequent precipitation during irradiation would be expected to reduce this benefit and
the resulting densification, though small, would also
effectively reduce the measured swelling. Swelling
data for a number of ST alloys, which were irradiated
in the AA-1 rig in EBR-II, are shown in Figure 2;
data are shown for two withdrawals, at peak fluences
of 14.7 Â 1026 n mÀ2 and 25.3 Â 1026 n mÀ2, with measurements for Inconel 600 and Inconel 625 reported
at both fluence levels, data for Nimonic PE16 and
Inconel 706 at the lower level, and data for Incoloy
800 and Hastelloy X at the higher level. The nickel
contents of the alloys range from about 34% in Incoloy 800 to 75% in Inconel 600. Swelling remained
relatively low in the three Inconel alloys and in PE16.
However, both Incoloy 800 and Hastelloy X exhibited
high swelling at some temperatures, with swelling in
the latter reaching $80% at 593  C. The reason for
such high swelling in neutron-irradiated Hastelloy
X (nickel content $48%) is unclear, but it was noted

that densification up to 3% occurred at the lower
irradiation temperatures – indicating microstructural
instability and possibly signaling changes in the composition of the matrix which may have affected the
swelling behavior. (Note that Hastelloy X was identified as a low-swelling alloy in the Ni2+ ion irradiation
experiments described by Johnston et al.12)
Some data for different heat-treated conditions of
PE16 at the higher fluence level were reported by

127

Garner and Gelles,22 and are compared for irradiations at 538  C (more or less corresponding to the
peak swelling temperature for PE16 in the AA-1
experiment) with lower fluence data from Bates
and Powell19 in Figure 3. The heat-treated conditions indicated in Figure 3 are ST (ST 4 h at
1080  C), A1 (ST and aged 16 h at 705  C), A2 (ST
and aged 1 h at 890  C plus 8 h at 750  C), and OA
(ST and aged 24 h at 840  C). Note that the silicon
content of the PE16 used in these experiments was
much lower at 0.01% than the level of $0.2% typically found in UK heats of the alloy. Overall, the data
appear to show little effect of initial heat treatment
on the swelling of PE16, except that the OA condition exhibited the most swelling (5.2%) at the higher
fluence.
Although it is clear that the swelling behavior of
austenitic alloys is largely dependent on nickel content, there is ample evidence to show that minor solute
additions can have significant effects. Much of the
work on minor solutes has focused on steels similar
to type 316, but some data are available for higher
nickel alloys. For example, Mazey and Hanks24 used
46.5 MeV Ni6+ ion irradiations to examine the effects
of Si, Ti, and Al additions on the swelling response of

model alloys with base compositions approximating
that of the matrix phase in PE16. Solute additions
of $0.25% Si or 1.2% Ti reduced swelling, but the
addition of $1.2% Al (in the absence of Si or Ti)
markedly increased it. The beneficial effect of Si was
believed to arise from its high diffusivity in solution
(this is discussed further in Section 4.04.2.2), whereas
that of Ti appeared to be related to the formation of Z
phase (hexagonal-structured Ni3Ti). The addition of
Al resulted in an increase in the concentration of
voids, the surfaces of which were coated in a thin
layer of the g0 phase (Ni3Al). A beneficial effect of Si
on the swelling response of modified Incoloy DS alloys
under Ni6+ ion irradiation was also reported by Mazey
et al.25 However, it should be noted that high Si contents can give rise to the formation of radiationinduced phases which are enriched with Ni and Si,
such as the Ni3Si form of g0 and the silicide G-phase
(M6Ni16Si7, where M is usually Ti, Nb, or Mn).
G-phase particles are generally found in association
with large voids and their formation may therefore
give rise to an increase in the swelling rate.26,27
Swelling data derived from density measurements
for neutron irradiated, modified Incoloy DS alloys,
with Si contents ranging from 0.19 to 2.05% (compared to a specified level of 1.9–2.6% in the commercial alloy), are compared with data for a ‘PE16 matrix


128

Radiation Effects in Nickel-Based Alloys

5.0


4.0

Fluence

15

Swelling (%)

3.5
3.0
10
2.5
2.0
1.5
5

1.0
0.5

Fluence (1026 n m−2, E > 0.1 MeV)

In 600
In 625
PE16
In 706

4.5

0.0

−0.5
350

400

450

500

550

600

650

0
700

Temperature (°C)
100

80

In 600
In 625
In 800
Hast X

Fluence


Swelling (%)

70

25

20

60
50

15

40
30

10

20
10

5

Fluence (1026 n m−2, E > 0.1 MeV)

90

0
−10
350


400

450

500

550

600

650

0
700

Temperature (°C)
Figure 2 Void swelling of nickel-based alloys irradiated in AA-1 rig in Experimental Breeder Reactor-II. Based on data
from Bates, J. F.; Powell, R. W. J. Nucl. Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S. In Effects of Radiation
on Materials: 14th International Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S., Eds.; American Society for
Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046.

alloy’ and Nimonic PE16 in Figure 4. The materials
were all in ST condition apart from PE16 which was
in an STA condition (aged 4 h at 750  C). The alloys
were irradiated in the UK-1 rig in EBR-II to fluences
in the range of 9–16 Â 1026 n mÀ2 (E > 0.1 MeV) at
temperatures of $390–640  C. These data are previously unpublished except those for STA PE16
(heat DAA 766) which were reported by Boothby.28
Swelling in the modified Incoloy DS alloys generally

decreased with increasing Si content. The 0.19% Si

alloy exhibited high swelling at all temperatures with
indications of swelling peaks at about 440 and 640  C.
Increased Si levels tended to suppress the high temperature swelling peak and reduce the magnitude of
swelling at lower temperatures. The PE16 matrix
alloy containing 0.24% Si exhibited a high temperature swelling peak but moderate swelling below
$550  C, suggesting a beneficial effect of Mo (this
being the main compositional difference between the
PE16 matrix alloy and the modified Incoloy DS alloys)


Radiation Effects in Nickel-Based Alloys

4.0

at lower temperatures. However, swelling in the PE16
matrix alloy remained significantly higher at all temperatures than in STA Nimonic PE16 (containing
0.15% Si), indicating a significant benefit of the g0
forming elements Al and Ti.

3.0

4.04.2.2

6.0
68 dpa

116 dpa


Swelling (%)

5.0

2.0
1.0
0.0
ST

A1

A2

OA

Figure 3 Effect of heat treatment on void swelling of
Nimonic PE16 irradiated in Experimental Breeder Reactor-II
at 538  C. Adapted from Bates, J. F.; Powell, R. W. J. Nucl.
Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S.
In Effects of Radiation on Materials: 14th International
Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S.,
Eds.; American Society for Testing and Materials:
Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM
STP 1046.

12.0
15

8.0
10

6.0

4.0

5

Fluence (1026 n m–2, E > 0.1 MeV)

10.0

Swelling (%)

129

2.0

0.0
300

350

400

450
500
550
Temperature (°C)

PE16 STA
PE16 Matrix

DS (0.19 Si)

600

0
650

DS (0.56 Si)
DS (1.05 Si)
DS (2.05 Si)

Figure 4 Void swelling data derived from density
measurements for Nimonic PE16, a PE16 matrix alloy,
and modified Incoloy DS alloys, irradiated in the UK-1 rig
in Experimental Breeder Reactor-II. Unpublished data from
Boothby, R. M.; Cattle, G. C. Void Swelling in EBR-2
Irradiated Nimonic PE16 and Incoloy DS; FPSG/P(90)10,
with permission from AEA Technology Plc.

Void-Swelling Models

Point defects created by atomic displacements are
lost either through mutual recombination or by
migration to sinks. Void swelling requires a mobile
population of excess vacancies and can only occur
over a limited temperature range, typically $350–
700  C in neutron-irradiated steels and nickel-based
alloys. Rapid diffusion at higher temperatures reduces
the concentration of radiation-induced vacancies to
near thermal equilibrium levels. Recombination dominates at lower temperatures, where reduced vacancy

mobility prevents the formation of voids as the necessary counter-migration of matrix atoms cannot occur.
In the swelling regime, an increased bias for interstitials over vacancies at dislocation sinks gives rise to the
surplus vacancies which agglomerate to form voids.
The flux of point defects to sinks, including void
surfaces, dislocations, and grain boundaries, results in
the segregation of particular solute atoms at the sinks
and the depletion of others. In austenitic steels and
nickel-based alloys, it is generally found that nickel
segregates at the point defect sinks. This is generally
attributed to the inverse Kirkendall effect described
by Marwick,29 whereby faster diffusing solutes such
as Cr move in the opposite direction to the vacancy
flux and are depleted at the sink, and slower diffusing
solutes such as Ni are enriched. One of the earliest
observations of nickel segregation at void surfaces
due to the inverse Kirkendall effect was made by
Marwick et al.30 in an alloy with a composition corresponding to that of the matrix phase in Nimonic
PE16. (For more detailed discussions on radiationinduced segregation effects, see the reviews of
Wiedersich and Lam,31 and Rehn and Okamoto.32)
Venker and Ehrlich33 recognized that differences
in the partial diffusion coefficients of alloy constituents might account for the effects of composition on
swelling. Any effect of this kind would generally be
expected to be more significant the larger are the
differences between the partial diffusion coefficients
of the alloy components. Garner and Wolfer34 examined Venker and Ehrlich’s conjecture and concluded
that the addition of even small amounts of a fastdiffusing solute such as silicon to austenitic alloys
would greatly increase the effective vacancy diffusion


130


Radiation Effects in Nickel-Based Alloys

coefficient (i.e., would enhance the diffusion rate for
all matrix elements). The overall effect is analogous to
an increase in temperature – resulting in an effective
decrease in the vacancy supersaturation and hence a
reduction in the void nucleation rate. This mechanism is generally accepted as the explanation for the
beneficial effect of silicon in reducing swelling in
austenitic steels and nickel-based alloys. Although
this relies on the diffusion of silicon via vacancy exchange, silicon is also generally observed to segregate
to point defect sinks and since it is an undersized
solute, this is believed to occur by the migration of
interstitial–solute complexes. There is, however, no
reason to suppose that both diffusion mechanisms
cannot operate simultaneously.
Garner and Wolfer34 originally considered that
since nickel diffuses relatively slowly in austenitic
alloys, an increase in nickel content would have the
opposite effect to silicon. However, a later assessment
made by Esmailzadeh and Kumar,35 based on diffusion data reported by Rothman et al.,36 indicated that
the void nucleation rate in Fe–15Cr–Ni alloys would
decrease with an increase in nickel content from 20 to
45%. This result is obtained because, although nickel
remains the slowest diffusing species, the effective
vacancy diffusion coefficient of the system is calculated to increase at the higher nickel content. Esmailzadeh and Kumar’s calculations also confirmed
the beneficial effect of silicon, with the addition of
1% Si predicted to be as effective in suppressing
void nucleation as increasing the nickel content
from 20 to 45%. Effects at nickel contents above

45% could not be examined due to a lack of appropriate diffusion data.
As well as affecting the nucleation of voids, differences in the diffusion rates of the various solutes
might also be expected to influence void growth.
Simplistically, this can be thought of as being partly
due to the segregation of slower diffusing solutes
reducing the rate of vacancy migration in the vicinity
of the voids. However, a further consequence of such
nonequilibrium solute segregation was identified by
Marwick,29 who realized that it would give rise to
an additional vacancy flux which would oppose the
radiation-induced flux to the sink. As discussed by
Marwick, this additional flux (the Kirkendall flux)
may itself be an important factor in limiting void
growth, since it will reduce the probability of vacancy
annihilation at sinks and increase the likelihood of
point defect recombination.
The effect of nickel content on void swelling was
considered further in a model developed by Wolfer

and coworkers.37,38 The model examined the compositional dependence of the void bias and focused on
the effects of nickel segregation at void surfaces.
Wolfer’s model indicated that the compositional gradients produced by radiation-induced segregation
give rise to additional drift forces which affect the
point defect fluxes and thereby modify the bias terms.
These additional drift forces arise from the effects of
composition on point defect formation and migration
energies, on the lattice parameter and the elastic moduli, and from the Kirkendall flux. Wolfer’s calculations
for binary Fe–Ni alloys indicated that the effect of the
Kirkendall flux is small for interstitials but significant
for vacancies. Nevertheless, it was considered that the

overall effect of compositional gradients on the bias
terms is likely to be greater for interstitials than for
vacancies due to other factors, particularly the effect of
variations in the elastic moduli. As noted by Garner
and Wolfer,39 an increase in the shear modulus in the
segregated regions around voids would reduce the bias
for interstitials and therefore help to stabilize voids. It
is difficult to predict the significance of this effect in
complex alloys, however, since depletion of Cr in the
segregated region will tend to reduce the shear modulus, whereas enrichment of Ni in high-Ni alloys will
tend to increase it.38 A more significant result of the
model with regard to the effect of nickel on swelling is
that there is a reversal in the sign of the Kirkendall
force for vacancies in Fe–Ni alloys at $35% Ni. Below
this level, vacancies are predicted to be attracted into
regions of higher Ni concentration, but above it,
the opposite occurs. Wolfer et al. considered that this
effect may account for the dependence of swelling on
Ni content in austenitic alloys containing less than
35% Ni.
A generalized description of the swelling behavior
of austenitic alloys, which was consistent with the
model developed by Wolfer et al., was put forward
by Garner40 (see also Chapter 4.02, Radiation
Damage in Austenitic Steels). Garner’s ideas were
largely based on the results of the EBR-II irradiation
experiments and the earlier ion bombardment work
of Johnston et al., both of which showed a strong
dependence of swelling on nickel content. It was
considered that swelling was characterized by a transient period followed by a regime in which the

swelling rate became constant. In neutron-irradiated
alloys, the swelling rate in the posttransient regime
was generally found to be $1% per dpa. In swellingresistant alloys, however, it was argued that such high
swelling rates might not be observed owing to extended transient periods. The duration of the


Radiation Effects in Nickel-Based Alloys

transient regime was shown to be dependent on alloy
composition and could extend for many tens of dpa in
low-swelling materials. The duration of the transient
regime was implicitly linked to the completion of
void nucleation but, at the time these ideas were
put forward, relatively few measurements of void
concentrations were available, as swelling data were
mainly derived from dimensional or density changes.
Factors that were proposed to account for the
influence of nickel on the void nucleation rate included the effect on vacancy diffusivity described by
Esmailzadeh and Kumar35; a possible correlation with
the development of fine scale compositional fluctuations by a spinodal-like decomposition process
(observed by Dodd et al.41 in ion-irradiated ternary
Fe–Cr–Ni alloys); and an effect of nickel on the
minimum critical radius for the formation of stable
voids.42 Voids are unstable below a critical size, and
will generally shrink unless stabilized by gas atoms;
the minimum stable void radius is dependent on a
number of factors, including temperature and defect
bias, and Coghlan and Garner suggested that the
compositional dependence of the vacancy diffusivity
would also affect this critical size. In other words, it

was considered that the transition from gas bubble to
void would require a larger bubble size in highnickel alloys, particularly at relatively high temperatures in the swelling regime where void nucleation
becomes increasingly difficult. Hoyt and Garner43
subsequently argued that the minimum critical void
radius concept might account for the minimum in
swelling found at the intermediate nickel contents,
provided that a compositional-dependent bias factor
for dislocations was also incorporated into the model.
The compositional dependence of the bias factor
arises from solute segregation, which reduces the
strain energy of dislocations and decreases the ratio
of the bias for interstitials compared to vacancies.
It is of interest that early evidence for the operation of the bubble to void transition was obtained by
Mazey and Nelson,44 who implanted Nimonic PE16
(STA condition) and a PE16 matrix alloy (ST condition) with 1000 appm He to produce a high density
of gas bubbles before subsequent irradiation with
46.5 MeV Ni6+ ions. The PE16 matrix alloy used
in this particular experiment was a low Si variant
(<0.02 wt%) which was known to exhibit relatively
high swelling. The mean bubble size following
helium implantation at 625  C was higher by a factor
of about two in the matrix alloy ($11 nm diameter)
than in the commercial PE16 alloy ($5 nm diameter).
Examination of the alloys following subsequent

131

irradiation also at 625  C revealed high swelling
(12% at 60 dpa) with a uniform distribution of large
voids but no remaining helium bubbles in the matrix

alloy, and low swelling ($1% at 60 dpa) with a
bimodal distribution of bubbles plus voids in the
standard PE16 alloy (see Figure 5). These results
were interpreted as providing evidence for the concept of a critical stable void size, with only a small
fraction of bubbles in the commercial PE16 alloy,
but all of the bubbles in the matrix alloy, being
sufficiently large to grow as voids. Although not
specifically discussed by Mazey and Nelson, the
compositional differences between the two alloys
suggest that the presence of Si and/or the g0 forming
solutes Al plus Ti may help to reduce void nucleation
in PE16.
The belief advanced by Garner,40 that sluggish void
nucleation generally accounted for low swelling in
nickel-based alloys, persisted for some time. However,
data reported by Muroga et al.45,46 largely overturned
this view. Muroga et al. carried out microstructural
examinations of a series of EBR-II-irradiated
Fe–15Cr–Ni ternary alloys with Ni contents ranging
from 15 to 75 wt%, and of archived samples of similar
alloys from the heavy-ion bombardment experiments
of Johnston et al.12 Examination of alloys irradiated in
EBR-II at 510  C showed that the saturation void concentration was dependent on nickel content and was
minimized at $35–45% Ni, but revealed that there was
no increase in void numbers in any of the materials
above a fluence of 2.6 Â 1026 n mÀ2 (E > 0.1 MeV) (see
Figure 6). Alloys containing 19% and 30% Ni exhibited high swelling rates at higher fluences, but swelling
remained low in higher nickel alloys. Similar effects
were found in the ion-bombarded samples, where, for
example, it was shown that there was no significant

change in the void concentration in Fe–15Cr–45Ni
at doses above 50 dpa in irradiations at 675  C, yet a
marked increase in swelling rate occurred above
120 dpa. Thus, contrary to earlier ideas, these investigations clearly demonstrated that the onset of a high
swelling rate was not related to the cessation of void
nucleation. It follows that the transition to a high rate of
swelling must be due to an increase in the growth rate
of existing voids.
Muroga et al.45,46 observed that the total dislocation density in the irradiated Fe–15Cr–Ni alloys was
only weakly dependent on nickel content. This suggested that at the intermediate nickel levels, where
the void concentration was low, dislocations were
weak sinks (for both vacancies and interstitials)
relative to voids. In addition, it was observed that


132

Radiation Effects in Nickel-Based Alloys

20

400
300

16

0.2 dpa

Bubbles


0.2 dpa

12
200

30 dpa

300
200

15
10

Bubbles
Voids

100

5

0

0
Bubbles

300

60 dpa

15


Number of cavities in interval Nc ϫ 10–19 m–3

0

Number of voids in interval Nv ϫ 10–19 m–3

Number of bubbles in interval Nb ϫ 10–19 m–3

8
100

Voids

4
0
16

30 dpa

12
8
4
0
16

60 dpa

12


200

10

100

5

4

0

0

8

0
0
(a)

5

10

15

20

25


30

35

40

45

50

55

0

60

Cavity diameter d nm

(b)

10 20 30 40 50 60 70 80 90 100 110 120 130
Cavity diameter, d (nm)

Figure 5 Histograms showing size distributions of bubbles/voids in (a) solution treated and aged Nimonic PE16 and
(b) solution treated PE16 matrix alloy, irradiated with Ni6þ ions at 625  C to damage levels of 30 and 60 dpa following
implantation with 1000 appm. He (producing $0.2 dpa) at the same temperature. Reproduced from Mazey, D. J.;
Nelson, R. S. J. Nucl. Mater.1979, 85–86, 671–675.

dislocation loops persisted to higher doses at the
intermediate nickel contents, indicating a lower

growth rate for the loops – again implying an effect
of nickel on dislocation sink strength. Based on these
observations, Muroga et al. suggested that a reduced
dislocation bias for interstitials at the intermediate nickel contents might explain the influence of
nickel on the early stages of void development. An
additional factor was required to account for the
eventual transition to a high swelling rate. Microchemical data presented by Muroga et al.46 suggested
that this transition was related to the depletion of
nickel in the matrix owing to its enrichment at void
surfaces.
A complete description which incorporates all of
the composition-dependent factors which affect the
nucleation and growth of voids is lacking. However,
there is a general consensus that the major influence
of alloy composition arises through its effects on the
effective vacancy diffusivity and on segregation arising from the inverse Kirkendall effect. A correlation
between the magnitude of void swelling and radiation-induced segregation was shown for Fe–Cr–Ni

ternary alloys by Allen et al.48 The compositional
dependence of radiation-induced segregation was
determined using a model based on the earlier work
of Marwick,29 which incorporates both the vacancy
flux to the voids and the back-diffusion of vacancies
due to the solute gradients set up by the inverse
Kirkendall effect. Vacancy diffusivities for various
alloy compositions were determined by the measurements of grain boundary segregation in protonirradiated samples. Swelling data for ion and
neutron-irradiated alloys were then compared with
the expected swelling propensity defined by the ratio
of the forward to back diffusion terms calculated at the
appropriate irradiation temperature. The materials

for which vacancy diffusivity data were determined
included Fe-based alloys containing 16–24% Cr and
9–24% Ni, and Ni-based alloys containing 18% Cr
and either zero or 9% Fe. This work did not specifically examine 40–50% Ni alloys corresponding to the
highest swelling resistance, though the results indicated that swelling generally decreased with increasing levels of nickel enrichment and chromium
depletion at void surfaces.


Radiation Effects in Nickel-Based Alloys

133

4.0
Immersion density
Garner and Kumar47

19Ni

3.5

DFR ~ 17 dpa
DFR ~ 80 dpa

3.0

10

35Ni

75Ni

45Ni

Swelling (%)

20

0
Void density (m−3)

30Ni

1021

2

4
6
8
Fluence (1026 n m-2)

1.5

0.0
300

75Ni
30Ni
35Ni
45Ni


0

2.0

0.5

19Ni

1020

2.5

1.0

Fe–15Cr–XNi
510 ЊC

350

400

450 500 550
Temperature (°C)

600

650

600


650

600

650

1022

10

12

Figure 6 Fluence dependence of swelling and void
density of Fe–15Cr–Ni alloys irradiated in Experimental
Breeder Reactor-II at 510  C. Swelling data obtained by
immersion density measurements by Garner and Kumar47
are also shown. Reproduced from Muroga, T.; Garner, F. A.;
Ohnuki, S. J. Nucl. Mater.1991, 179–181, 546–549.

Void concentration (m−3)

Swelling (%)

30

1021

1020

1019


1018
300

4.04.2.3 Swelling Behavior of NeutronIrradiated Nimonic PE16

350

400

450 500 550
Temperature (°C)

100
90
80
Void diameter (nm)

Brown et al.49 compared the swelling behavior of
STA Nimonic PE16 and two cold-worked austenitic
steels (M316 and Nb-stabilized FV548) which were
irradiated in DFR as fuel pin cladding. Two PE16
clad pins were examined, which were irradiated to
burn-ups of 6.1% and 21.6% of heavy atoms,
corresponding to peak damage levels of about 17
and 80 dpa, respectively. Void concentrations and
swelling were lower in PE16 than in the austenitic
steels. Swelling data, void concentrations, and void
diameters for the two PE16 pins examined by Brown
et al. are shown in Figure 7. Note that Brown et al.49

only showed trend lines for void concentration and
void size in the less highly irradiated pin and compared the swelling tendencies of the two pins; the
individual data points were not plotted and those
shown in Figure 7 are previously unpublished data
obtained by Sharpe. Brown et al. stated that the void
concentration in PE16 decreased with increasing
irradiation temperature but did not alter greatly
with an increasing dose above $17 dpa. It should be
noted, however, that swelling measurements for the
higher burn-up pin were restricted to temperatures

DFR ~ 17 dpa
DFR ~ 80 dpa

DFR ~ 17 dpa
DFR ~ 80 dpa

70
60
50
40
30
20
10
0
300

350

400 450 500 550

Temperature (°C)

Figure 7 Swelling data, void concentrations, and void
diameters for Nimonic PE16 Dounreay Fast Reactor fuel pin
cladding. Unpublished data from Sharpe, R. M. Void Swelling
in Fast Reactor Irradiated Commercial High Nickel Alloy;
DFMC/P(82)27, with permission from AEA Technology Plc.

below 525  C, so that a direct comparison of void
concentrations in the two pins cannot be made at
higher temperatures. Although there were fewer
voids in PE16 than in the two steels, the voids


10.0

Swelling (%)

8.0

STA DAA766
OA DAA766
STA Z260D
STA Z184

Fluence

15

10


6.0
4.0

5
2.0
0.0
300

350

400

450

500

550

600

0
650

600

650

600


650

Temperature (°C)
1022
Void concentration (m–3)

appeared to be homogeneously distributed and to have
developed during the early stages of irradiation; once
nucleated, the growth rate of voids in PE16 remained
low. These observations are clearly contrary to early
models which suggested that low swelling rates result
from incomplete void nucleation and extended transient regimes. Rather, in agreement with the more
recent observations of Muroga et al.,45,46 it appears
that the swelling resistance of PE16 is due to a combination of a comparatively low saturation void concentration, which is reached at a relatively low
displacement dose, and a low void growth rate. There
does not appear to be any evidence of an accelerated
swelling rate in PE16 once void nucleation is complete.
Additional data on void concentrations in neutronirradiated PE16 are available from Cawthorne et al.,8
Sklad et al.,50 and Boothby.28 The results presented
by Cawthorne et al. for PE16 fuel pin cladding irradiated in DFR to a peak fluence of 5.6 Â 1026 n mÀ2
($28 dpa) differ from those shown in Figure 7 in
that, although void number densities are similar for
irradiations at $380–520  C, void concentrations are
about an order of magnitude higher at 350  C and
600–630  C. Such discrepancies might arise from
uncertainty and/or variability in irradiation temperatures. Another possibility is that void nucleation was
incomplete at the higher irradiation temperatures in
the lower burn-up pin examined by Brown et al. Data
from Sklad et al. show an increase in void numbers in
unstressed PE16 specimens irradiated in EBR-II at

500  C from an average (for two differently heat treated conditions) of about 4 Â 1019 to 1.2 Â 1020 mÀ3
with increasing fluence from 1.2 Â 1026 to
4.0 Â 1026 n mÀ2 (E > 0.1 MeV), that is, from $6 to
20 dpa. In this case, the void concentration and overall
swelling of $0.2% at $20 dpa remain below the levels
shown in Figure 7 for the DFR-irradiated pin
at $17 dpa; this may reflect the effect of stress on
swelling for fuel pin cladding.
Void swelling data determined from Transmission
electron microscope (TEM) examinations of three
heats of PE16 which were irradiated in the UK-1
rig in EBR-II are shown in Figure 8, which includes
previously unpublished results for the low boron
(4 ppm) heat Z184 as well as data for heats DAA766
and Z260D (with 18 and 70 ppm boron, respectively)
which were reported by Boothby.28 Data are shown
for all three heats in the STA condition (ST 1020  C
and aged 4 h at 750  C) and for DAA766 in the OA
condition (a multistage heat treatment that included
aging at 900  C, slow cooling to 750  C, and then
aging for 16 h at that temperature, resulting in the

Fluence (1026 n m-2, E > 0.1 MeV)

Radiation Effects in Nickel-Based Alloys

1021

STA DAA766
OA DAA766

STA Z260D
STA Z184

1020

1019

1018
300

350

400 450 500 550
Temperature (°C)

100
90
Void diameter (nm)

134

80
70

STA DAA766
OA DAA766
STA Z260D
STA Z184

60

50
40
30
20
10
0
300

350

400 450 500 550
Temperature (°C)

Figure 8 Swelling data, void concentrations, and void
diameters for Nimonic PE16 samples irradiated in UK-1 rig
in Experimental Breeder Reactor-II. Adapted from Boothby,
R. M. J. Nucl. Mater.1996, 230, 148–157; Unpublished data
for Boothby, R. M. The Microstructure of EBR-II Irradiated
Nimonic PE16; AEA TRS 2002 (FPSG/P(90)23), with
permission from AEA Technology Plc.

precipitation of TiC and an overaged g0 structure).
Swelling data derived from the density measurements of STA PE16 heat DAA 766 from the same
experiment are shown in Figure 4. An example of the


Radiation Effects in Nickel-Based Alloys

200 nm
Figure 9 Void structure in PE16 (OA condition)

irradiated in Experimental Breeder Reactor-II to 58 dpa at
513  C. Reproduced from Boothby, R. M. J. Nucl.
Mater.1996, 230, 148–157.

void distribution in the OA condition is shown in
Figure 9. Note that the voids in neutron-irradiated
PE16 tend to be cuboidal and that enhanced growth
of voids attached to TiC precipitates (located at the
site of a prior grain boundary) has occurred.
Neutron fluences and irradiation temperatures
in the UK-1 experiment were similar to those for the
first withdrawal of the AA-1 rig for which data is
shown in Figure 2. Void concentrations for heats
DAA766 and Z260D shown in Figure 8 appear to be
less temperature-dependent than for the fuel pin cladding data shown in Figure 7. Void numbers are generally lower than in the cladding at temperatures
up to $550  C, but are intermediate between the
results of Brown et al.49 and Cawthorne et al.8 for irradiations at $600  C. Void concentrations for PE16
irradiated to fast neutron fluences (E > 0.1 MeV) of
9.4–12.3 Â 1026 n mÀ2 at 477–513  C in the UK-1 experiment were very similar to those determined by
Sklad et al.50 for 4.0 Â 1026 n mÀ2 at 500  C. The low
boron heat Z184 showed atypical behavior, with a very
high concentration of small voids and low swelling at
438  C, but high swelling owing to increased void sizes
at normal void concentrations at temperatures above
513  C. It is probable that the effect of boron on swelling
is related to the formation of boron–vacancy complexes,
which can give rise to the nonequilibrium segregation
of boron in the presence of quenched-in thermal vacancies as well as to radiation-induced effects.51
Some variability in the swelling response of
Nimonic PE16 in PFR (Prototype Fast Reactor)

components was reported by Brown and Linekar.52

135

Increased swelling in PE16 subassembly and guide
tube wrappers in PFR compared to expectations
based on the performance of DFR pin cladding
appeared to be related to temperature fluctuations,
particularly at temperatures below 400  C during the
early operation of PFR. Void concentrations were
reported to be higher in the PFR components, and
it was suggested (by Cawthorne, unpublished data)
that this may have been due to the release of vacancies from vacancy loops which had formed during
lower temperature excursions. In fact, the void concentration reported by Cawthorne et al.8 for DFR pin
cladding irradiated at 350  C was higher than the
highest value reported for the PFR components by
a factor of about 3, but this comparison was not made
by Brown and Linekar. There were also indications of
heat-to-heat variability and effects of the fabrication
route on the swelling of PE16 wrappers in PFR.
Nevertheless, swelling of PE16 wrappers, although
higher than expected, remained low in absolute terms
and did not give rise to any operational problems.
Although PE16 was originally selected as the reference wrapper material for PFR and as an alternative to cold-worked M316 steel for fuel pin cladding,
PE16 was favored as a cladding material with 12%Cr
ferritic–martensitic steel wrappers in subsequent
subassembly designs.53 The 12%Cr steel was chosen
as a wrapper material because of its superior swelling
resistance, but its use was limited to relatively low
temperatures owing to inadequate strength at the

higher operating temperatures experienced by pin
cladding. Design calculations for PE16 fuel pin cladding made by Cole54 indicated that cladding hoop
stresses, which arise from the internal pressure from
the gaseous fission products released from the fuel,
were much lower than the yield stress of the material
and were generally expected to remain below about
70 MPa. In addition, the void swelling and irradiation
creep behavior of PE16 were considered to be well
matched to the fuel swelling, so that fuel–clad interaction stresses also remain low. Fuel pins with PE16
cladding successfully attained high burn-ups in PFR,
with some 3500 pins exceeding dose levels of 100 dpa
and 265 pins reaching maximum doses of 155 dpa.55
Very few failures of PE16 clad pins were recorded –
three failures occurred in pins which had reached
burn-ups over 17 at.%, with one failure at 11.3 at.%
burn-up which was believed to have resulted from
a fabrication defect.56 In addition to the four PE16
cladding failures in PFR, Plitz et al.57 recorded 14
failures in austenitic steel cladding, all at lower burnups than in PE16. The failures in PE16 cladding were


136

Radiation Effects in Nickel-Based Alloys

regarded as benign and permitted continued operation, with no significant loss of fuel into the primary
circuit coolant. A peak burn-up of 23.2 at.%,
corresponding to a peak dose in the PE16 cladding
of 144 dpa, was achieved in PFR in an experimental
fuel cluster. Postirradiation examinations of pins

from this cluster and a high burn-up subassembly
(18.9 at.%, with a peak cladding dose of 148 dpa)
were carried out by Naganuma et al.58 Maximum
diametral strains of less than 1% were measured,
attributable to the combined effects of void swelling,
creep deformation arising from internal gas pressure
in the pins, and small contributions from mechanical
interactions between the fuel and cladding in the
lower part of the pins.

4.04.3 Irradiation Creep
A detailed discussion of irradiation creep mechanisms is beyond the scope of this chapter, which will
instead concentrate on experimental data which
enable comparisons to be made between nickelbased alloys and austenitic steels. However, some
insight into irradiation creep mechanisms is given
in Section 4.04.4.1, where the effect of stress on the
evolution of dislocation structures is described. Irradiation creep mechanisms are discussed more fully
in Chapter 1.04, Effect of Radiation on Strength
and Ductility of Metals and Alloys. Several reviews
of irradiation creep data are available in the literature,
for example, by Harries,59 Ehrlich,60 and Garner,61
and although these have tended to focus on austenitic
steels, the behavior of nickel-based alloys generally
appears to be similar.
Different types of test specimen, including pressurized tubes and helical springs, have been used to
measure irradiation creep strains. The data are
therefore generally converted to effective strain e
 values, using the Soderberg
and effective stress s
formalism60:

e=
s ¼ e=s ¼ g=3t ¼ 4eH =3sH
where e, g, and eH are tensile, surface shear and
hoop strains; and s, t, and sH are tensile, surface
shear and hoop stresses, respectively.
Irradiation creep experiments carried out in DFR
used helical spring specimens, which were loaded in
tension and periodically removed for measurements.
DFR data for austenitic steels and Nimonic PE16
were reviewed by Mosedale et al.62 and Harries,59
and results for PE16 were reported in full by

Lewthwaite and Mosedale.63 Average irradiation
temperatures for PE16 specimens ranged from
about 280 to 340  C, with displacement doses up to
a maximum of $13 dpa (N/2). For austenitic steels,
the irradiation creep strain was found to be linearly
dependent on the applied stress and the displacement
dose, comprising transient and steady-state components as follows:
g ¼ At þ Bd t
where d is the displacement dose and A and B are
material-dependent creep coefficients. For PE16 in
a STA condition (1 h at 1080  C plus 16 h at 700  C),
creep at dose rates of $5 Â 10À7 dpa (N/2) sÀ1 was
characterized by an initial period of low strain and an
increased creep rate at higher displacement doses.
Mosedale et al.62 described the g=t versus dpa creep
curve for STA PE16 as parabolic, though the maximum observed creep rate was similar to that in austenitic steels and Harries59 represented the creep
strain above a threshold dose of 8 dpa (N/2) by
g ¼ 4:3 Â 10À6 tðd À 8Þ

where t is in MPa; converting to effective strain/
stress values and to NRT units of displacement dose
(assuming 1 dpa (N/2) ¼ 0.8 dpa (NRT-Fe)) would
reduce the creep coefficient by a factor of 2.4. Data
presented by Lewthwaite and Mosedale63 showed
that ST PE16 behaved similarly to the STA condition, though OA conditions exhibited higher creep
strains due to a combination of increased creep rates
and low threshold doses (around 1 dpa). An apparent
dose-rate dependency was observed, with steadystate creep coefficients for STA and OA PE16
increased by factors of $2 at lower damage rates of
$0.5–1.5 Â 10À7 dpa (N/2) sÀ1 and threshold doses
reduced to $0.5 dpa or less. A similar effect of
dose rate on the creep strain per dpa was also
reported for austenitic steels.64 Steady-state creep
coefficients (MPaÀ1 dpaÀ1) and creep strain rates
(MPaÀ1 sÀ1) for PE16 as a function of dose rate are
compared with data for cold-worked steels M316 and
FV548 in Figure 10. The data plotted in Figure 10
are derived from the results of Lewthwaite and Mosedale63,64 but are converted to effective strain/stress
values and NRT(Fe) dpa units to enable comparison
with other published data. It is evident that the irradiation creep behavior of STA and OA (24 h at
800  C) PE16 is similar to that of the austenitic steels.
Creep rates at higher dose rates are generally lower
than would be indicated from the linear extrapolation
of low dose rate data. Lewthwaite and Mosedale63


Radiation Effects in Nickel-Based Alloys

Creep coefficient, B (10–6 MPa–1 dpa–1)


6.0
M316
FV548
PE16 OA
PE16 STA

5.0

4.0

3.0

2.0

1.0

0.0
0.0

1.0

2.0

3.0

4.0

5.0


Displacement rate (10–7 dpa s–1)
10.0

Creep strain rate (10–13 MPa–1 s–1)

9.0
8.0

M316
FV548
PE16 OA
PE16 STA

7.0
6.0
5.0
4.0
3.0
2.0
1.0
0.0
0.0

1.0

2.0

Displacement rate

3.0


4.0

5.0

(10–7 dpa s–1)

Figure 10 Steady-state creep coefficients and creep
strain rates for Nimonic PE16 and austenitic steels, derived
from the measurements of Lewthwaite and Mosedale.
Adapted from Lewthwaite, G. W.; Mosedale, D. In
Proceedings of International Conference on Irradiation
Behaviour of Metallic Materials for Fast Reactor Core
Components, Ajaccio, Corsica, June 4–8, 1979; Poirier, J.,
Dupouy, J. M., Eds.; Le Commissariat a l’Energie
Atomique (CEA): Saclay, France, 1979; pp 399–405;
Lewthwaite, G. W.; Mosedale, D. J. Nucl. Mater. 1980, 90,
205–215.

considered that the measured irradiation creep rates
for PE16 at low dose rates were in close agreement
with the expected rates for SIPA-(stress-induced preferred absorption of interstitials at dislocations) controlled creep. It was suggested by Mosedale et al.62

137

that reduced creep rates at higher dose rates might
be attributable to increased recombination rates for
vacancies and interstitial atoms, although a more
detailed assessment of this effect by Lewthwaite and
Mosedale64 proved inconclusive and a dose-rate dependency has not generally been observed in other

experiments.60 Garner and coworkers65,66 considered
that the higher creep rates measured by Lewthwaite
and Mosedale at lower displacement rates were an
aberration due to transient effects at low dpa levels.
Nevertheless, this does not alter the finding that the
irradiation creep behavior of PE16 is comparable to
that of austenitic steels.
Paxton et al.67 examined the in-reactor creep
behavior of a number of alloys, including Nimonic
PE16, Inconel 706, and Inconel 718, as well as austenitic and ferritic steels, in pressurized tube experiments carried out in EBR-II at 540  C to fluences
up to 4 Â 1026 n mÀ2 (E > 0.1 MeV). Diametral strains
measured in pressurized tubes (with hoop stresses
in the approximate range of 25–175 MPa) were corrected for void swelling and/or densification observed
in unstressed specimens (though this does not allow
for any effects of stress on swelling or precipitation
processes). Precipitation-hardened alloys exhibited
lower creep strains than solid solution strengthened
steels, with the Inconel alloys superior to PE16 at
540  C. The creep resistance of the precipitationhardened materials was also dependent on heat treatment, with ST conditions generally superior to aged
conditions. However, it was noted that ST conditions
also exhibited greater densification – giving rise to
the possibility of increased fuel–clad interactions in
fuel elements. In-reactor creep strains were discussed
in terms of a widely used model which includes a
term for creep enhancement due to swelling. The
total effective creep strain e is given by
e ¼ B0 ft s
 þ DS
s
where B0 is the creep compliance, ft is the neutron

fluence, D is the creep–swelling coupling coefficient,
and S is the fractional swelling. A contribution from
thermal creep may be expected at 540  C, but data to
correct for this component were not available and
hence the creep coefficients could not be determined
precisely. The stress dependence of the measured
creep strain was approximately linear in the low
swelling precipitation-hardened alloys, though nonlinearity attributable to the effects of stress on
swelling was observed in the solid solution alloys.
An approximate value of B0 of 1.5 Â 10À28 MPaÀ1
(n cmÀ2)À1, which is equivalent to 3 Â 10À7 MPaÀ1


138

Radiation Effects in Nickel-Based Alloys

dpaÀ1, was derived by Paxton et al. for the Inconel
alloys. Ehrlich60 subsequently made estimates of B0
for the other materials included in this study, which
ranged from $1.4 Â 10À6 MPaÀ1 dpaÀ1 for ST PE16
to $10À5 MPaÀ1 dpaÀ1 for cold-worked 316 steel.
Paxton et al. noted that values of the creep–swelling
coefficient D appeared to be much larger for the solid
solution strengthened steels than for the precipitationhardened alloys, with the higher values being attributable to increased thermal creep components and/or
the effects of stress on swelling.
Gilbert and Chin68 examined the nonisothermal
creep behavior of EBR-II-irradiated PE16 and
Inconel 706. Both materials were in ST conditions.
Pressurized tubes, with nominal hoop stresses of

100 MPa for PE16 and 200 MPa for Inconel 706,
were irradiated at 425, 540, and 590  C, both isothermally and with temperature steps. Diametral strains
for isothermally irradiated PE16 increased with
increasing fluence and temperature as expected. Following temperature changes from 540 to 590  C or
425  C, the creep rate for PE16 adjusted to the isothermal rate at the new temperature. For Inconel 706,
however, the isothermal creep rate was highest at
425  C, and an upward step to 540  C resulted in a
reduced creep rate; a downward step from 540 to
425  C gave rise to an increased creep rate that
exceeded the isothermal rate at 425  C; and an
upward step from 540 to 590  C reduced the creep
rate, even though the isothermal creep rate was
higher at 590 than 540  C. The complex in-reactor
creep behavior of Inconel 706 appeared to be related
to the stability of the ordered body-centered tetragonal, Ni3Nb g00 phase and its effect on thermal creep
resistance. Gilbert and Chin considered that the inreactor deformation of Inconel 706 was primarily
controlled by thermal rather than irradiation creep
processes, since similar creep rates were reported
to occur in thermal control tests. Microstructural
examinations made by Thomas69 indicated that g00
precipitated during irradiation above $500  C but
dissolved at lower temperatures, thereby reducing
the creep strength of the material. Gelles70 subsequently reported that the dissolution of g00 at low
irradiation temperatures appeared to be promoted
by the application of stress since more of this phase
was retained in unstressed material.
Toloczko et al.5 investigated the swelling and
creep behavior of five austenitic alloys which were
irradiated in PFR in a pressurized tube experiment
at $420  C. The materials examined included the

solid solution strengthened steels 316 and D9, and

the higher-Ni precipitation-hardened alloys D21,
D68, and D66. Dose rate variations were examined
by positioning specimens at different axial locations within the reactor core. The tubes were
removed periodically for diameter measurements,
with peak doses of $50 dpa being attained at the
highest flux level. Hoop stresses ranged from 0 to
150 MPa, and swelling as a function of dose was
estimated from measurements on unstressed tubes
assuming that densification effects were completed
during the first irradiation cycle. There was some
scatter in the results but the creep coefficient B0
was found to be relatively independent of alloy
composition and dose rate, with typical values
of $1.0–1.4 Â 10À6 MPaÀ1 dpaÀ1 (though higher
values were determined for type 316 steel). The
creep–swelling coupling coefficient D was also
independent of dose rate but appeared to be material dependent (with values in the approximate
range of 0.4–1.6 Â 10À2 MPaÀ1), though this variability could not be associated with any particular
compositional factor. Similar results for two precipitation-hardened high-nickel alloys (with similar compositions to Nimonic PE16, but with
additions of $0.5% Nb), which were irradiated in
a pressurized tube experiment in the Russian fast
reactor BN-350 to $90 dpa at 400  C, were also
reported by Porollo et al.71

4.04.4 Microstructural Stability
4.04.4.1

Dislocation Structures


Dislocation structures in irradiated pressurized tube
samples were examined by Gelles et al.72 The materials which were examined included stressed and
unstressed samples of ST PE16, and stressed samples
of ST and STA Inconel 706. A subsequent paper by
Gelles73 extended these investigations to the stressed
samples of PE16 in STA and OA conditions. Further
details of this work were also provided by Garner and
Gelles74, and by Gelles.70
Examination of ST PE16, which was irradiated at
550  C to 2 Â 1026 n mÀ2 (E > 0.1 MeV) at hoop stresses of 0 and 167 MPa, revealed that the distribution of
Frank dislocation loops was similar on all the four
{111} planes in the unstressed sample but was anisotropic in the stressed material. In the stressed sample,
the loop density on any particular {111} plane
increased with increasing magnitude of the normal
stress component on that plane. A near-linear relationship between the loop density and the normal


Radiation Effects in Nickel-Based Alloys

component of the deviatoric stress tensor, sDN
(¼ sN À sH , where sN is the normal component of
the applied stress on a particular plane and sH is the
hydrostatic stress), was found for PE16. This result is
in line with the SIPA loop growth model described by
Garner et al.75 No such correlation was found in the
similarly irradiated and stressed Inconel 706 samples,
however, possibly because the low creep rate of this
material at 550  C did not allow the relaxation of
internal stresses.

Unfaulting of Frank dislocation loops with a/3
{111} Burgers vectors proceeds via interaction with
a/6{112} Shockley partials to produce perfect a/2
{110} line dislocations. Gelles70 described how this
occurs via a two-step process, with the necessary
partial dislocations (two per interstitial loop) first
being nucleated by an interaction of the faulted
loop with a suitable perfect dislocation and then
sweeping across the loop to reestablish the perfect
dislocation. Gelles73 examined the distribution of
Burgers vectors among the six possible a/2{110}
perfect dislocation types in irradiated pressurized tube samples of PE16. The samples examined
included the stressed ST condition irradiated at
550  C, and STA and OA conditions which were
both irradiated at 480  C to a fluence of
8 Â 1026 n mÀ2 at a hoop stress of 331 MPa. The
results showed highly anisotropic distributions in
the Burgers vectors of perfect dislocations in all the
three heat-treated conditions, with dislocation densities of the various types differing by factors of up to
10–40 in each sample. The level of anisotropy produced in the population of perfect dislocations was
significantly greater than in the dispersion of Frank
loops. This is a feasible outcome since, in principle,
all loops may be unfaulted by just two variants of the
six a/2{110} perfect dislocation types. In effect, the
development of anisotropic dislocation structures
is a response of the material to produce the strain
which is required to accommodate the applied stress.
Furthermore, it was found that the perfect dislocations in the irradiation creep samples of PE16 were
primarily of edge type lying on {100} planes rather
than {111} slip planes, indicating that they could

only contribute to the creep strain via climb (i.e., by
the SIPA mechanism) and not by processes involving
dislocation glide.
4.04.4.2

Precipitate Stability

Early models of precipitate stability under irradiation
were based on the ideas of Nelson et al.,76 who

139

suggested that precipitates would evolve to an equilibrium size determined by competing processes
affecting their growth, via the radiation enhanced
and/or thermal diffusion of solutes, and their simultaneous dissolution due to damage arising in collision
cascades. Two dissolution mechanisms were suggested: recoil dissolution due to the displacement of
atoms from the precipitate into the matrix, and disordering dissolution of ordered phases such as g0 ,
with the latter predicted to be the more effective.
The model predicted that fine precipitates would
continue to grow to some equilibrium size (dependent on temperature, dose rate, and solute levels), but
that precipitates greater than this size would shrink.
Experimental evidence for the dissolution of large
preexisting Ni3Al g0 precipitates in heavy-ionirradiated Ni–Al alloys was shown by Nelson et al.76
These ideas were developed further and applied
to g0 precipitates in ion and neutron-irradiated alloys
by Baron et al.77 The model developed by Baron et al.
indicated that, at a given particle size, a higher
solute supersaturation was required under irradiation
than in a purely thermal environment. The model
appeared to be consistent with the observed coarsening behavior of g0 precipitates during irradiation,

though no evidence for the shrinkage of large particles
was presented. For example, data for PE16 irradiated
at fluences up to 7.5 Â 1026 n mÀ2 at 560  C, which
were reported by Chang and Baron,78 only examined
the growth of g0 particles up to a maximum radius
of $15 nm under conditions where the predicted
maximum equilibrium radius was $35 nm.
However, detailed examinations of g0 structures in
neutron-irradiated Nimonic PE16 which were made
by Gelles79 found no evidence to indicate that irradiation-induced dissolution mechanisms limited the
particle size. Microstructural examination of PE16,
originally in ST, STA, and OA conditions, irradiated
in EBR-II to $27 dpa (5.4 Â 1026 n mÀ2, E > 0.1 MeV)
at 600  C, revealed that preexisting g0 dispersions in
aged material were maintained but continued to
coarsen even in the OA condition, and that a fine
dispersion formed in ST material. Coarsening of
the g0 particles in the OA material was accompanied
by the formation of fine background precipitates in
some regions. Further in-reactor precipitation of g0
also occurred at point defect sinks, including void
surfaces and dislocations, in all the heat-treated conditions. Additional examinations by Gelles80 of ST
PE16, irradiated to $30–50 dpa at temperatures in
the range of 430–650  C, indicated that g0 coarsening
was controlled by radiation-enhanced diffusion


140

Radiation Effects in Nickel-Based Alloys


below 600  C with an activation energy that (in
agreement with theoretical predictions for a process
governed by point defect recombination) was approximately a quarter of that for thermal diffusion.
As described in Section 4.04.5.1 in relation to
irradiation embrittlement effects, Yang81 examined
an identically irradiated set of ST PE16 samples as
Gelles, focusing on the precipitation of g0 at grain
boundaries. Similar g0 structures to those described
by Gelles and Yang were also observed by Boothby28
in the aged conditions of EBR-II-irradiated PE16,
though at higher irradiation temperatures (!540  C
for the STA condition, and !600  C for the OA
condition), where doses were in the range 66–74 dpa,
the spherical g0 precipitates which formed during
thermal aging were almost entirely replaced by ‘skeletal’ forms nucleated at point defect sinks. Figure 11
shows an example of the g0 distribution, imaged in
dark field, in STA PE16 irradiated to 69 dpa at
570  C; although small spherical precipitates were
retained in a narrow region adjacent to the grain
boundary, a much coarser dispersion is evident at
the boundary itself and within the bulk of the grain.

4.04.5 Irradiation Embrittlement
The effects of fast neutron irradiation on the tensile
properties of several precipitation-hardened nickelbased alloys were investigated in the 1970s and 1980s.

The materials examined included a number of g0 /g00 hardened alloys, such as the Inconel alloys 706 and
718 and the developmental alloys D68 and 7818,
as well as g0 -hardened alloys similar to Nimonic

PE16. Earlier work by Broomfield et al.82 on thermal
reactor irradiated materials indicated that PE16
was more susceptible to irradiation embrittlement
at elevated test temperatures than austenitic steels.
Broomfield83 found that thermal neutron irradiated
PE16 was most severely embrittled in low strain tests
at $550–650  C, and attributed this to an increased
tendency for intergranular failure arising from the
effects of helium generated from the 10B(n,a)7Li
reaction. Boron itself is considered to have a beneficial effect on (unirradiated) creep rupture life, as it
segregates to grain boundaries and inhibits intergranular cracking, and additions of a few 10s of ppm are
therefore, generally made to nickel-based alloys,
including PE16.84 Nickel is also a major source of
helium in neutron-irradiated alloys, with the twostage 58Ni(n,g)59Ni(n,a)56Fe reaction becoming the
dominant source at high thermal neutron fluences,
and nickel threshold reactions accounting for the
greater part of helium production in fast neutron
spectra.85 For example, the rate of helium generation
in fast reactor irradiated PE16 was estimated by
Boothby28 to be $1.2 appm per dpa, with about
85% of the helium being generated from nickel
threshold reactions (see also Chapter 1.06, The
Effects of Helium in Irradiated Structural Alloys).
Nevertheless, other factors, including irradiationinduced strengthening and grain boundary segregation and precipitation effects, have been implicated
in
the
embrittlement
of
fast
neutron

irradiated nickel-based alloys.

4.04.5.1 Fast Neutron Irradiation
Experiments

200 nm

Figure 11 Dark field, transmission electron
micrograph, illustrating the distribution of g0 precipitates
in solution treated and aged Nimonic PE16 irradiated
in Experimental Breeder Reactor-II to 69 dpa at 570  C.
Unpublished data from Boothby, R. M. The
Microstructure of EBR-II Irradiated Nimonic PE16; AEA
TRS 2002 (FPSG/P(90)23), with permission from AEA
Technology Plc.

Rowcliffe and Horak86 investigated the tensile properties of Inconel 706 (in a multistep ‘fully aged’
condition) and Inconel 718 (ST condition) following
irradiation in EBR-II to fluences of 4–5 Â 1026 n mÀ2
(E > 0.1 MeV). Irradiation temperatures (Ti) ranged
from 450 to 735  C, with tensile tests being performed at a strain rate of 4 Â 10À4 sÀ1 at temperatures
corresponding to Ti and to Ti þ 110  C. Yield stresses
and total elongation data for Inconel 706 are shown in
Figure 12 and for Inconel 718 in Figure 13. Data for
Inconel 706 showed very high (>1000 MPa) yield
stresses and ultimate tensile strengths (UTS) in


Radiation Effects in Nickel-Based Alloys


specimens irradiated at temperatures up to and
including 500  C. This high tensile strength was
maintained in a specimen irradiated at 500  C but
tested at 610  C. Although there was some reduction
in strength in specimens irradiated at 560  C and
above, the UTS remained above 650 MPa in specimens irradiated at 625  C. The very high tensile

strengths exhibited at the lower irradiation temperatures were attributed to the instability of the (ordered
body-centered tetragonal) g00 phase below 525  C and
its consequent dissolution, leading to the reprecipitation of nickel and niobium as (ordered face-centered
cubic) g0 on dislocation loops. At higher irradiation
temperatures, both g0 and g00 were stable, but

1400

20
Yield at Ti + 110 °C
Elong. at Ti + 110 °C

1200

Yield at Ti

18
16
14
12

800


10
600

8
6

400

Total elongation (%)

Elong. at Ti

1000
Yield stress (MPa)

141

4
200
2
0
400

450

500

550

600


650

700

750

0
800

Temperature (°C)
Figure 12 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110  C for Experimental
Breeder Reactor-II-irradiated Inconel 706. Based on data from Rowcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans. 1981,
38, 266–267.

20

1400
Yield at Ti
Elong. at Ti

1200

Yield at Ti + 110 °C
Elong. at Ti + 110 °C

18

Yield stress (MPa)


1000

14
12

800

10
600

8
6

400

Total elongation (%)

16

4
200
2
0
400

450

500

550

600
650
Temperature (°C)

700

750

0
800

Figure 13 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110  C for Experimental
Breeder Reactor-II-irradiated Inconel 718. Based on data from Rowcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans.1981,
38, 266–267.


142

Radiation Effects in Nickel-Based Alloys

precipitate coarsening resulted in lower tensile
strength. Elongations to failure for tests carried out
at the irradiation temperature were between 1.5%
and 3% up to 625  C, compared to >8% in unirradiated material. Irradiation embrittlement was generally more severe in tests at Ti þ 110  C, particularly
at 610–735  C where the lowest recorded ductility
was 0.2%. Fractures in irradiated Inconel 706 were
predominantly intergranular, with failure believed to
be facilitated by the decohesion of Z phase (hexagonal Ni3(Ti,Nb)) platelets which were formed at grain
boundaries during the initial heat treatment.
Rowcliffe and Horak’s data for ST Inconel 718

showed similar trends to Inconel 706. Precipitation
of the g0 and g00 phases occurred during the irradiation of Inconel 718, resulting in yield strengths in
excess of 1000 MPa at irradiation temperatures up to
560  C and above 800 MPa at 625  C. The ductility of
Inconel 718 was reduced from more than 30% in the
unirradiated condition to 0.2% or less in specimens
which were irradiated at 500–560  C and tested at
Ti þ 110  C. In contrast to Inconel 706, failures in
irradiated Inconel 718 were reported to be predominantly transgranular. Crack propagation in Inconel
718 appeared to have been via a ‘channel’ fracture
mechanism, that is, with deformation occurring by
highly localized planar slip and consequent linkage of
radiation-induced voids.

Bajaj et al.87 examined the tensile properties of
Nimonic PE16 irradiated in EBR-II to neutron
fluences up to a maximum of $7 Â 1026 n mÀ2
(E > 0.1 MeV), at temperatures in the range of 450–
735  C. The alloy was in a STA (1 h at 900  C plus 8 h
at 750  C) condition, and appears to have been the
same low-Si heat of PE16 that was subsequently used
in the AA-1 swelling experiment described by Garner
and Gelles.22 Tensile tests were carried out at 232  C
(to simulate refueling conditions), at the irradiation
temperature Ti and at Ti þ 110  C (to simulate reactor transients), at a strain rate of 4 Â 10À4 sÀ1, and
with a small number of tests at 4 Â 10À3 sÀ1. Irradiated specimens tested at 232  C generally showed
a substantial increase in yield stress and a small
increase in UTS over the unirradiated values
(although samples irradiated at the highest temperature of 735  C exhibited some softening), and
retained good levels of ductility with total elongation values above 10%. Yield stress and total elongation data for PE16 at higher test temperatures are

shown in Figure 14 for specimens irradiated to a fast
neutron fluence of 4.3 Â 1026 n mÀ2 (enabling direct
comparison with the data for the similarly irradiated
Inconel alloys shown in Figures 12 and 13). Specimens tested at the irradiation temperature again
showed strengthening at temperatures in the range of
450–625  C and softening at 735  C, with good
20

1000
Yield at Ti + 110 °C
Elong. at Ti + 110 °C

800

Elong. at Ti

18
16

Yield stress (MPa)

14
12

600

10
8

400


6

Total elongation (%)

Yield at Ti

4

200

2
0
400

450

500

550
600
650
Temperature (°C)

700

750

0
800


Figure 14 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110  C for Experimental
Breeder Reactor-II-irradiated Nimonic PE16. Based on data from Bajaj, R.; Shogan, R. P.; DeFlitch, C.; et al. In Effects
of Radiation on Materials: 10th Conference; Kramer, D., Brager, H. S., Perrrin, J. S., Eds.; American Society for Testing
and Materials: Philadelphia, PA, 1981; pp 326–351, ASTM STP 725. Reprinted, with permission, from ASTM STP725-Effects
of Radiation on Materials, copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428.


Radiation Effects in Nickel-Based Alloys

ductility at 450  C but total elongations reduced
to $3% at 560–625  C. Tests at Ti þ 110  C showed
further increases in tensile strength (consistent with the
greater hardening expected from irradiation at a lower
temperature) and more severe embrittlement with
ductility levels at 670–735  C reduced to 0.3% at
a fluence of 4.3 Â 1026 n mÀ2 and to zero (i.e., failure
before yield) in higher dose samples (7.1 Â 1026 n mÀ2).
Tests at Ti at the higher strain rate resulted in an
improvement in ductility by a factor of two or three.
Examination of fracture surfaces showed that failures
were predominantly intergranular in irradiated samples tested above $550  C, transgranular at 232  C,
and mixed mode at 450–550  C. Bajaj et al. considered
that the irradiation embrittlement of PE16 evident at
high temperatures could simply be explained by matrix
hardening with little or no change in the grain boundary fracture strength – evidenced by increases in yield
strength but no significant changes in true (as opposed
to engineering) UTS values – so that mechanisms
relying on the weakening of grain boundaries could
be discounted for the test conditions studied.

Sklad et al.50 reported tensile data for two aged
conditions of Nimonic PE16 which were irradiated
in EBR-II to 1.2 Â 1026 n mÀ2 (E > 0.1 MeV) at
500  C and tested at strain rates from $3 Â 10À5 to
3 Â 10À3 sÀ1. There was no significant difference in
the postirradiation properties of the two differently
aged conditions, although one aging treatment (2 h at
800  C plus 16 h at 700  C) resulted in an unirradiated
yield stress $25% higher than the other condition (1 h
at 900  C plus 8 h at 750  C). No effect of strain rate on
tensile properties was evident in tests at the irradiation
temperature, where total elongations remained above
10%. Tests at higher temperatures were made only at
the lowest strain rate, with failure elongations being
reduced to 1.6% at 600  C and 0.5% at 700  C. The
low ductility failures were associated with an increased
tendency toward intergranular fracture, and additional
tests, in which samples irradiated to 4 Â 1026 n mÀ2 at
500  C were fractured in situ in an Auger spectrometer,
revealed helium release from samples which fractured
intergranularly as well as the segregation of Ni, P, and
S to grain boundaries. Helium release was estimated
at $0.03 He atoms per grain boundary atom. No grain
boundary helium bubbles were observable by TEM,
and it was therefore considered that helium either
remained in solution as a partial monolayer or was
present in unresolved bubbles less than 1–2 nm in
diameter.
The presence of grain boundary helium bubbles in
Nimonic PE16 was reported by Fisher et al.88 in


143

sections of AGR (advanced gas-cooled reactor) tie
bars irradiated at 512  C and above. AGR tie bars,
which are approximately 10 m long and are under
load only during charging and discharging of the
fuel element stringers, operate at temperatures from
325 to 650  C from bottom to top, with peak doses
of $3 dpa occurring at around the 4 m position.
Stress-rupture testing at 600  C at an applied stress
of 500 MPa showed a trough in properties (i.e., a
minimum in failure times) and intergranular failures
in sections of some tie bars which were irradiated
at temperatures in the range of 350–400  C where
grain boundary helium bubbles were not generally
observed. Even so, grain boundary cavitation was
observed in a fractured tie bar section which was
irradiated at 360  C, with the cavities appearing to
be nucleated (possibly at submicroscopic helium
bubbles) at the intersection of slip bands with the
boundary. The trough in stress-rupture properties
occurred in tie bar sections which exhibited both
high yield strengths (attributable to high concentrations of dislocation loops and small voids) and high
levels of grain boundary segregation. EDX (energy
dispersive X-ray) analyses showed a significant
enrichment of Ni and Si, and a depletion of Fe, Cr,
and Mo, at the grain boundaries of sections irradiated
at 335–585  C. In addition, high levels of Si were
detected in sections irradiated at 335–512  C in the

g0 phase that precipitated at the surface of voids, with
the Si content increasing with decreasing irradiation
temperature. Although the presence of Si-enriched g0
phase at grain boundaries could not be confirmed, it
was suggested that its formation may have contributed to the minimum in stress-rupture life, which was
thought to result from the weakening of the boundaries relative to the matrix.
Grain boundary helium bubbles were also observed
by Boothby and Harries89 and Boothby28 in PE16
irradiated in DFR and EBR-II at 535  C and above.
Tensile testing of DFR-irradiated PE16, exposed
to $20 dpa at 465–635  C, and strained at a rate of
2.5 Â 10À6 sÀ1 at temperatures approximating those of
irradiation, revealed severe embrittlement with minimum elongations of $0.2% at 550  C; TEM examination of strained specimens provided evidence of
intergranular cavitation, and the ductility data were
interpreted using a model for the diffusion-induced
growth of cavities nucleated at grain boundary helium
bubbles.89
The postirradiation tensile properties and microstructure of developmental g0 (D21, D25, and D66)
and g0 /g00 (D68) strengthened alloys were discussed


144

Radiation Effects in Nickel-Based Alloys

by Yang et al.4 The alloys were all irradiated in a ST
condition; additionally, D25 was irradiated in an aged
(24 h at 700  C) condition (STA), and D66 in a 30%
cold-worked plus aged (11 h at 800  C plus 2 h at
700  C) condition (CWA). Specimens were irradiated

at 450–735  C to a fast neutron fluence of
4 Â 1026 n mÀ2 (E > 0.1 MeV) in EBR-II, and were
tested at Ti, Ti þ 110  C and 232  C. Severe irradiation
embrittlement was evident in the ST alloys and STA
D25, particularly in tests at Ti þ 110  C. Zero ductility
was recorded in the lower-Ni alloy D21 (25Ni–8Cr)
irradiated and tested at 550 and 600  C. Severe ductility losses were associated with intergranular failures,
which were attributed to irradiation-induced solute
segregation and consequent precipitation of brittle g0
layers at grain boundaries. However, reasonable levels
of ductility, ranging from 2 to 6%, coupled with transgranular failures, were obtained at all temperatures in
irradiated CWA D66 (45Ni–12Cr). The preirradiation
grain boundary structure of this material, comprising
a ‘necklace’ of small recrystallized subgrains plus large
g0 particles and discrete Laves particles, remained
stable with no indication of irradiation-induced g0
layers. Yang et al. considered that the radiation-induced
segregation of g0 forming solutes to grain boundaries
was inhibited by the introduction of a high density of
dislocation sinks by cold working.
Vaidyanathan et al.90 and Huang and Fish91
examined the embrittlement of EBR-II-irradiated,
precipitation-hardened alloys, using ring ductility
tests. In this test, small sections of tubing are compressed and the ductility, defined as the strain at the
initiation of cracking, is deduced from the change
in the sample radius of curvature at maximum load.
Both experiments included Inconel 706 and Nimonic
PE16 in ST conditions, while Vaidyanathan et al.
also examined the developmental alloys D25 and
D68 in ST and STA conditions. Peak fluences in

these experiments were around 6–7 Â 1026 n mÀ2
(E > 0.1 MeV) and irradiation temperatures were in
the range 460–616  C. All the materials exhibited low
ductility failures at high test temperatures, particularly in tests at about Ti þ 110  C where ductilities
were generally below 0.1%, though Vaidyanathan
et al. found that postirradiation heat treatments (typically of 4 h at 785  C) produced a moderate recovery
in ductility. Based largely on observations reported
by Yang81 for irradiated ST PE16, Vaidyanathan et al.
and Huang and Fish considered that the irradiationinduced embrittlement of precipitation-hardened
alloys could generally be attributed to the formation
of brittle g0 layers at grain boundaries. However, the

arguments presented were far from conclusive –
microstructural examinations of the developmental
alloys which were reported by Vaidyanathan et al.
showed only weak indications of g0 precipitation in
D25 even within the grains, and evidence for g0 precipitation at grain boundaries in D68 was not found in
the low ductility tested samples but only in material
irradiated to a higher fluence. Yang81 examined the
microstructure of a low Si (0.01%) heat of ST PE16,
which was irradiated in EBR-II to doses of about 30
and 50 dpa at temperatures from 425 to 650  C. Grain
boundary g0 layers were observed in ST PE16 samples
which were irradiated at 510  C or above but not at
425  C, and helium bubbles were detected at boundaries in samples irradiated at 600–650  C. It was
considered by Yang that the irradiation-induced
embrittlement of ST PE16 was mainly attributable to
the cleavage fracture of grain boundary g0 layers and
that any effects of helium were of secondary importance. However, although grain boundary precipitation
of g0 was observed by Boothby28 in PE16 irradiated to

relatively high doses in EBR-II, there was no evidence
for the formation of intergranular g0 layers in the aged
conditions of PE16 which exhibited low ductility failures following irradiation in DFR to $20 dpa.89 Thus,
although it remains possible that the formation of grain
boundary g0 layers may aggravate the embrittlement,
it was considered by Boothby28 that the irradiation
embrittlement of PE16 is primarily due to helium.
A breach in solution-annealed Inconel 706 fuel
pin cladding, irradiated to 5% burn-up in EBR-II,
was reported by Yang and Makenas.92 The rupture
occurred from 12.7 to 18.4 cm from the bottom of the
pin, corresponding to irradiation at 447–526  C at a
fluence of 6 Â 1026 n mÀ2 (E > 0.1 MeV). Microstructural examinations revealed a brittle intergranular
fracture, with failure being attributed to a combination of matrix hardening due to g0 precipitation and
grain boundary weakening due to the formation of
interconnected Ni3(Ti,Nb) Z phase particles. In contrast to the work of Rowcliffe and Horak86 where
grain boundary Z phase was precipitated during a
preirradiation aging treatment, this phase formed
during the irradiation period in the solutionannealed cladding. Precipitation of Z was considered
to be irradiation enhanced because it was not formed
in long-term thermal annealing experiments at 480–
540  C. Grain boundary precipitation of Z phase was
also observed at the hot (650  C) end of the fuel pin
cladding, with both g0 and g00 in the matrix.
Cauvin et al.93 and Le Naour et al.94 also attributed
irradiation embrittlement effects in Inconel 706


Radiation Effects in Nickel-Based Alloys


cladding to the combined effects of matrix hardening
and the precipitation of Z at grain boundaries.
Inconel 706 fuel pin cladding, fabricated from four
heats with Nb contents varying from 1 to 3% and in
two heat-treated conditions (solution annealed or
aged), was irradiated in the Phenix fast reactor up
to a maximum of 100 dpa. Tensile tests on cladding sections were carried out at a strain rate of
3 Â 10À4 sÀ1. Tensile tests performed at ambient temperature showed high UTS (>1000 MPa) along the
full length of the pins with peak values of $1500 MPa
in sections irradiated near 500  C; ductility values
(uniform elongations only were given) remained
low (<2%) for irradiation temperatures up to
550  C and then increased sharply. In tests carried
out at the irradiation temperatures, however, premature failures occurred above about 500  C, with tensile strengths reduced to 300 MPa and ductilities
close to zero. All of the materials examined by Cauvin and Le Naour et al. showed similar properties,
with no systematic influence of composition or heat
treatment. Plitz et al.57 listed three fuel pin failures in
Inconel 706 cladding in the Phenix reactor, with a
further 12 failures in austenitic steel cladding; two of
the Inconel 706 failures were associated with long
periods of low power operation followed by a rise to
full power conditions, resulting in mechanical interaction between the fuel and the low ductility cladding
material.
4.04.5.2

Helium Implantation Experiments

Relatively few experiments have used helium implantation to investigate the embrittlement of nickel-based
alloys, although this technique has been more widely
used for austenitic steels. However, some data for

high-Ni alloys have been published by Shiraishi
et al.95 and Boothby.96
Shiraishi et al.95 compared the effects of helium
injection and neutron irradiation on the tensile properties of developmental g0 -Ni3(Ti,Al) (alloy 7817)
and g00 -Ni3Nb (alloy 7818) precipitation-hardened
40Ni–15Cr alloys. A number of alloy conditions,
including ST, aged, and cold worked, were tested.
Cyclotron injections of helium were made at 650  C
to levels of 5 or 10 appm. Neutron irradiations were
made at the same temperature to a fast fluence
(E > 1 MeV) of 1.7 Â 1024 n mÀ2 and a thermal fluence of 5.9 Â 1024 n mÀ2. The helium content of
the reactor-irradiated specimens was estimated to
be $45 appm, produced mainly from the thermal
neutron reaction with 10B. Tensile tests were carried

145

out at the implantation/irradiation temperature at a
strain rate of $5 Â 10À4 sÀ1. The results showed similar trends in helium-implanted and neutronirradiated specimens, with the total elongation values
tending to decrease with increasing tensile strength.
Variations in tensile strength for each alloy were
largely attributable to variations in the initial heat
treatment and working schedules. However, there
were some indications of softening and reduced
ductility in the neutron-irradiated specimens compared to those injected with helium. Overall, the g0 hardened alloy 7817 exhibited relatively high tensile
strength (typically >700 MPa) but low ductility following helium implantation or neutron irradiation
(with total elongation values generally <10% and
as low as 1–2% in the highest strength conditions).
The g00 -hardened alloy 7818 showed lower tensile
strength (typically 500–600 MPa) but maintained

good ductility, with total elongation values always
exceeding 10% and generally being above 20% in
STA conditions.
Boothby96 examined the effects of helium and/or
lithium injection on the tensile properties of STA
(ST 1050  C, aged 8 h at 700  C) Nimonic PE16.
The effects of lithium were examined since this element is also produced from the 10B(n,a)7Li reaction
in neutron-irradiated alloys. Helium and lithium
were implanted at ambient temperature, either singly
or in combination, to levels of 10 appm each. Samples
were tested in the as-implanted condition or following an additional aging treatment of 72 h at 750  C.
Tensile tests at a strain rate of 3 Â 10À5 sÀ1 on asimplanted samples showed no effect of helium or
lithium on ductility at 200–550  C. However, at
650  C, the as-implanted samples were all embrittled
to a similar extent, with the total elongations generally reduced to about half of the unimplanted levels
regardless of whether He, Li, or (He þ Li) ions were
injected. Postimplant aging of samples containing He
or (He þ Li) resulted in further ductility loss in tests
at 650  C, with significant embrittlement also evident
at 550  C though not at 450 or 200  C. Postimplant
aging of samples containing only lithium, however,
resulted in some recovery in ductility compared to
the as-implanted condition at 650  C but some ductility loss at 550  C. Thus, although it was clear that
lithium had a detrimental effect on ductility, it did
not appear to exacerbate the effects of helium.
A mechanism for lithium embrittlement was not
identified, though ductility loss in both Li and He
implanted samples was associated with an increased
propensity for intergranular fracture.



146

Radiation Effects in Nickel-Based Alloys

Some additional, previously unpublished, data
from a helium injection experiment conducted by
Boothby and Cattle are given in Table 2. In this
experiment, helium was implanted at ambient temperature to levels of 2, 10, and 50 appm into Nimonic
PE16 that had been given a two-stage aging treatment (ST 1050  C, aged 4 h at 800  C plus 16 h
750  C). As before, helium-doped samples were
either retained in the as-implanted condition or
given an additional aging treatment to coarsen the
dispersion of gas bubbles. Tensile tests were carried
out at 650  C at strain rates of 3 Â 10À5 and
3 Â 10À6 sÀ1. The results show a significant loss of
ductility even at 2 appm helium. In tests carried out
at the higher strain rate, the total elongation values
decreased progressively with increasing helium content and were further reduced by postimplant aging.
The ductility of as-implanted samples was generally
lower but less sensitive to helium concentration in
tests at a strain rate of 3 Â 10À6 than at 3 Â 10À5 sÀ1.
However, there was little effect of the strain rate on
the ductility of the postimplant aged samples.
Figure 15 illustrates grain boundary structures in
a tensile-tested PE16 sample which had been aged
subsequent to helium injection. Failure in this case
appeared to occur by the growth and coalescence of
cavities which were nucleated at grain boundary gas
bubbles.97 The nucleation of unstable cavities at grain

boundary helium bubbles requires the application of
a critical tensile stress, which is an inverse function of
the bubble radius, normal to the boundary. Cavity
growth then occurs via the stress-induced absorption
of vacancies. In the as-implanted condition, however,
Table 2

Tensile properties of helium-implanted Nimonic PE16 at 650  C

Strain rate (sÀ1)
À5

3 Â 10

3 Â 10À6

the helium dispersion was too fine to enable grain
boundary cavities to nucleate during tensile testing.
Reduced ductility in the as-implanted samples
appeared to be associated with grain boundary
wedge cracking, where, as discussed by van der
Schaaf and Marshall98 in relation to helium embrittlement of type 316 steel, the role of helium may be
simply to decrease the effective surface energy for
fracture.
Although it is evident from simulation experiments that helium alone can largely account for irradiation embrittlement, it is more difficult to assess
the significance of other radiation-induced effects such
as matrix hardening and grain boundary segregation
and/or precipitation. One experiment which examined
the effect of the radiation-induced precipitation of the
Ni3Si g0 phase on the ductility of a binary Ni-8 at.% Si

alloy was described by Packan et al.99 In this experiment, thin foil tensile specimens were bombarded with
either protons or a-particles to damage levels of 0.1–
0.3 dpa at 750 K; irradiation with a-particles resulted in
the introduction of high helium concentrations of
about 750 appm per 0.1 dpa. Proton and a-particle
irradiations both resulted in the formation of g0 layers
about 20–30 nm thick at the grain boundaries, but the
material remained relatively ductile, exhibiting transgranular failures, in tensile tests carried out at a strain
rate of $3 Â 10À4 sÀ1 at room temperature and, for the
proton-irradiated case only, 720 K. Unfortunately, no
tests were carried out at higher temperatures and samples which were irradiated with a-particles only at
750 K were not tested except at room temperature.
However, low ductility intergranular failure was

Helium (appm)

0.2% PS (MPa)

UTS (MPa)

Uniform elongation (%)

Total elongation (%)

0
2
10
50
2a
10a

50a
0
2
10
50
2a
10a
50a

487
495
432
505
445
443
403
434
473
430
404
425
404
430

575
603
538
569
500
499

483
491
491
479
458
438
431
439

3.9
5.2
4.9
3.7
4.2
2.8
2.2
2.3
0.8
1.2
1.1
0.6
2.3
0.4

35.4
19.3
11.7
7.1
5.4
4.4

2.4
31.8
6.8
7.2
6.0
5.8
4.3
1.3

a
Postimplant aged 72 h at 750  C.
Unpublished data from Boothby, R. M.; Cattle, G. C. Development of g0 -hardened 25Ni-based Alloys for Fast Reactor Core Applications;
FPSG/P(91)9, with permission from AEA Technology PLC.


Radiation Effects in Nickel-Based Alloys

100 nm

147

200 nm

Figure 15 Transmission electron micrographs illustrating (left) bubble dispersion on a grain boundary parallel to the
tensile axis and (right) cavitation on a boundary approximately normal to the tensile axis in Nimonic PE16 (implanted with
10 appm helium at ambient temperature, and subsequently aged for 60 h at 750  C prior to tensile testing at 650  C).
Reproduced from Boothby, R. M. J. Nucl. Mater. 1990, 171, 215–222.

induced in a test carried out at 720 K in a sample which
was preimplanted with 1000 appm He at 970 K, then

irradiated to 0.3 dpa, introducing an additional
2300 appm He at 750 K. Preimplantation of helium at
970 K produced grain boundary bubbles which were
10–20 nm in diameter, compared to 1.5–2 nm in material that was only irradiated with a-particles at 750 K.
The results of this experiment therefore indicated
that the radiation-induced precipitation of g0 at grain
boundaries did not give rise to embrittlement unless
helium was also implanted into the specimens.

4.04.6 Concluding Remarks
The effects of irradiation on the microstructure and
mechanical properties of nickel-based alloys are complex and, although the main factors affecting their
behavior have been identified, a full understanding of
radiation-induced effects remains elusive. This is particularly true of the precipitation-hardened alloys,
typified by Nimonic PE16 and Inconel 706, where the
role of the hardening phases – which confer high thermal creep strength, but are redistributed during irradiation and may possibly influence swelling behavior and
contribute to intergranular embrittlement – is unclear.
The radiation-induced effects considered in this
chapter – void swelling, irradiation creep, the evolution
of precipitate and dislocation structures, and irradiation
embrittlement – are interrelated in several ways, but
particularly through the effects of point defect fluxes
and the consequent redistribution of solute atoms.
The beneficial effect of nickel on the swelling
resistance of austenitic alloys is well known, but a

clear explanation for the minimum in swelling that
is found in alloys containing about 40–45% Ni has
not been forthcoming. There is general agreement
that the major influence of alloy composition on

swelling arises through its effects on the effective
vacancy diffusivity and on segregation via the
inverse Kirkendall effect. However, on what appears
to be the mistaken assumption that the swelling resistance of nickel-based alloys derives from an extended
void nucleation period, swelling models have largely
focused on factors affecting the nucleation rather than
the growth of voids. Data for neutron-irradiated Nimonic PE16, for example, indicate that its swelling resistance is due to a combination of a comparatively low
saturation void concentration, which is reached at a
relatively low displacement dose, and a low void
growth rate. The minimum critical void radius concept
appears to offer the most plausible explanation for the
minimum in swelling found at intermediate nickel
contents, although experimental data comparing the
behavior of PE16 and a nonprecipitation hardenable
alloy with a similar matrix composition indicate that, in
addition to the Ni content of the alloy, the presence of
Si and/or the g0 forming solutes Al plus Ti may also be
important. The dependence of the void growth rate on
Ni may be related to the effects of radiation-induced
segregation on the bias terms for the absorption of
point defects at sinks, though again there is evidence
that minor solutes, including Si, B, and Mo, as well as
the g0 -forming elements, have a beneficial effect on the
overall swelling behavior of nickel-based alloys.
The irradiation creep behavior of nickel-based
alloys generally appears to be similar to that of
austenitic steels, though the higher thermal creep