Comprehensive nuclear materials 1 04 effect of radiation on strength and ductility of metals and alloys

1.04 Effect of Radiation on Strength and Ductility of
Metals and Alloys
M. L. Grossbeck
University of Tennessee, Knoxville, TN, USA

ß 2012 Elsevier Ltd. All rights reserved.

1.04.1
1.04.2
1.04.3
1.04.4
1.04.5
1.04.6
1.04.7
1.04.7.1
1.04.7.2
1.04.7.3
1.04.8
1.04.8.1
1.04.8.2
1.04.9
1.04.10
References

Introduction
Mechanisms of Irradiation Hardening
Tensile Behavior
Effects of Neutron Spectrum
Tensile Ductility
Effect of Test Temperature
Ferritic–Martensitic Alloys

Introduction
Tensile Behavior
Helium Effects
Refractory Metals
Tensile Behavior
Helium Effects
Amorphous Metals
Conclusions

Abbreviations
A1
ANN
appm
ASTM
ATR
bcc
BR2
CW
DBTT
dpa
EBR-II
fcc
FFTF
HFBR
HFIR
HFR
JPCA

Lowest equilibrium temperature at which
the austenite phase exists in steel

Annealed
Atomic parts per million
ASTM International
Advanced Test Reactor, Idaho Falls, ID,
USA
Body-centered cubic
Belgian Reactor-2, Mol, Belgium
Cold worked
Ductile-brittle transition temperature
Displacements per atom
Experimental Breeder Reactor-II, Idaho
Falls, ID, USA
Face-centered cubic
Fast Flux Test Facility, Richland, WA,
USA
High Flux Beam Reactor, Brookhaven,
Upton, NY, USA
High Flux Isotope Reactor, Oak Ridge,
TN, USA
High Flux Reactor, Petten, The
Netherlands
Japanese Prime Candidate Alloy

99
100
101
102
103
107
108

108
108
113
116
116
118
119
120
121

LMFBR Liquid Metal Fast Breeder Reactor
LWR
Light Water Reactor
ORR
Oak Ridge Research Reactor, Oak Ridge,
TN, USA
PCA
Prime candidate alloy, adopted by the US
Fusion Program in mid-1970s
ppm
Parts per million
Unirr
Unirradiated

1.04.1 Introduction
The most commonly considered mechanical properties of metals and alloys include strength, ductility,
fatigue, fatigue crack growth, thermal and irradiation
creep, and fracture toughness. All these properties are
important in the design of a structure that is to experience an irradiation environment. While determining
the mechanical properties of irradiated materials, tensile properties, typically yield strength, ultimate tensile

strength, uniform elongation, total elongation, and
reduction of area are the most commonly considered
because they are usually the simplest and the least
costly to measure. In addition, the tensile properties
can be used as an indicator of the other mechanical

99


100

Effect of Radiation on Strength and Ductility of Metals and Alloys

properties. Space in a reactor or in an accelerator target
is often so limited that the larger specimens required
for fatigue and fracture toughness testing are not practical; consequently, the number of specimens that can
be irradiated is so small that a meaningful test matrix
is not possible. Shear punch testing of 3-mm diameter disks, typically used as transmission microscopy
specimens, was developed to address the problem of
irradiation space. Although much information can be
obtained from shear punch testing, the tensile test
remains the most reliable indicator of strength and
ductility. For these reasons, the tensile test is usually
the first mechanical test used in determining the
irradiated properties of new materials. This chapter
addresses the tensile strength and ductility of alloys.

1.04.2 Mechanisms of Irradiation
Hardening
Irradiation introduces obstacles to dislocation motion,

which results in plastic deformation, in the form of
defects resulting from atomic displacement and from
transmutation products. Small Frank loops and defect
clusters, known as black dots, large Frank loops (about
an order of magnitude larger), precipitates, and cavities
(either voids or bubbles) contribute to hardening in an
irradiated alloy. Frank loops unfault and eventually
contribute to the network dislocation density. Precipitates are certainly present in the unirradiated alloy,
but additional precipitation results from the segregation of elements during irradiation and from the
irradiation-induced changes that shift the thermodynamic stability of phases. Transmutation production of
new elements in the alloy can also result in the formation of new precipitates. The production of insoluble
species, most importantly helium, also results in precipitation, especially in the form of bubbles.
Defects are divided into two classes: long range
and short range. Short-range obstacles are defined as
those that influence moving dislocations only on the
same slip plane as opposed to long-range obstacles,
which impede dislocation motion on slip planes not
containing the obstacle.1 Coherent precipitates and
large loops are long-range obstacles, but for this
analysis, only network dislocations will be considered
as long-range obstacles, a reasonable simplification
from observations. As recommended by Bement,2 the
contributions from short-range obstacles are added
directly,
DFTS ¼ DFLR þ DFSR

½1Š

where the quantities in eqn [1] are total stress, longrange contribution to stress, and short-range contribution to stress. The contributions from the short-range
obstacles are added in quadrature as follows3:

ðDFSR Þ2 ¼ ðDFSMloop Þ2 þ ðDFLGLoop Þ2
þ ðDFPRECIP Þ2 þ ðDFCAVITY Þ2

½2Š

where the term on the left represents the contribution
from all short-range obstacles, and the terms on the
right represent the stress contributions from small
loops, large loops, precipitates, and cavities, either
voids or bubbles.
The contribution to hardening by network dislocations may be expressed by
pffiffiffiffiffi
½3Š
tnet ¼ aGb rd
where tnet is the increment in shear stress, G is the
shear modulus, b is the Burgers vector, and rd is
the dislocation density. The constant a is dependent
upon the geometry of the dislocation configuration
and is usually determined experimentally. However,
Taylor has calculated a to be between 0.15 and 0.3,4
and Seeger has determined the value to be 0.2, incorporating the assumption of a random distribution of
dislocation directions.5 Short-range defects such as
small and large Frank loops and precipitates are
treated as hard impenetrable obstacles where dislocations bow around them by the Orowan mechanism.
The stress increment is expressed by
pffiffiffiffiffiffiffiffiffiffiffi
Dt ¼ Gb Nd =b
½4Š
where N is the defect density and d is the diameter.
The constant b ranges between 2 and 4 as suggested

by Bement2 or 6 as suggested by Olander.6 Voids and
bubbles are also treated as hard obstacles using the
same expression. Precipitates and bubbles have been
observed in austenitic stainless steels to nucleate and
grow together.7 In this case, the bubbles and precipitates are considered as one obstacle where the hardening increment is calculated assuming rod geometry
using a treatment by Kelly expressed by8:
pffiffiffiffiffiffi pffiffiffi 
6d
0:16Gb Nd
pffiffipffiffiffiffiffiffi ln
Bubble-precip ¼
½5Š
6
3b
1À
Nd
3

where the parameters are the same as for eqn [4].
From the previous discussion, it can be inferred that
because the nature of the irradiation-induced defects
determines the degree of hardening, and because the
nature, size, and density of defects is a strong function of
temperature, radiation strengthening will be a strong
function of irradiation temperature. Figure 1 illustrates


Effect of Radiation on Strength and Ductility of Metals and Alloys

101

Relative contribution
to strength

1
Black dots

0.8

Frank loops

0.6

Bubbles-precipitates
Network dislocation

0.4
0.2
0
0

100

200
Temperature (ЊC)

300

400

Figure 1 Relative contribution to strengthening from irradiation-induced defects in the austenitic stainless steel, PCA,
irradiated to 7 dpa in the Oak Ridge Research Reactor. Reproduced from Grossbeck, M. L.; Maziasz, P. J.; Rowcliffe, A. F.
J. Nucl. Mater. 1992, 191–194, 808.

160

Irrad. temp. » Test Temp.
Strain rate ~ 4 ´ 10-5 S-1

1000

140

900

371 ЊC

Yield strength (MPa)

800

Test
Symbol Temp. (ЊC)

427 ЊC

700

371
427

483
538
593
649
704
760
816

483 ЊC

600
500
538 ЊC
593 ЊC

400
300
760 ЊC

100

816 ЊC

0

100
80
60

649 ЊC

200

0

1

2

120

Yield strength (ksi)

1100

40

704 ЊC
20

3

4

5

6

7

8

9

10

12

0

Neutron fluence (n cm−2 ) (E > 0.1 MeV)

Figure 2 Yield strength of 20% cold-worked type 316 stainless steel irradiated in the EBR-II. Reproduced from
Fish, R. L.; Cannon, N. S.; Wire, G. L. In Effects of Radiation on Structural Materials; Sprague, J. A., Dramer, K., Eds.;
ASTM: Philadelphia, PA, 1979; ASTM STP 683, p 450. Reprinted, with permission, from Effects of Radiation on Structural
Materials, copyright ASTM International, West Conshohocken, PA.

strengthening from individual types of defects as a
function of irradiation temperature for the austenitic
stainless steel PCA.7
As can be seen from Figure 1, the black dot
damage characteristic of low temperatures vanishes
at temperatures over 300  C as Frank loops emerge.
Bubbles and precipitates also become major contributors to hardening above 200  C.

1.04.3 Tensile Behavior
Tensile behavior is determined by the irradiationinduced defect structure previously discussed. Austenitic stainless steels will again be used for the
example since they are typical of fcc alloys and in

many respects to other alloys (see Chapter 2.09,

Properties of Austenitic Steels for Nuclear Reactor Applications and Chapter 4.02, Radiation
Damage in Austenitic Steels). The behavior of
other example classes of alloys will be discussed in
later sections of this chapter. The tensile behavior
characteristic of austenitic stainless steels is shown in
Figure 2, where yield strength is plotted as a function
of fluence and displacement level.9 Saturation in
strength is clear with the saturation time becoming
shorter as irradiation temperature is increased. At
temperatures above about 500  C, saturation is evident, but in this case, strength decreases.
This decrease is a result of recovery of the coldworked microstructure of the 20% cold-worked type
316 stainless steel presented in Figure 2. Figure 3


102

Effect of Radiation on Strength and Ductility of Metals and Alloys

850

600

Ultimate tensile strength (UTS)

650 ЊC

20% cold-worked

400

800
750

200
Strength (MPa)

Annealed
0
538 ЊC

Yield strength (MPa)

600

20% cold-worked
400

Yield strength (YS)

650
600
550

200

450

427 ЊC
800
600

Annealed

200
0

1

2

3

4

5

6

7

8

0

10

20
30
Dose (dpa)

40

50

Figure 4 Strength properties of 20% cold-worked type
316 stainless steel irradiated in EBR-II. Reproduced from
Allen, T. R.; Tsai, H.; Cole, J. I.; Ohta, J.; Dohi, K.;
Kusanagi, H. Effects of Radiation on Materials; ASTM:
Philadelphia, PA, 2004; ASTM STP 1447, p 3. Reprinted,
with permission, from Effects of Radiation on Structural
Materials, copyright ASTM International, West
Conshohocken, PA.

20% cold-worked

400

20% CW 316 stainless steel
irradiation temp 375–385 ЊC
Test temperature 370 ЊC

500

Annealed

0

0

700

9

10

Neutron fluence (1022 n cm−2 )
Figure 3 Yield strength of type 316 stainless steel
irradiated in the EBR-II. Reproduced from Garner, F. A.;
Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl.
Mater. 1981, 103 & 104, 803.

shows yield strength resulting from the recovery of a
cold-worked dislocation structure and the generation
of a radiation-induced microstructure, resulting in a
saturation strength independent of the initial condition of the alloy.10 Again, it is seen that the approach
to saturation is faster with increasing temperature,
with saturation achieved between 5 and 10 dpa at 538
and 650  C, but 15–20 dpa is necessary to achieve
saturation at 427  C.
Saturation is observed in yield strength curves for
fluences as high as 9 Â 1022 n cmÀ2 in a fast reactor
(45 dpa), but more recent data show a hint of softening above 50 dpa,11,12 and other fast reactor data
have shown a reduction in strength even for displacements below 50 dpa, as shown in Figure 4.13
This could result from coarsening of the microstructure or depletion of interstitial elements from the
matrix due to precipitation. This effect is also
observed in martensitic steels irradiated to high dpa
levels in the FFTF, but in this class of alloys, recovery
of the martensitic lath structure is also a factor.12
However, even in austenitic steels, it is difficult
to attribute such softening with certainty to an
irradiation effect because of the strong influence

of irradiation temperature on strength.14 Indeed,
uncertainties in irradiation temperature are an inherent difficulty in neutron irradiation experiments.

1.04.4 Effects of Neutron Spectrum
This discussion has used neutron irradiations for
illustration purposes. Reactors provide an effective
instrument for achieving high neutron exposures
under conditions relevant for most nuclear applications. However, reactor irradiations suffer from many
difficult-to-control and, sometimes, uncontrolled
variables. The neutron energy spectrum is responsible for large differences in irradiation effects between
different reactors.
The mechanism of atomic displacement is well
understood.15 With a known neutron energy spectrum, neutron atomic displacements can be calculated as a function of fluence for a given reactor.
Transmutation of elements in the material under
study, which is a strong function of neutron spectrum,
results in wide variation in some mechanical properties. This is of particular importance in applying
fission reactor results to fusion. In a fusion device,
helium and hydrogen will be generated through (n,a)
and (n,p) reactions in nearly all common structural
materials. Hydrogen has a very high diffusivity in
metals so that an equilibrium concentration will be


Effect of Radiation on Strength and Ductility of Metals and Alloys

established at a level that is believed to be benign.16
By contrast, helium is insoluble in metals, segregating
at grain boundaries and other internal surfaces and
discontinuities.

Although helium is produced in all nuclear reactors, the thermal spectrum is responsible for the
highest concentrations. The largest contributors to
helium in a thermal reactor are boron and nickel by
the following reactions:
10

58

Bðn; aÞ7 Li

Niðn; gÞ59 Ni

59

Niðn; aÞ56 Fe

Boron is present as a trace element in most alloying elements but only at ppm levels. Nickel is a
major constituent of many alloys and a minor constituent of still others. The two nickel reactions
constitute a two-step generation process for
helium, which starts slowly and accelerates as
59
Ni builds up in the alloy, limited only by the
supply of 58Ni, which for practical purposes is
often unlimited. In austenitic alloys, the high flux
isotope reactor (HFIR) has generated over
4000 appm He in austenitic stainless steels. The
generation rate is so high that multistep absorber
experiments have been conducted to reduce the
helium generation rate to that characteristic of
fusion reactors, 12 appm He per dpa in austenitic

stainless steels.17 (see Chapter 1.06, The Effects
of Helium in Irradiated Structural Alloys).
Other transmutation products may also complicate reactor irradiation studies. Examples are the
transmutation of manganese to iron by the following reaction: 55Mn (n,g) 56Mn ! 56Fe and the transmutation of chromium to vanadium by 50Cr (n,g)
51
Cr ! 51V. The first reaction leads to loss of an
alloy constituent, and the second leads to doping
with an extraneous element. However, neither of
these reactions has been shown to significantly affect
mechanical properties of steels.18
Helium remains the most studied transmutation
product, and it can have profound effects on tensile
properties, especially at high temperatures. Experiments have been conducted in various reactors
throughout the world to assess the effects of helium
on mechanical properties of alloys.19 An interesting
result is that helium has little effect on strength.
This is illustrated in Figure 5 where a comparison
has been made between austenitic steels irradiated
in Rapsodie, a fast spectrum reactor, and steels
irradiated in HFIR, a mixed-spectrum reactor with

103

a very high thermal flux. The saturation yield
strength of all alloys remains within a single scatter
band.20,21
The tramp impurity elements sulfur and phosphorus have significantly high (n,a) cross sections at
high energies, as shown in Figure 6. Although the
cross section for phosphorus is large only at energies
characteristic of fusion, a boiling water reactor produces 500 appm He from sulfur and 40 appm He from

phosphorus in eight years of operation. An Liquid
Metal Fast Breeder Reactor (LMFBR) can produce
100 times these concentrations. All these elements
are expected to enhance embrittlement when segregated to grain boundaries, but it remains to be
determined which is more detrimental, helium, sulfur, or phosphorus.

1.04.5 Tensile Ductility
Tensile ductility is a more vulnerable parameter than
strength to radiation effects since it tends to be very
high in unirradiated austenitic stainless steels and is
often reduced to quite low levels by irradiation. It is
also of more concern since strengthening, although
not reliable due to its slow initiation, is usually a
beneficial change. In contrast, embrittlement is
always detrimental. Like strength, ductility exhibits
saturation with increasing fluence, although the
behavior is significantly more complex than that of
strength. The general trends in type 316 stainless
steel are shown in Figure 7 for material irradiated
in the EBR-II. These data are for the same specimens
for which the yield strength was shown in Figure 2.9
Fast reactor data are used here to avoid the complication of helium effects. Once stabilization of the
dislocation microstructure is achieved, a smooth
curve approaching an apparent saturation is observed.
More information can be gleaned from ductility
data if they are viewed in terms of irradiation and
test temperature. Figure 822 shows total tensile elongation for a series of irradiated austenitic alloys at
a displacement level of 30 dpa in both annealed
and cold-worked conditions. The room temperature
ductility exceeds 10%, but it decreases rapidly with

increasing temperature up to approximately 300  C
and then exhibits the expected increase with temperature observed for unirradiated alloys. Beyond
500  C, ductility again decreases with an onset of
intergranular embrittlement resulting from helium
introduced through transmutations in the thermal
flux of the HFIR.


104

Effect of Radiation on Strength and Ductility of Metals and Alloys

900
1.4988 sa
1.4970 sa + cw + a
AISI 316 cw
AISI 304 sa

800

Saturation yield strength (MPa)

700

600

500

400

US 316 20% cw HFIR
JPCA SA HFIR
JPCA 15% cw HFIR

300

316 20% cw EBR-II
US PCA SA + 800 ЊC, 8 h HFIR

200
Typical yield strength values of unirradiated
solution annealed austenitic stainless steel
100
350

400

450

500

550

600

650

Irradiation and test temperature (ЊC)
Figure 5 Saturation yield strength as a function of temperature for austenitic alloys irradiated in Rapsodie, EBR-II, and
high flux isotope reactor showing similar saturation strength. Reproduced from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C.

J. Nucl. Mater. 1990, 174, 264.

1.00E + 03

(n, a) cross-section (barns)

1.00E + 02

LMFBR flux (arb. units)
LWR flux (arb. units)
Sulfur (n, a)
Phosphorus (n, a)

1.00E + 01

1.00E + 00

1.00E - 01

1.00E - 02

1.00E - 03
1.E - 11

1.E - 08

1.E - 05
1.E - 02
Energy (MeV)

Figure 6 Cross-section for (n,a) reactions as a function of neutron energy.

1.E + 01

1.E + 04


Effect of Radiation on Strength and Ductility of Metals and Alloys

105

17

Irrad. temp. » Test temp.
Strain rate ~ 4 ´ 10-5 S−1

16
15
14

Test
Temp. (ЊC)
371
427
483
538
593
649
704
760

816

Symbol

13
Total elongation (%)

12
11
10
9
8

593 ЊC

7

538 ЊC

6
5
4
3

816 ЊC

2
1
0

760 ЊC

0

1

2

649 ЊC
371 ЊC

704 ЊC

3
4
5
6
7
8
Neutron fluence (n cm−2) (E > 0.1 MeV)

10 ´ 1022

9

Figure 7 Total elongation of 20% cold-worked type 316 stainless steel irradiated in EBR-II. Reproduced from
Fish, R. L.; Cannon, N. S.; Wire, G. L. In Effects of Radiation on Structural Materials; Sprague, J. A., Dramer, K., Eds.;
ASTM: Philadelphia, PA, 1979; ASTM STP 683, p 450. Reprinted, with permission, from Effects of Radiation on Structural
Materials, copyright ASTM International, West Conshohocken, PA.

ORNL-DWG 89-13395
20
30 DPA

+
+
J

+
Total elongation (%)

15

*

0
C

US PCA 25% CW HFIR
JPCA ANN HFIR
JPCA 15% CW HFIR
US 316 20% CW HFIR
US PCA B3 HFIR
316 20% CW EBR-II

J
J

10

+
J
J*
0
0

+

5

J
J
0
C.
0

0

100

200

J
C.

300
400
Temperature (ЊC)

500

C.J
C.
00
600

700

Figure 8 Total elongation as a function of irradiation and test temperature for fast (EBR-II) and mixed-spectrum
(high flux isotope reactor) reactor irradiation.

Uniform elongation, the elongation at the onset of
plastic instability, or necking, appears to be most sensitive to the effects of irradiation and, in general, is
less dependent on specimen geometry than other parameters such as total tensile elongation. The low values
of uniform elongation are often cause for great concern, which is usually justified. However, it should be
borne in mind that if stresses remain below the yield

stress of a metal, elongation becomes a secondary concern. As long as limited plastic deformation relieves the
stress that produced it, a structure remains intact.
The high level of irradiation strengthening
observed at temperatures below 300  C, which is
due to black dot defect clusters and small loops,
also results in low ductility throughout this temperature range. Small helium bubbles and helium-defect


106

Effect of Radiation on Strength and Ductility of Metals and Alloys

20

Uniform elongation (%)

10 dpa
250 ЊC
Sym. (%)

15

0.20
0.20
0.33
0.29
0.31
0.28
0.23
0.18
0.30
0.20
0.10

10

5

316 SS

PCA

0

0

100

200

300
400
Temperature (ЊC)
Symbols

500

600

316 20% CW EBR-II

316 ANN EBR-II

US 316 20% CW ORR

US PCA 25% CW HFIR

US PCA 25% CW ORR

US 316 20% CW HFIR

US PCA 25% CW HFR

US 316 20% CW HFR

316 20% CW DO HFIR

US PCA 25% CW BR2

US 316 20% CW BR2

EC 316 ANN HFIR

JPCA ANN HFIR

EC 316 ANN BR2

JPCA ANN HFR

EC 316 ANN HFR

JPCA 15% CW HFIR

US PCA B3 HFIR

700

Figure 9 Uniform elongation as a function of irradiation and test temperature at a displacement level of 10 dpa.
The trend curves are for type 316 stainless steel and PCA. Reproduced from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C.
J. Nucl. Mater. 1990, 174, 264.

clusters also contribute to hardening and reduction
in ductility, but this form of helium embrittlement
is not related to the severe intergranular embrittlement that is observed above 500  C. Both these

effects are apparent in Figure 9 where uniform
elongation for an extensive set of austenitic alloys
irradiated in thermal and fast spectrum reactors is
shown.11 The specimens irradiated in the fast spectrum (<5 appm He) exhibit consistently higher ductility than the mixed-spectrum reactor specimens
(500–1000 appm He) even at this low displacement
level, especially above 600  C, where helium embrittlement is certain to control.
A similar pattern is exhibited at 30 dpa where a
very limited uniform elongation characteristic of
lower temperatures is apparent. After a restoration
of ductility above 400  C, ductility again decreases
above 500  C due to the onset of intergranular
helium embrittlement. Differences in alloy behavior,
especially in the case of titanium-modified alloys
somewhat clouds the understanding of helium

embrittlement observed in Figure 10.11 However, at
50 dpa, where helium levels exceed 4000 appm, the
trend becomes clear with the fast reactor specimens
showing uniform elongations several times larger
than those observed in mixed-spectrum reactors
(Figure 11).11 What is less expected is the recovery
of ductility at 50  C at 50 dpa compared to the
results at 30 dpa. This irradiation annealing effect
has also been observed at 230  C by Ehrlich, where
strength of the alloy 1.4988 decreased continuously
from 10 to 30 dpa.20 Results from an experiment in
the Oak Ridge Research Reactor (ORR), where the
spectrum was tailored to produce a ratio of He per
dpa characteristic of a fusion reactor, show similar
low levels of uniform elongation for cold-worked

alloys at low temperatures, but high uniform elongations were observed in annealed type 316 stainless
steel at 60  C. This high ductility was drastically
reduced between 200 and 330  C before the microstructure characteristic of higher temperatures became
effective.21


Effect of Radiation on Strength and Ductility of Metals and Alloys

107

Uniform elongation (%)

20
316 20% CW EBR-II
US 316 20% CW ORR
US PCA 25% CW ORR
US PCA 25% CW HFR
316 20% CW DO HFIR
US 316 20% CW BR2
JPCA ANN HFIR
JPCA ANN HFR
JPCA 15% CW HFIR

15

10

316 ANN EBR-II
US PCA 25% CW HFIR
US 316 20% CW HFIR

US 316 20% CW HFR
US PCA 25% CW BR2
EC 316 ANN HFIR
EC 316 ANN BR2
EC 316 ANN HFR
US PCA B3 HFIR

30 dpa

5

0

0

100

200

300
400
Temperature (ЊC)

500

600

700

Figure 10 Uniform elongation of austenitic stainless steels irradiated in fast and thermal reactors to a displacement level

of 30 dpa. Severe helium embrittlement is shown at 600  C. Reproduced from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C.
J. Nucl. Mater. 1990, 174, 264.

ORNL-DWG 90-14298

20

Uniform elongation (%)

50 dpa
US PCA 25% CW HFIR
JPCA SA HFIR
JPCA 15% CW HFIR
316 20% CW EBR-II
US PCA SA + 800 ЊC, 8 h HFIR
J 316 20% CW HFIR
J 316 SA HFIR

15

10
64 dpa

5

0

78 dpa

0

100

200

300
400
Temperature (ЊC)

500

600

700

Figure 11 Uniform elongation of austenitic stainless steels irradiated to 50 dpa in high flux isotope reactor (HFIR) and
78 dpa in EBR-II showing embrittlement from helium generated in the mixed-spectrum reactor, HFIR. Reproduced
from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C. J. Nucl. Mater. 1990, 174, 264.

1.04.6 Effect of Test Temperature
An interesting phenomenon is observed when irradiated alloys are tested at temperatures different
from the irradiation temperature. Figure 12 shows
total elongation data from cold-worked type 316
stainless steel irradiated to displacement levels of
48–63 dpa in the FFTF, where elongation is plotted

against the increment of the test temperature above
the irradiation temperature.23 Although there is
significant scatter in the data, the elongations below
1% obtained by test temperatures about 100  C

above the irradiation temperature are cause for
concern. This phenomenon has also been observed
in higher nickel alloys. The cause of this phenomenon remains elusive, pending further testing with


108

Effect of Radiation on Strength and Ductility of Metals and Alloys

32
Fluence = 9.13 ´ 1022 n cm–2
(E > 0.1 MeV)

Total elongation (%)

28
24
20
16
12
8

f t = 12 ´ 1022 n cm–2

4
0
- 700 - 600 - 500 - 400 - 300 - 200 - 100

0

100

200

300

400 500

Test
Irradiation (ЊC)
–
temperature temperature
Figure 12 Total elongation of 20% cold-worked type 316 stainless steel irradiated in FFTF to displacement levels of
48–63 dpa. Hamilton, M. L.; Cannon, N. S.; Johnson, G. D. In Effects of Radiation on Materials; Brager, H. R., Perrin, J. S.,
Eds.; ASTM: Philadelphia, PA, 1982; ASTM STP 782, p 636. Reprinted, with permission, from Effects of Radiation on
Structural Materials, copyright ASTM International, West Conshohocken, PA.

different holding times at various temperatures
before tensile testing. Migration of interstitial solutes
to moving dislocations is a candidate mechanism for
this phenomenon.

1.04.7 Ferritic–Martensitic Alloys
1.04.7.1

Introduction

The class of ferritic–martensitic alloys with chromium concentrations in the range of 9–12% has
attracted interest in the fast reactor programs because
of its radiation resistance, in particular, very low

swelling and low irradiation creep. Alloys such as
Sandvik HT-9 (12Cr1Mo.6Mn.1Si.5W.3V)) and
other alloys of this class were irradiated in the
EBR-II,24 in research reactors25 and with heavy
ions.26 The quantitative results from the ion irradiations in this class of alloys and the low neutron
absorption cross section led to inclusion of ferritic
alloys into the fast reactor alloy development programs, in particular in the United States in the
mid-1970s. The radiation resistance has been confirmed to displacement levels of 70 dpa.12,14
Further interest in this class of alloys was initiated
by the fusion reactor programs in Europe and the
United States when the necessity for low neutron
activation structural materials was realized. Further

research on martensitic alloys by fusion programs in
Europe, the United States, and Japan led to the development of low-activation alloys by replacing elements
that result in long-term activation products. Molybdenum and niobium, both of which result in longlived activation products, were replaced by tungsten
and tantalum. This research led to radiation-resistant
alloys with a fracture toughness superior to that of the
commercial alloys even in the unirradiated condition.27 The compositions of representative members
of this class of alloys referred to in this chapter are
presented in Table 1. An excellent review of irradiation behavior of this class of alloys has been published
by Klueh and Harries.27 Details of the metallurgy of
martensitic alloys appears in Chapter 4.03, Ferritic
Steels and Advanced Ferritic–Martensitic Steels.
1.04.7.2

Tensile Behavior

Unlike the tensile behavior of fcc metals, where there
is a smooth increase in strength as plastic deformation proceeds and work hardening progresses, bcc

metals typically exhibit a load drop almost immediately following the onset of plastic deformation.
Interstitial solutes such as carbon in steels effectively
lock dislocations leading to a longer period of elastic
deformation after which generation of new dislocations results in a load drop, or yield point, until


Effect of Radiation on Strength and Ductility of Metals and Alloys

Table 1

109

Nominal or typical compositions of ferritic–martensitic alloys cited

Steel type

12Cr–MoVW
8Cr–2WVTa
9Cr–1WVTa
9Cr–1GeV
12Cr–1MoVNiNb
10Cr–MoVNiNb
9Cr–2WVTa
9Cr–2WVTa
9Cr–2WVTa
7Cr–2WVTa
2.25Cr–2WVTa
12Cr–2WVTa
12Cr–2WVTa
9Cr–1MoVNb

9Cr–2Mo–1Ni
9Cr–2W
9Cr–2W

Designation

HT-9
F82H
OPTIFER Ia
OPTIFER II
MANET I
MANET II
9Cr–2WVTa
JLF-1
JLF-2
JLF-3
JLF-4
JLF-5
JLF-6
T91
JFMS
NFL-0
NFL-1

Composition (wt%)
C

Si

Mn

Cr

Ni

Mo

V

0.20
0.1
0.1
0.1
0.14
0.1
0.1
0.1
0.1
0.09
0.1
0.09
0.10
0.1
0.05
0.10
0.10

0.38

0.60

11.95
8.0
9.0

0.60

1.0

10.8
10.0
8.7
9.0
9.16
7.03
2.23
11.99
12.00
9.0
9.6
8.65
9.01

0.9
0.7

0.30
0.2
0.25
0.3

0.2
0.2
0.23
0.19
0.20
0.20
0.20
0.19
0.19
0.2
0.12
0.25
0.26

0.2

0.3
0.67
0.056
0.042

0.4
0.46
0.45
0.45
0.50
0.48
0.46
0.4
0.58

0.050
0.53

0.94

0.75
0.6

1.0
2.3

v

Nb

W

B

Other

0.52
1.0
0.16
0.15

0.009
0.007
2.2
1.97

1.93
1.97
1.97
1.98
1.94

0.04 Ta
0.07 Ta
1.1Ge
0.06 Zr
0.03 Zr
0.06 Ta
0.07 Ta
0.07 Ta
0.07 Ta
0.07 Ta
0.07 Ta
0.07 Ta

0.08
0.06
1.92
2.06

0.0032

Source: Maloy, S. A.; Toloczko, M. B.; McClellan, K. J.; et al. J. Nucl. Mater. 2006, 356, 62; Klueh, R. L.; Harries, D. R. High-Chromium
Ferritic and Martensitic Steels for Nuclear Applications; ASTM: Philadelphia, PA, 2001; Kohno, Y.; Kohyama, A.; Hirose, T.;
Hamilton, M. L.; Narui, M. J. Nucl. Mater. 1999, 271 & 272, 145; Kurishita, H.; Kayano, H.; Narui, M.; Kimura, A.; Hamilton, M. L.;
Gelles, D. S. J. Nucl. Mater. 1994, 212–215, 730.

(´ 103)
120
1 ´ 1020 nvt

100

5 ´ 1018 nvt

Stress (psi)

80

Unirradiated
1.7 ´ 1019 nvt

60
40
20
0

0

0.05

0.10

0.15
0.20
0.25

Strain (in. in.–1)

0.30

0.35

0.40

Figure 13 Stress–strain curves for low-carbon steel weld material irradiated at 80  C (nvt is fluence in neutrons/cm2).
Reproduced from Wilson, J. C. Effects of irradiation on the structural materials in nuclear power reactors. In Proceedings
of the Second United Nations International Conference on the Peaceful Uses of Atomic Energy, United Nations, 1958;
Vol. 5, p 431.

terminated by the work hardening mechanism of
dislocation interaction.28
Upon irradiation, the load drop is frequently
masked by an early termination of work hardening,
leading to very low values of uniform elongation.
This behavior is evident even at displacement levels

below 0.01 dpa and is illustrated in Figure 13 from
research presented at the Second Atoms for Peace
Conference in 1958.29 Extreme irradiation hardening
and severe plastic instability are clearly illustrated by
this early research. More recent alloys with more
careful control of impurities and controlled processing


110

Effect of Radiation on Strength and Ductility of Metals and Alloys

1400

JFMS specimens tested at 25 ЊC after irradiation in FFTF

Engineering stress (MPa)

1200

35.3 dpa, Tirr = 390 ЊC
9.8 dpa, Tirr = 373 ЊC
22.2 dpa, Tirr = 390 ЊC
44 dpa, Tirr = 427 ЊC

1000
800

0 dpa

600
400
200
0
0

2

4

12
14
16
8
10
Engineering strain (%)

6

18

20

22

Figure 14 Stress–strain curves for JFMS alloy irradiated in FFTF to 44 dpa at temperatures of 373–427  C and tested at
25  C. Reproduced from Maloy, S. A.; Toloczko, M. B.; McClellan, K. J.; et al. J. Nucl. Mater. 2006, 356, 62.

800
MATRON/tensile(S)
700

656–700 K

Yield stress (MPa)

600
500
400
300

200
100
0

0

10

20

30

40

50

60

70

Displacement damage (dpa)

Figure 15 Yield stress as a function of displacement level for martensitic alloys irradiated in FFTF or EBR-II.
Reproduced from Kohno, Y.; Kohyama, A.; Hirose, T.; Hamilton, M. L.; Narui, M. J. Nucl. Mater. 1999, 271 & 272, 145.

have led to ferritic alloys with appreciable work hardening even at high displacement levels. Figure 14
shows tensile curves for the 9Cr–2Mo–1Ni steel,
JFMS, neutron irradiated and tested at room temperature. Uniform elongations of several percent are
evident, a reasonable value for irradiated steels.12
A plot of yield stress as a function of displacement

damage level is shown in Figure 15 for low-activation

ferritic alloys irradiated in fast reactors,30 and plots
of yield stress and total elongation are shown in
Figure 16 for fast and mixed-spectrum reactors.31
Unlike the austenitic alloys, the martensitic alloys
rapidly reach a peak in strength then soften with
further irradiation followed by near saturation in
strength beginning at about 30 dpa. Total elongation
follows a corresponding pattern, demonstrating


Effect of Radiation on Strength and Ductility of Metals and Alloys

111

800

JLF-5
(12Cr–2WVTa)

Total elongation (%)

JLF-5
(12Cr–2WVTa)

600

Yield stress (MPa)

30

JLF-4
(2.25Cr–2WVTa)

700

500
F82H
10B

400
300

F82H
STD

JLF-1
(9Cr–2WVTa)

F82H
(8Cr–2WVTaB)

JLF-1
(9Cr–2WVTa)

20

F82H
STD

F82H
(8Cr–2WVTaB)

10

JLF-3
(7Cr–2WVTa)

200

F82H
10B

: Irradiated in HFIR (spec. SS3)
Others in FFTF (spec. TS(s))

100

JLF-4
(2.25Cr–2WVTa)

0
0

10

20

30

40

50

60

Displacement damage (dpa)

0

0

10

20

30

40

JLF-3
(7Cr–2WVTa)

50

60

Displacement damage (dpa)

Figure 16 Yield strength of martensitic alloys following irradiation at 400  C in FFTF or high flux isotope reactor.

Reproduced from Kohyama, A.; Hishinuma, A.; Gelles, D. S.; Klueh, R. L.; Dietz, W.; Ehrlich, K. J. Nucl. Mater. 1996,
233–237, 138.

inverse behavior. This appears at first to be a phenomenon totally different from that which occurs in
austenitic alloys. However, the operable mechanisms
are really the same, just occurring at lower fluences.
Hardening mechanisms are similar, occurring by
point defect clusters at low irradiation temperatures
and transitioning to loops and network dislocations
and precipitation as temperature is increased. The
martensitic alloys are more complex in that the initial
microstructure is determined by the heat treatment.
The alloys are used in the normalized and tempered
condition produced by austenitizing the alloy and
quenching below the A1 temperature to produce
martensite. The martensite is then tempered below
the A1 temperature. Precipitates, primarily carbides
such as M23C6, form on prior austenite grain boundaries and on the martensite laths. In the case of the bcc
alloys, more rapid radiation-enhanced diffusion
results in irradiation-induced recovery and precipitate growth at lower fluence than in the fcc alloys.
Recall that in the austenitic alloys, saturation was
reached and sustained for a long period until microstructural coarsening resulted in a slight decrease in
strength at high displacement levels. The pronounced
peak in strength results from rapid hardening due to
irradiation-produced defects, but the effect of irradiation hardening is offset by irradiation-enhanced
recovery, resulting in a decrease in strength and hence
a peak in strength.32 The martensitic alloys demonstrate the same phenomena but at a lower fluence.

Figure 17 shows tensile test results from alloys
irradiated in FFTF to approximately 30 dpa. Tests at

the irradiation temperatures show high and nearly
constant values of total tensile elongation at temperatures above 425  C.33 Somewhat similar behavior
of elongation was also exhibited by 9Cr–1MoVNb
steel irradiated to 12 dpa in EBR-II.34 This alloy
also exhibits an increase in total elongation at 550  C
(Figure 18).
The apparent saturation in strength above about
450  C is also in agreement with hardness measurements. In a study of irradiated HT-9 and
9Cr–1MoVNb, Hu and Gelles observed that hardness
retained its unirradiated value upon irradiation to
26 dpa when irradiated at temperatures above about
450  C in EBR-II.35 From this behavior, it would
normally be concluded that fracture properties
would remain unchanged or improve with increasing
temperature. Although fracture toughness testing is
used extensively in the study and certification of
irradiated alloys, the Charpy impact test is more
commonly used in alloy development and fundamental research because the test requires smaller specimens and can be conducted more easily. Ductile to
brittle transition temperature (DBTT) is a useful tool
for comparison of alloys and assessment of radiation
damage. Charpy impact testing was conducted by
Hu and Gelles who observed that in the case of
the 9Cr–1MoVNb, the DBTT did in fact retain
essentially its unirradiated value for irradiation


112

Effect of Radiation on Strength and Ductility of Metals and Alloys

12
Total elongation (%)

Yield stress, sy (MPa)

800

600

400

Unirr.

200

8

4

0

(b)

600

700

800

900

1000

Uniform elongation (%)

Ultimate tensile stress, su (MPa)

(a)

0

800
600
Unirr.
400

Total elongation, er (%)

8

9Cr–1 MoVNb steel
Unirradiated
Aged
Irradiated
Test temperature @ Irradiation
temperature
@ Aging temperature

4

200
0
0

600

700

800

900

400

500
Test temperature (ЊC)

600

Figure 18 Elongation of 9Cr–1MoVNb irradiated in EBR-II
to 0.9 dpa. Reproduced from Klueh, R. L.; Vitek, J. M.
J. Nucl. Mater. 1985, 132, 27.

30

20
Unirr.
10

0

(c)

12

600
700
800
Irradiation temperature (K)

900

NFL-0
NFL-1
Figure 17 Tensile properties of two Fe–9Cr–2W steels
with and without small additions of boron, irradiated in FFTF
to approximately 30 dpa and tested at room temperature.
Reproduced from Kurishita, H.; Kayano, H.; Narui, M.;
Kimura, A.; Hamilton, M. L.; Gelles, D. S. J. Nucl. Mater.
1994, 212–215, 730.

temperatures above 450  C at 26 dpa but not for lower
irradiation temperatures, as shown in Figure 19.35
However, the DBTT of HT-9 failed to retain its
unirradiated value despite the absence of an increase
in strength and hardness. In both cases the upper shelf
energy was reduced by the irradiation, indicating
some changes in the irradiation microstructure.

The shift in the DBTT for both alloys at 13 and
26 dpa is shown as a function of irradiation temperature in Figure 20. The retention of the increase in the

DBTT at high temperatures illustrates the caution
that must be used in assessing ferritic alloys. In impact
testing, fracture is generally initiated at carbide particles. Even though coarsening of the carbides results
in less impediment to dislocation motion, and thus
less hardening, the stress intensity factor increases
for a nucleating crack at a hard particle so that the
effective crack length is the crack nucleus plus
the diameter of the carbide particle.36 The result is
that fracture toughness can increase with irradiation
temperature.
Despite the caution that must be taken in making
generalizations based upon tensile behavior, development of low-activation martensitic alloys has led to
alloys with very favorable fracture properties. The
DBTT is shown in Figure 21 for two low-activation
alloys, 9Cr–2WVTa and 9Cr–2WVTa–2Ni, with the
conventional Ht-9 and 9Cr–1MoVNb for comparison. The irradiation was done in EBR-II at irradiation
temperatures from 376 to 405 and to displacement
levels of 23–33 dpa. After irradiation, the two


Effect of Radiation on Strength and Ductility of Metals and Alloys

113

9Cr–1Mo base metal (TV series), 26 dpa
Ti = 390 ЊC
Ti = 500 ЊC

Normalized fracture energy (J cm−2)

600

Control

480
TV11
360

240
TV22

120

0
-150

-100

-50
0
50
Test temperature (ЊC)

100

150

HT-9 base metal (TT series), 26 dpa

Normalized fracture energy (J cm−2)

400

320

Ti = 390 ЊC
Ti = 450 ЊC
Ti = 500 ЊC
Ti = 550 ЊC

Control

240

160

TT23

TT15
TT20

80

0
-100

-50

0

TT28

50
100
150
Test temperature (ЊC)

200

250

Figure 19 Charpy impact test results for 9Cr–1Mo and HT-9 irradiated in EBR-II to 26 dpa. Reproduced from Hu, W. L.;
Gelles, D. S. Influence of Radiation on Material Properties; ASTM: Philadelphia, PA, 1987; ASTM STP 956, p 83.
Reprinted, with permission, from Influence of Radiation on Material Properties, copyright ASTM International, West
Conshohocken, PA.

conventional alloys had DBTT values above room
temperature, whereas the values for the two
reduced-activation steels were below À75  C.32 The
figure clearly shows that HT-9 is not the alloy of
choice for nuclear applications. Despite the Ni content of 9Cr–2WVTa–2Ni, this class of alloys can
prove useful in nuclear applications such as fast
reactors where the thermal flux is small and the
very high-energy neutrons characteristic of a fusion
reaction are absent.

1.04.7.3

Helium Effects

Helium effects are important for systems that generate high-energy neutrons such as fusion reactors and
spallation targets that encounter high-energy protons. The 14 MeV neutrons produced by D–T fusion
will produce (n,a) reactions in nearly all common
structural elements such that ferritic steels are not
exempt from helium generation. However, in the
absence of nickel, helium generation rates are lower.


114

Effect of Radiation on Strength and Ductility of Metals and Alloys

140
Shift in transition temperature (ЊC)

13 dpa HT9

120

26 dpa HT9

100

26 dpa Mod 9Cr–1Mo

13 dpa Mod 9Cr–1Mo

80
60
40

20
0
500
450
Irradiation temperature (ЊC)

400

550

Figure 20 Shift in ductile to brittle transition temperature
as a function of irradiation temperature for 13 and
26 dpa following EBR-II irradiation. Reproduced from
Hu, W. L.; Gelles, D. S. Influence of Radiation on Material
Properties; ASTM: Philadelphia, PA, 1987; ASTM STP
956, p 83. Reprinted, with permission, from Influence of
Radiation on Material Properties, copyright ASTM
International, West Conshohocken, PA.

75

Transition temperature (ЊC)

50

Unirradiated
Irradiated

25
0

-25

9Cr-1MoVNb
12Cr-1MoVW

-50
-75
-100
-125

9Cr-2WVTa
9Cr-2WVTa-2Ni

Figure 21 Ductile to brittle transition temperature before
and after irradiation in EBR-II to 23–33 dpa at 376–405  C
comparing conventional alloys with irradiation-resistant
martensitic alloys. Reproduced from Klueh, R. L.; Sokolov,
M. A.; Hashimoto, N. J. Nucl. Mater. 2008, 374, 220.
Reprinted, with permission, from Influence of Radiation on
Material Properties, copyright ASTM International, 100 Barr
Harbor Drive, West Conshohocken, PA 19428.

In austenitic stainless steels, the high nickel content is
used to introduce helium to simulate the fusion environment. However, the nickel content is so high that
unrealistically high concentrations of helium are produced in a thermal or mixed-spectrum reactor. As a
result, spectral tailoring is necessary to achieve the
correct He per dpa ratio.37

In the case of martensitic alloys, several techniques have been used to generate helium in fission
reactors. Doping with natural nickel has been used

in 9Cr and 12Cr alloys27,38,39 and isotopically separated nickel has been used to discriminate against
the effect of nickel as opposed to the effect of He.
The isotope 59Ni has been used as it will generate
He but 60Ni, used as a control, will not.40 Doping
with 10B has also been used to introduce He into this
class of alloys, leading to concentrations of several
hundred parts per million.41–43 Nickel doping experiments with 12Cr–1MoVW were conducted in the
HFIR using 1 and 2% Ni to generate up to about
300 appm He. This experiment demonstrated what
appeared to be a strengthening effect of helium.44
However, it was determined that these results were
clouded by the fact that doping with Ni lowers the
A1 temperature of the alloy, leading to untempered
martensite upon cooling from the tempering
temperature. Adjustments were made in the tempering temperature, but the additional variables cast
doubt on the results.44 The isotope separation experiments and the boron doping experiments found no
clear indication of a helium effect.40,41 Several
mechanisms for hardening by helium have been identified,45 but Klueh and Harries27 have, after a careful
review, come to the conclusion that the effect of
helium on strength and ductility below 500  C is
inconclusive and that if there is an effect, it is probably small and of minor significance for the conditions
examined.
As with austenitic stainless steels, the effect of
hardening, if it exists for ferritic alloys, is of minor
importance compared with high-temperature intergranular embrittlement. High He concentrations and
higher temperatures have been investigated in martensitic alloys by means of helium implantation with
accelerators. Hasegawa implanted He at 600  C to
levels of 500 appm and performed tensile tests at
this temperature.46 However, all fractures were transgranular. Bae et al.47 implanted similar levels of
helium, 500 appm, in the 12Cr-steel, MANET, at

temperatures as high as 500  C. They observed hardening, shown in Figure 22, but again they were only
transgranular fractures.
Bae et al.47 also observed no synergistic effect of
500 appm H and 500 appm He and no effect of hydrogen alone at this level.
Jung et al.48 found that He implanted in the 9Cr
alloy, EUROFER97, to concentrations as high as
1250 appm produced both hardening and reduction
in ductility when implanted at 250  C and tested at


Effect of Radiation on Strength and Ductility of Metals and Alloys

115

20

900
15
Total elongation (%)

MANET

Yield stress (MPa)

700

500

MANET

10

5
300

DSA
100

0
0

600
200
400
Irradiation and test temperature (ЊC)

800

0

200
400
600
Irradiation and test temperature (ЊC)

800

0.32 dpa, 500 appm H, 500 appm He
0.30 dpa, 500 appm He
0.02 dpa, 500 appm H

Unirradiated

Figure 22 Yield strength and total elongation for the 12Cr steel, MANET 1 cyclotron implanted with He and H. Reproduced
from Bae, K. K.; Ehrlich, K.; Mosalang, A. J. Nucl. Mater. 1992, 191–194, 905.

1200

0.125

EUROFER97
0.25

Timp = 250 ЊC

s (MPa)

1000
0.25
0.06
0.125
0.06

800
600

cHe(at.%)
0
Ttest = 25 ЊC

400

0
Ttest = 250 ЊC

200
0

0

5
e (%)

10

Figure 23 Stress–strain curves for EUROFER97 for
cyclotron implantation of helium. Jung, P.; Henry, J.;
Chen, J. J. Nucl. Mater. 2005, 343, 275.

either 25 or 250  C. Their data demonstrate both
hardening and loss of work hardening with helium,
as clearly shown in Figure 23.
Severe intergranular embrittlement was, in fact,
demonstrated by Jung et al. in the 9Cr steels T91 and
EM10 when levels of 5000 appm were attained by
cyclotron implantation.49 In cases where He was
implanted at 250  C and tested at 25  C and at
250  C, clear intergranular fracture was experienced.
For implantation temperatures of 550  C, where

intergranular embrittlement would be expected in
austenitic stainless steels, the fracture surfaces indicated a return of some ductility with little or no

intergranular fracture, particularly when tested at
550  C, Figure 24. The latter case can perhaps be
explained by the loss of strength at 550  C where
plastic deformation blunts any nucleating crack
before the local stress at a grain boundary is sufficiently high for grain separation. In the case of
implantation at 550  C and testing at 25  C, the fracture mechanism is less clear. Here, the temperature is
sufficiently high for diffusion of He to the grain
boundaries, but perhaps capture by other sinks prevents high levels of He at the boundaries. More
investigation of the high-temperature He embrittlement is necessary to determine the underlying mechanism of fracture. It is clear, however, that helium
embrittlement is more effective in austenitic stainless
steels than in ferritic–martensitic steels. Since nickel
is the largest source of helium in a fast neutron
environment, ferritic alloys clearly have a lower He
generation rate than austenitic steels.45 The combination of higher resistance to helium embrittlement
and the lower generation rate of He in ferritic alloys
makes this class of alloys more favorable with respect
to He embrittlement.


116

Effect of Radiation on Strength and Ductility of Metals and Alloys

250/25

550/25

250/250

550/550

Figure 24 Fracture surfaces of T91 following helium
implantation at 250 and 550  C and tested at the
implantation temperatures indicated. Reproduced from
Jung, P.; Henry, J.; Chen, J.; Brachet, J. C. J. Nucl. Mater.
2003, 318, 241.

1.04.8 Refractory Metals
1.04.8.1

Tensile Behavior

The refractory metals are the metals in groups V and
VI of the periodic table: vanadium, niobium, and
tantalum in group V and chromium, molybdenum,
and tungsten in group VI. All have the characteristic
of a high melting point, hence the term refractory.
The group VI metals are typically brittle, even without irradiation. For example, chromium is almost
never used pure or as a major alloy element, although
it is invaluable as a minor alloying element. Molybdenum and tungsten are both brittle in nature but can
be made into useful structural alloys by controlling
interstitial impurities and by the addition of minor
elements. In contrast to the brittle behavior of the
group VI metals, the group V metals are inherently
ductile. Structural alloys based upon this group have
been developed, primarily for very high temperature
and space applications. The primary disadvantage of
the refractory metals is their formation of volatile
oxides as opposed to protective oxide layers. Vanadium and molybdenum oxides have melting points
below metal working temperatures so that the metals

become wet and can have liquid oxide drip off them.
Unlike the tensile behavior of fcc metals, where
there is a smooth increase in strength as plastic deformation proceeds and work hardening progresses, bcc

metals typically exhibit a load drop, or yield point,
almost immediately following the onset of plastic
deformation, as discussed in Section 1.04.7.2.
In the case of refractory metals, mechanical properties are largely determined by interstitial solutes.
High purity refractory metals do not exhibit a yield
point but behave more like fcc metals. Niobium
alloys irradiated in Li at 1200 C for over three
months in EBR-II had total elongations of about
60%. Despite any irradiation hardening, the near
absence of oxygen resulted in a very soft material at
these high temperatures.50 However, since interstitials are nearly always present, tensile behavior is
more typically characteristic of bcc metals.
Irradiation-produced defects interact with interstitial elements, resulting, in some cases, in severe
embrittlement. The tantalum alloy, Westinghouse T111 (Ta–8W–2Hf) is used in Figure 25 to illustrate a
commonly observed phenomenon of plastic instability.51 Plastic deformation becomes local, with high
levels of slip on closely spaced planes where dislocations sweep out the irradiation-generated defects
giving rise to local channels of very high deformation.
This phenomenon, called channel deformation, is very
common in irradiated metals. The result is a sudden
and severe load drop with the fracture surface showing
what appears to be a completely ductile chisel point
fracture.52 Addition of 405 wt ppm oxygen to T-111
results in a cleavage fracture with no measurable plastic
deformation, as shown in Figure 26.51 In both Figures
25 and 26, corresponding unirradiated alloys are
shown demonstrating ductile behavior. In the unirradiated condition, the addition of 400 wt ppm oxygen

has only minor effects on strength and ductility, as can
be concluded by a comparison of Figures 25 and 26.
However, irradiation hardening superimposed upon
the oxygen interstitial hardening appears to raise the
yield stress above the cleavage stress for the alloy. It is
suggested that interstitial solutes such as oxygen diffuse to irradiation-produced defect clusters, enhancing their hardening effect.53,54 All three behaviors are
observed in irradiated refractory metals: ductile with
hardening, plastic instability, and cleavage fracture in
the elastic range.55
The synergism between interstitial hardening and
irradiation hardening does not necessarily lead to
immediate catastrophic embrittlement. This behavior
is shown in Figure 27 for vanadium containing a very
high level of oxygen, 2100 wt ppm. Irradiation to a
fluence level of 1.5 Â 1019 (E > 1 MeV) leads to the
familiar plastic instability but with several per cent
plastic strain.53


Effect of Radiation on Strength and Ductility of Metals and Alloys

117

1600
111T

1400

Unirradiated
EBR-II irradiated

ft (E > 0.1 MeV) = 1.6 × 1026 n m–2
Tirrad = Ttest = 873 K

Stress (MPa)

1200
1000
800

*

600
400
200
0

*
0

2

4

6

8

10

12 14 16

Strain (%)

18

20

22

24

26

28

Figure 25 Stress–strain curves for the Ta alloy, T-111 showing characteristic tensile behavior following irradiation in EBR-II
to 8 dpa at 600  C. Reproduced from Grossbeck, M. L.; Wiffen, F. W. In Space Nuclear Power Systems; El-Genk, M. S.,
Hoover, M. D., Eds.; Orbit Book Co.: Malabar, FL, 1986; Vol. III, p 85.

1400
111 T + 405 wt. ppm

oxygen
Unirradiated
EBR-II irradiated
ft (E > 0.1 MeV) = 1.4 × 1026 n m–2
Tirrad = Ttest = 853 K

*

1200

Stress (MPa)

1000
800
600
400

*

200
0

0

2

4

6

8

10

12 14 16
Strain (%)

18

20

22

24

26

28

Figure 26 Stress–strain curves for oxygen-doped T-111 T unirradiated and irradiated in EBR-II to 7 dpa at 580  C.
Reproduced from Grossbeck, M. L.; Wiffen, F. W. In Space Nuclear Power Systems; El-Genk, M. S., Hoover, M. D., Eds.;
Orbit Book Co.: Malabar, FL, 1986; Vol. III, p 85.

Interstitial solutes, especially oxygen, may be
controlled by the addition of gettering elements. In
the vanadium system, titanium has been successful.
Alloys in the V–Cr–Ti system have been studied for
application to fusion reactors. In refractory metal

alloys, it is the oxygen in solution that is detrimental, so that the oxygen must be combined with the
titanium.56 This usually requires a heat treatment of
sufficiently long times and high temperatures to precipitate the oxygen. In the vanadium–titanium system,


118

Effect of Radiation on Strength and Ductility of Metals and Alloys

90

Vanadium–2100 wt ppm oxygen
Tensile tests at 300 ЊK
and 1.67 ´ 10-4 s-1

80

Irradiated, 1.5 ´ 1019 n cm−2 (E > 1 MeV)

70

Stress (Kpsi)

60
Unirradiated

50
40
30
20
10
0

Uniform
elongation
0

2

4

6

8

12
14
10
Strain (%)

16

18

20

22

24

Figure 27 Stress–strain curves for oxygen-doped vanadium irradiated at 85  C at the Ames Laboratory Research
Reactor. Reproduced from Wechsler, M. S.; Alexander, D. G.; Bajaj, R.; Carlson, O. N. In Defects and Defect Clusters in B.C.
C. Metals and Their Alloys, Nuclear Metallurgy; Arsenault, R. J., Ed.; National Bureau of Standards: Gaithersburg, MD, 1973;
Vol. 18, p 127.

a heat treatment of two hours at 950  C has been
found sufficient to achieve ductility, whether or not
it is to be irradiated.57,58 Confusion over oxygen present in solution and oxygen combined in precipitates is
believed to be one reason for the disparity in tensile
data for this class of alloys and perhaps accounts
for the relatively high level of ductility observed in

Figure 27.
Upon irradiation of alloys in the range of V–3–
5Cr–3–5Ti in the HFIR, no cleavage fracture without plastic deformation was observed.59,60 However,
plastic instability was commonly observed at irradiation and test temperatures below 400  C. Irradiations
in the range of 4–6 dpa in the HFIR produced
uniform elongations from 0.2 to 0.6% and total elongations below 4%. Corresponding irradiations at
500  C did not reveal plastic instability and produced
uniform elongations in the range of 2–5%.59,60 Irradiations to 3–5 dpa in the advanced test reactor
(ATR) demonstrated plastic instability for irradiation
and test temperatures of about 200  C, with uniform
elongations below 0.5%.61 Irradiations conducted in
the high flux beam reactor (HFBR) at exposures of
only 0.1 and 0.5 dpa corroborated these results and
demonstrated a transition in the fracture mechanism
between 300 and 400  C, resulting in a significant
increase in ductility at temperatures above 400  C,
Figure 28.62

1.04.8.2

Helium Effects

Helium is conveniently introduced into nickelbearing alloys through thermal neutron irradiation.
Although helium is usually detrimental, especially if
the material is to be subsequently welded,63 it offers
a method to simulate the production of helium
expected in the very hard spectrum of a fusion reactor. In the case of vanadium and other refractory
metal alloys, the effect of helium has been studied
using two primary methods of introduction of
helium. One method is implantation of a-particles

with an accelerator; the other is the use of the decay
of tritium. Tritium rapidly diffuses into group
V refractory metals at elevated temperatures. The
elevated temperature serves more to dissolve the
protective oxide layer than to accelerate the kinetics
of dissolution. The tritium thus introduced is permitted to decay, by b-decay with a residual nucleus of
3
He. Helium-doped specimens have subsequently
been neutron irradiated to study the synergistic
effects of helium and atomic displacement damage.
A limited number of experiments have used techniques to simultaneously implant He and produce
atomic displacements through an irradiation environment of Li. The concept of introducing tritium into
an irradiation capsule with the specimens in contact
with lithium has been investigated to study vanadium


Effect of Radiation on Strength and Ductility of Metals and Alloys

12

140
0.0002

6
0.5 dpa
(V1–V3)

4
2
0

Ttest = Tirr
0

100

200

300

Molybdenum

120

0.1 dpa
V4

8

Ultimate tensile stress (1000 psi)

Uniform elongation (%)

10

-2

119

400

500

600

100

700

40

0
60
Elongation (%)

alloys with the He/dpa ratio characteristic of a fusion
environment. The tritium charge, the production of
tritium from lithium, and the production of tritium
from 3He are some of the important considerations in
the design of the experiment.64 Although conceptually
valid, the desired results have not yet been obtained
with experiments of this type.
Cyclotron-implanted helium has been used, also
to study the effects of the fusion irradiation environment. Tanaka showed severe embrittlement with
the introduction of 90 and 200 appm He at 700  C
in V–20Ti.65 Grossbeck and Horak showed that a
level of 80 at. ppm He implanted as part of the same
experiment had no significant effect on elongation in
V–15Cr–5Ti at 700  C.52 Braski also observed no significant effect on ductility in V–15Cr–5Ti at 600  C
with similar levels of He introduced from decay of

tritium.66 The alloys, Vanstar-7 and V–3Ti–1Si, were
also investigated, in some cases with an improvement
in ductility upon introduction of helium.66 Following irradiation, severe embrittlement was observed
in V–15Cr–5Ti at 600  C in tritium trick samples
by Braski66 and at 625  C in cyclotron-implanted
samples by Grossbeck and Horak.52 Irradiation
experiments with refractory metals, unless using a
Li environment, frequently subject the specimens to
contamination by interstitial impurities, also leading
to embrittlement.52,67
Unalloyed molybdenum, Mo–0.5Ti, and Mo–50Re
were irradiated in EBR-II by Wiffen at exposures of

0.0002

60

20

Test temperature (ЊC)

Figure 28 Uniform elongation of V–4Cr–4Ti irradiated in
the high flux beam reactor. Reproduced from Snead, L. L.;
Zinkle, S. J.; Alexander, D. J.; Rowcliffe, A. F.;
Robertson, J. P.; Eatherly, W. S. Fusion Reactor Materials
Semiannual Progress Report for Period Ending, Dec 31,
1997; DOE/ER-0313/23, p 81.

80

Controls
3.5 or 4.0 ´ 1022 n cm−2 at 455 ЊC
6.1 ´ 1022 n cm−2 at 857–1136 ЊC
4.0 ´ 1022 n cm−2 at 455 ЊC + 1050 ЊC anneal

40
20
0.0002
0
-200

0

200
400
Test temperature (ЊC)

0.0002
600

Figure 29 Ultimate tensile strength and total elongation
for molybdenum in the unirradiated condition and irradiated
in EBR-II to 20–30 dpa. Dashed lines connect results where
irradiation conditions or strain rates were not constant.
Reproduced from Wiffen, F. W. In Defects and Defect
Clusters in B.C.C. Metals and Their Alloys, Nuclear
Metallurgy; Arsenault, R. J., Ed.; National Bureau of
Standards: Gaithersburg, MD, 1973; Vol. 18, p 176.

3.5–6.1 n cmÀ2 (E > 0.1 MeV) (18–32 dpa).68 Although

Mo alloys are known to exhibit increased ductility
with increasing temperature in the unirradiated condition, at temperatures above 400–550  C, all three
materials suffered plastic instability with uniform
elongations below about 0.5%. This effect is shown
in Mo in Figure 2968 where irradiation temperature
is shown to be the critical parameter and where specimens irradiated at 455 and 1136  C were embrittled
even in room temperature tests.
This class of alloys is discussed further in Chapter
4.06, Radiation Effects in Refractory Metals and
Alloys.

1.04.9 Amorphous Metals
Stable metallic glasses may be produced, commonly in intermetallic compounds. Interest in the


Irradiation dose required for amorphization (dpa)

120

Effect of Radiation on Strength and Ductility of Metals and Alloys

16.0
14.0
12.0
40Ar
irradiation

10.0
8.0
6.0

4.0
2.0
0

0

100

200 300 400 500 600
Irradiation temperature (K)

700

Figure 30 Irradiation displacement level as a function of
temperature for 0.9 MeV electron and 0.5–1.5 MeV Ar ion
irradiation. The family of curves is for several dpa rates of
1.04–1.83 mdpa sÀ1. Reproduced from Howe, L. M.;
Phillips, D.; Motta, A. T.; Okamoto, P. R. Surface Coatings
Tech. 1994, 66, 411.

irradiation properties of this class of materials
resulted from preliminary tests that showed that
these materials actually became more ductile upon
irradiation.69 Other intermetallic compounds have
been shown to become amorphous upon irradiation.
Although semiconductors such as Si and Ge are
susceptible to amorphization under irradiation, the
phenomenon is almost exclusively restricted to intermetallic compounds.70 To mention only a few, Zr3Al,
Mo3Si, Nb3Ge, and Fe2Mo are compounds that have
been studied in the amorphous state. Results and a

detailed review of mechanisms and theories of amorphization have been published by Motta.70 In simple
terms, the lattice disruption and defect generation
from irradiation disrupts long-range order in the
system. Thermal annealing competes with the disordering so that there is a critical temperature above
which amorphization is not possible. Figure 30
shows a plot of the irradiation exposure necessary
for amorphization as a function of temperature for
Zr3Fe.71 The critical temperatures and the necessary
exposures are both functions of the material as well as
the impinging particle. Once formed, the amorphous
phases are stable under irradiation, but the critical
temperatures are typically lower than would be experienced for structural materials in nuclear systems.
They are of interest, however, because some intermetallic phases, such as Fe2Mo and Fe3B found
in commercial alloys, become amorphous under

irradiation.70,72 In the example of Zr3Fe, the critical
temperature under argon ion irradiation is approximately 250  C, a temperature too low for most, but
not all, reactors.
The intermetallic alloys that can be produced in
the amorphous state before irradiation are of more
interest as potential structural materials, although
they remain in the category of research interest at
the present time. In addition to the increase in ductility upon irradiation, the absence of a crystalline
structure with interacting dislocations was further
incentive to investigate the irradiation properties of
this class of materials. Metallic glasses containing
boron, such as Fe40Ni40B20 and (Mo.6Ru.4)82B18 are
a few examples, with the former receiving the most
attention in terms of mechanical properties.69,73–75
Amorphous alloys are complex systems where

changes in free volume and segregation into clusters
of differing composition result in changes in behavior
as irradiation proceeds. Investigation of the Fe–Ni–B
alloy has shown that ductility first decreases and then
increases with increasing fluence due to the competing effects of free volume and formation of regions of
boron-depleted and boron-rich clusters.73 For sufficiently high fluences, the result is severe embrittlement. In the case of alloys based on the intermetallic
Zr3Al, very severe embrittlement upon irradiation
is attributed to the formation of new amorphous
phases.76
Even though a crystal structure is absent, the
atoms may be dislodged from their locations, creating
additional free volume. Without the bonds present
from a crystal lattice, the low binding energy results
in high displacement levels for fluence levels that
what would be considered low in crystalline alloys.
Fluence levels in the range of 1016–1021 n cmÀ2 have
been investigated resulting in displacement levels
exceeding 100 dpa. However, simply having similar
displacement levels does not permit a true comparison with crystalline materials. Much research is necessary before this class of materials becomes of
commercial importance.

1.04.10 Conclusions
The relationship between the irradiation-induced
microstructure and tensile properties has been briefly
presented using representative classes of alloys. The
austenitic stainless steels are an important class of
alloys, and they are less complex than the martensitic
steels. In the unirradiated condition, the austenitic

Effect of Radiation on Strength and Ductility of Metals and Alloys

alloys are primarily hardened by dislocation reactions
leading to conventional work hardening, and the
martensitic steels are hardened by phase transformations requiring careful heat treatments. The primary
irradiation effects are similar, but they influence
microstructure and, therefore, behavior in different
ways. Both types of alloys have important applications in the nuclear field. Helium embrittlement
might be the most important, considering the use of
alloys in a neutron environment at high temperatures. For the proper conditions, helium can nearly
always cause catastrophic failure. Repair welding of
alloys with as little as 1–10 appm helium can lead to
severe intergranular cracking.
The refractory metals are useful for space reactor
application because of their liquid metal compatibility and their high-temperature strength. Space reactors can lose heat only by thermal radiation,
necessitating high temperatures. However, this class
of alloys is most susceptible to embrittlement by
interstitial impurities, and synergism of impurities
with irradiation-induced defects is an area that must
be addressed further.
Amorphous alloys are a research curiosity in that
they have interesting properties with respect to
irradiation but little application at the present
time. Any new class of alloys must be understood
before it can be engineered, so research is the
essential beginning.

References
1.

2.
3.
4.
5.
6.

7.
8.
9.

10.
11.
12.

Simons, R. L.; Hulbert, K. A. In Effects of Radiation on
Materials; Garner, F. A., Perrin, J. S., Eds.; ASTM:
Philadelphia, PA, 1985; Vol. II, p 820.
Bement, A. L. In Proceedings on the Strength of Metals
and Alloys, ASM: Petals Park, 1970; p 693.
Koppenaal, T. J.; Kuhlmann-Wilsdorf, D. Appl. Phys. Lett.
1964, 4, 59.
Taylor, G. I. Proc. Royal Soc. 1934, 145, 362.
Seeger, A. Dislocations and Mechanical Properties of
Crystals; Wiley: New York, 1957; p 243.
Olander, D. R. Fundamental Aspects of Nuclear Reactor
Fuel Elements; Technical Information Center, ERDA: Oak
Ridge, TN, 1976; p 441.
Grossbeck, M. L.; Maziasz, P. J.; Rowcliffe, A. F. J. Nucl.
Mater. 1992, 191–194, 808.
Kelly, P. M. Scripta Met. 1972, 6, 647.

Fish, R. L.; Cannon, N. S.; Wire, G. L. In Effects of
Radiation on Structural Materials; Sprague, J. A.,
Dramer, K., Eds.; ASTM: Philadelphia, PA, 1979;
ASTM STP 683, p 450.
Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.;
Johnson, G. D. J. Nucl. Mater. 1981, 103 & 104, 803.
Grossbeck, M. L. J. Nucl. Mater. 1991, 179–181, 568.
Maloy, S. A.; Toloczko, M. B.; McClellan, K. J. et al.
J. Nucl. Mater. 2006, 356, 62.

13.

121

Allen, T. R.; Tsai, H.; Cole, J. I.; Ohta, J.; Dohi, K.;
Kusanagi, H. Effects of Radiation on Materials; ASTM:
Philadelphia, PA, 2004; ASTM STP 1447, p 3.
14. Maloy, S. A.; James, M. R.; Romero, T. J.; Toloczko, M. B.;
Kurtz, R. J.; Kimura, A. J. Nucl. Mater. 2005, 341, 141.
15. Norget, M. J.; Robinson, M. T.; Torrens, I. M. Nucl. Eng.
Des. 1975, 33, 91.
16. Wiffen, F. W.; Maziasz, P. J.; Bloom, E. E.; Stiegler, J. O.;
Grossbeck, M. L. In Proceedings of Symposium on the
Metal Physics of Stainless Steels, AIME, 1978.
17. Gabriel, T. A.; Bishop, B. L.; Wiffen, F. W. Calculated
Irradiation Response of Materials Using a Fusion-Reactor
First-Wall Neutron Spectrum, ORNL/TM-5956; Oak Ridge
National Laboratory: Oak Ridge, TN, 1977.
18. Garner, F. A.; Greenwood, L. R. Mater. Trans. JIM 1993,
34, 985.

19. Grossbeck, M. L. Tensile Properties of HFIR-Irradiated
Austenitic Stainless Steels at 250 to 400  C from the
European Community/U.S./Japan Fusion Materials
Collaboration; Fusion Reactor Materials Semiannual
Progress Report for Period Ending, Sept 30, 1986;
DOE/ER-0313/1, pp 254–263.
20. Ehrlich, K. J. Nucl. Mater. 1985, 133 & 134, 119.
21. Grossbeck, M. L.; Ehrlich, K.; Wassilew, C. J. Nucl. Mater.
1990, 174, 264.
22. Grossbeck, M. L. Development of Tensile Property
Relations for ITER Data Base; Fusion Reactor Materials
Semiannual Progress Report for Period Ending, Mar 31,
1989; DOE/ER-0313/06, p 243.
23. Hamilton, M. L.; Cannon, N. S.; Johnson, G. D. In Effects
of Radiation on Materials; Brager, H. R., Perrin, J. S., Eds.;
ASTM: Philadelphia, PA, 1982; ASTM STP 782, p 636.
24. Garr, K. R.; Rhodes, C. G.; Kramer, D. In Effects of
Radiation on Substructure and Mechanical Properties of
Metals and Alloys; ASTM: Philadelphia, PA, 1973; ASTM
STP 529, p 109.
25. Hawthorne, J. R.; Reed, J. R.; Sprague, J. A. In Effects of
Radiation on Materials; ASTM: Philadelphia, PA, 1985;
ASTM STP 870, p 580.
26. Smidt, F. A.; Malmberg, P. R.; Sprague, J. A.;
Westmoreland, J. E. In Irradiation Effects on the
Microstructure and Properties of Metals; ASTM:
Philadelphia, PA, 1976; ASTM STP 611, p 227.
27. Klueh, R. L.; Harries, D. R. High-Chromium Ferritic and
Martensitic Steels for Nuclear Applications; ASTM:
Philadelphia, PA, 2001.

28. Tetelman, A. S.; McEvily, A. J. Fracture of Structural
Materials; Wiley: New York, 1967; p 178.
29. Wilson, J. C. Effects of irradiation on the structural materials
in nuclear power reactors. In Proceedings of the Second
United Nations International Conference on the Peaceful
Uses of Atomic Energy, United Nations, 1958; Vol. 5, p 431.
30. Kohno, Y.; Kohyama, A.; Hirose, T.; Hamilton, M. L.;
Narui, M. J. Nucl. Mater. 1999, 271 & 272, 145.
31. Kohyama, A.; Hishinuma, A.; Gelles, D. S.; Klueh, R. L.;
Dietz, W.; Ehrlich, K. J. Nucl. Mater. 1996, 233–237, 138.
32. Klueh, R. L.; Sokolov, M. A.; Hashimoto, N. J. Nucl. Mater.
2008, 374, 220.
33. Kurishita, H.; Kayano, H.; Narui, M.; Kimura, A.;
Hamilton, M. L.; Gelles, D. S. J. Nucl. Mater. 1994,
212–215, 730.
34. Klueh, R. L.; Vitek, J. M. J. Nucl. Mater. 1985, 132, 27.
35. Hu, W. L.; Gelles, D. S. Influence of Radiation on Material
Properties; ASTM: Philadelphia, PA, 1987; ASTM STP 956,
p 83.
36. Klueh, R. L.; Shiba, K.; Sokolov, M. A. J. Nucl. Mater. 2008,
377, 427.
37. Grossbeck, M. L.; Horak, J. A. J. Nucl. Mater. 1988,
155–157, 1001.


122

Effect of Radiation on Strength and Ductility of Metals and Alloys

38. Klueh, R. L.; Vitek, J. M.; Grossbeck, M. L. J. Nucl. Mater.

1981, 103 & 104, 887.
39. Kasada, R.; Kimura, A.; Matsui, H.; Narui, M. J. Nucl.
Mater. 1998, 258–263, 1199.
40. Gelles, D. S.; Hankin, G. L.; Hamilton, M. L. J. Nucl. Mater.
1997, 251, 188.
41. Shiba, K.; Suzuki, M.; Hishinuma, A.; Pawel, J. E. Effects of
Radiation on Materials; ASTM: Philadelphia, PA, 1996;
ASTM STP 1270, p 753.
42. Shiba, K.; Hishinuma, A. J. Nucl. Mater. 2000, 283–287,
474.
43. Kiumura, A.; Morimura, T.; Narui, M.; Matsui, H. J. Nucl.
Mater. 1996, 233–237, 319.
44. Klueh, R. L.; Vitek, J. M.; Grossbeck, M. L. Effects of
Radiation on Materials; ASTM: Philadelphia, PA, 1982;
ASTM STP 782, p 648.
45. Mansur, L. K.; Grossbeck, M. L. J. Nucl. Mater. 1988,
155–157, 130.
46. Hasegawa, A.; Shiraishi, H.; Matsui, H.; Abe, K. J. Nucl.
Mater. 1994, 212–215, 720.
47. Bae, K. K.; Ehrlich, K.; Mosalang, A. J. Nucl. Mater. 1992,
191–194, 905.
48. Jung, P.; Henry, J.; Chen, J. J. Nucl. Mater. 2005, 343,
275.
49. Jung, P.; Henry, J.; Chen, J.; Brachet, J. C. J. Nucl. Mater.
2003, 318, 241.
50. Grossbeck, M. L.; Heestand, R. L. Effect of irradiation
on the tensile of niobium-base alloys. In Proceedings
of Fourth Symposium on Space Nuclear Power
Systems, Albuquerque, NM, Jan 15, 1987; CONF-870118,
p 151.

51. Grossbeck, M. L.; Wiffen, F. W. In Space Nuclear Power
Systems; El-Genk, M. S., Hoover, M. D., Eds.; Orbit Book
Co.: Malabar, FL, 1986; Vol. III, p 85.
52. Grossbeck, M. L.; Horak, J. A. In Influence of Radiation on
Material Properties; Garner, F. A., Henager, C. H.,
Igata, N., Eds.; ASTM: Philadelphia, PA, 1987; ASTM STP
956, p 291.
53. Wechsler, M. S.; Alexander, D. G.; Bajaj, R.; Carlson, O. N.
In Defects and Defect Clusters in B.C.C. Metals and Their
Alloys, Nuclear Metallurgy; Arsenault, R. J., Ed.; National
Bureau of Standards: Gaithersburg, MD, 1973; Vol. 18,
p 127.
54. Ohr, S. M.; Tucker, R. P.; Wechsler, M. S. Phys. Stat. Sol.
A 1970, 2, 559.
55. Wiffen, F. W. Effects of irradiation on properties of
refractory alloys with emphasis on space power reactor
applications. In Proceedings of Refractory Alloy
Technology for Space Nuclear Power Applications,
USDOE, 1984; CONF-8308130, p 252.
56. DiStefano, J. R.; Hendricks, J. W. Nucl. Tech. 1995, 110,
145.
57. DiStefano, J. R.; DeVan, J. H. J. Nucl. Mater. 1997, 249,
150.
58. Grossbeck, M. L.; King, J. F.; Hoelzer, D. T. J. Nucl. Mater.
2000, 283–287, 1356.

59.

60.

61.

62.

63.
64.
65.
66.

67.

68.

69.
70.
71.
72.

73.

74.
75.
76.

Yan, Y.; Tsai, H.; Pushis, D. O.; Smith, D. L.; Fukumoto, K.;
Matsui, H. In Tensile Properties of V–(Cr,Fe)–Ti Alloys After
Irradiation in the HFIR-11J Experiment; Fusion Reactor
Materials Semiannual Progress Report for Period Ending,
Jun 2000 DOE/ER-0313/28, p 62.
Fukumjoto, K.; Matsui, H.; Yan, Y.; Tsai, H.; Strain, R. V.;

Smith, D. L. Tensile Properties of V–(Cr,Fe)–Ti Alloys After
Irradiation in the HFIR-12J Experiment; Fusion Reactor
Materials Semiannual Progress Report for Period Ending,
Dec 1999; DOE/ER-0313/27, p 70.
Hamilton, M. L.; Toloczko, M. B.; Oliver, B. M.;
Garner, F. A. Effect of Low Temperature Irradiation in ATR
on the Mechanical Properties of Ternary V–Cr–Ti Alloys;
Fusion Reactor Materials Semiannual Progress Report for
Period Ending, Dec 1999; DOE/ER-0313/27, p 76.
Snead, L. L.; Zinkle, S. J.; Alexander, D. J.; Rowcliffe, A. F.;
Robertson, J. P.; Eatherly, W. S. Fusion Reactor Materials
Semiannual Progress Report for Period Ending, Dec 31,
1997; DOE/ER-0313/23, p 81.
Li, S.; Grossbeck, M. L.; Zhang, Z.; Shen, W.; Chin, B. A.
Welding of irradiated materials. Welding Journal.
Gomes, I. C.; Tsai, H.; Smith, D. L. J. Nucl. Mater. 1999,
271 & 272, 349.
Tanaka, M. P.; Bloom, E. E.; Horak, J. A. J. Nucl. Mater.
1981, 103 & 104, 895.
Braski, D. N. In Influence of Radiation on Material
Properties; Garner, F. A., Henager, C. H., Igata, N.,
Eds.; ASTM: Philadelphia, PA, 1987; ASTM STP 956,
p 271.
Bohm, H. In Defects and Defect Clusters in B.C.C. Metals
and Their Alloys, Nuclear Metallurgy; Arsenault, R. J., Ed.;
National Bureau of Standards: Gaithersburg, MD, 1973;
Vol. 18, p 163.
Wiffen, F. W. In Defects and Defect Clusters in B.C.C.
Metals and Their Alloys, Nuclear Metallurgy;
Arsenault, R. J., Ed.; National Bureau of Standards:

Gaithersburg, MD, 1973; Vol. 18, p 176.
Kramer, E. A.; Johnson, W. L. Appl. Phys. Lett. 1979, 35,
815.
Motta, A. T. J. Nucl. Mater. 1997, 244, 227.
Howe, L. M.; Phillips, D.; Motta, A. T.; Okamoto, P. R.
Surface Coatings Tech. 1994, 66, 411.
Harris, L. L.; Yang, W. J. S. Radiation-Induced Changes in
Microstructure; ASTM: Philadelphia, PA, 1987; ASTM STP
955, p 661.
Gerling, R.; Wagner, R. In Proceedings of Fourth
International Conference on Rapidly Quenched Metals,
Sendai, Japan, Aug 24–28, 1981; Masumoto, T.,
Suzuki, K., Eds.; The Japah Institute of Metals: Sendai,
Japan, 1982; p 767.
Gerling, R.; Schimansky, F. P.; Wagner, R. Scripta Met.
1983, 17, 203.
Gerling, R.; Wagner, R. J. Nucl. Mater. 1982, 107, 311.
Rosinger, H. E. J. Nucl. Mater. 1980, 95, 171.