ORNL-TM-2258.txt

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OAK RIDGE NATIONAL LABORATORY
operated by

 

 

UNION CARBIDE CORPORATION
NUCLEAR DIVISION e
for the

U.S. ATOMIC ENERGY COMMISSION

ORNL- TM-2258
(Ut~ CTEOZ T >

MASTER

IRRADIATION BEHAVIOR OF CLADDING AND STRUCTURAL MATERIALS

J. R.Weir, J. O, Stiegler, and E. E. Bloom

NOTICE This document contains information of a preliminary nature
ond was prepared primarily for internal use at the Oak Ridge Naticnal

Laboratory. it is subject to revision or correction and therefore does
not represent a final report.

DISTRIBUTION OF THIS COCUMENT 1§ UNLIMITED
— it = LEGAL NOTICE e —o o oo o

This report was prepored as an account of Government sponsored work., Neither the United Stotes,

nor the Commission, nor any person acting on behalf of the Commission:

A, Makes any warranty or representation, expressed or implied, with respect to the accuracy,
completeness, or usefulness of the informotion contained in this report, or that the use of
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B. Assumes any liabilities with respect to the use of, or for damages resulting from the use of

any information, apparatus, method, or process disclosed in this report.

As used in the ahove, ''person octing on beholf of the Commission® includes any employee or
contractor of the Commission, or employee of such contractor, to the extent that such employee
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provides access to, any information pursuant to his employment or contract with the Commission,

or his employment with such contracter.

 

f"
ORNL-TM-2258

Contract No. W-7405-eng-26

METALS AND CERAMICS DIVISION

'LEGAL NOTICE
This report was prepared as an account of Government sponsored work. Neither the United
States, nor the Commission, nor any person acting on behalf of the Commission:

A. Mazkes any warranty or representation, expressed or implied, with respect to the accu-
racy, completeness, or usefulness of the information contained in this report, or that the use
of any information, apparatus, method, or process disclosed in this reporti may not infringe
privately owned rights; or

B. Assumes any liabilities with respect to the use of, or for damages resulting from the
use of any information, apparatus, method, or process disclosed in this report.

As used in the above, ‘*person acting on behalf of the Commission™ includes any em-
ployee or contractor of the Commission, or employee of such contractor, to the extent that
such employee or contractor of the Commission, or employee of such contractor prepares,
disseminates, or provides access to, any information pursuant to his employment or contract
with the Commission, or his employment with such contractor.

IRRADTATION BEHAVIOR OF CILADDING AND STRUCTURAL MATERIALS

J. R. Weir, J. 0. Stiegler, and E. E. Bloom

Paper presented at the American Nuclear Society, National

Topical Meeting, Cincinnati, Ohi -
. io, April 2— .
in the proceedi;gs, ’ > April 24, 1968. To be published

SEPTEMBER 1968

OAK RIDGE NATTONAL LABORATORY
Cek Ridge, Tennessee
operated by
UNION CARBIDE CORPORATION
for the
U.S. ATOMIC ENERGY COMMISSION

i T E L e e el o T RRAR RGO TR DR Pan p Rl
CROTRIEL TN T T DOCUMUNE R U R

A
iit

CONTENTS

Page

Abstract . . .« . 4 4 4 e e e e e e e e e e e e 1
Intrecduction . . . . « .+« v o v 4 00 e e e e e 2
Production of Defects . . . +« ¢« .+ « « ¢ v v v v 0 0w o0 3
Effects of Irradiation on Mechanical Properties 7
Low Temperatures 8
Intermediate Temperatures . . . +« « « « v « « « « « « « « « 18

High Temperatures . . « v « v o o « & + o & o ¢ o « o « o« « 29
SUMMETY  « v & o+ v o v e e e e e e e e e e e e e e e e e ... 38
Acknowledgments . . v v v 4 4 v e e e e e e e e e e e e e 4D

ReFETeNCES v v v v e e e e e e e e e e e e e e e e e e e e e e 4
IRRADTATTON BEHAVIOR OF CLADDING AND STRUCTURAL MATERIALS

J. R. Weir, J. 0. Stiegler, and E., E. Bloom

ABSTRACT

The effects of irradiation on the mechanical and
physical properties of materials to be used as cladding
and structural components in fast reactors are of great
interest to the reactor designer. In this paper the
general aspects of the problem are discussed in terms
of the observed changes in properties and micro-
structure and the possible mechanisms that might
explain the observed effects. The discussion is
concerned primarily with the austenitic stainiess
steels and with changes in mechanical properties
which occur at test temperatures near the irradiation
temperatures. For convenience the problem is divided
into three ranges of irradiation temperature: Ilow
temperatures, T < 0.40 Ty; intermediate temperatures,
0.40 Ty < T < 0.55 T; and high temperatures,

T > 0.55 Ty. (Tp i1s the melting point on the absolute
temperature scale.) On the basis of data presently
available the damage appears to be significantly
different for each temperature range. In the
low-temperature range there is an increase in yield
strength and reduction of work-hardening coefficient
and uniform strain. These effects result primarily
from the interaction of dislocations with irradiation-
produced defects. At intermediate temperatures
irradiation-produced changes in the precipitation
process become important. In this same temperature
range the formation of voids and dislocation lcops
after irradiation to high fast neutron fluences cause
large increases in yield strength and large reductions
in ductility parameters. At high-irradiation temper-
atures strength properties are not affected; however,
ductility is severely reduced. These effects result
from helium produced by wvarious (n,&) reactions.

 
INTRODUCTION

Changes in mechanical and physical properties of fuel cladding and
reactor structural components which occur as a result of neutron irradia-
tion are of major importance to the reactor designer. For example,
large reductions in either the strength or ductility of the material
used as a fuel cladding would severely limit its ability to withstand
the imposed stresses without excessive deformation or fracture. Mate-
rials used in a fast reactor system must retain adequate strength proper-
ties under rather severe operating conditions. The fuel cladding will
operate at temperatures between 400 and 700°C, will be exposed to fast

neutron fluxes of 1 X 10%% to 1 X 10%'%® neutrons cm™? sec™t

and during
its lifetime in the reactor will receive fast neutron fluences in excess
of 1023 neutrons/cmz. Other structural components may operate at some-
what lower temperatures and neutron fluxes but because of their longer
residence time in the reactor they may receive significantly higher
neutron fluences.

Data describing the effects of such irradiation conditions on the
mechanical and physical properties of materials are very limited. It
is thus necessary to combine the relevant data obtained from irradiations
conducted in thermal reactors with the data from fast reactor irradia-
tions in order to evaluate the expected changes in mechanical and
physical properties.

We shall restrict our discussion mainly to the behavior of
austenitic stainless steels and include results from other alloy systems
only to demonsirate general conclusions. This limitation is imposed

because the first liquid metal fast breeder reactors will be constructed
of these alloys and because the effects of irradiation on mechanical and

physical properties are best understood in these alloy systems.

PRODUCTION OF DEFECTS

Neutron irradiation of a crystal has two basic effects. First,
neutrons collide with lattice atoms and may displace some atoms. A
single displacement leaves one lattice site vacant, a vacancy, and
locates one atom in an off-lattice position, an interstitial atom. The
second effect, transmutation, is initiated by a neutron capture and
results in a changed mass number of the capturing atom.

Vacancies and interstitials are produced primarily as a result of
collisions between moving particles (neutrons or displaced atoms) and
lattice atoms. Assuming that such collisions can be treated as elastic
collisions between hard spheres, the maximum energy transferred when a

particle of mass m1 and energy E strikes a particle of mass mp at rest is

dmimo
E = e Hy 1
mex - (my + mp)2 (1)

Since the neutron has a mass number of 1, this becomes

G /Ay, (2)

where A> is the mass number of the struck particle. The average energy
transfer is half the maximum amount. Now, if the energy transfer to

the struck atom exceeds some threshold value, usually estimated to be
about 25 ev, the atom will be displaced from its lattice site. ©Such an
atom, termed a primary knock-on, will interact with lattice atoms in its
vicinity, possibly displace some of them, and gradually come to rest. If

the struck atom receives a large amount of energy, its more loosely bound
electrons will be stripped from it, leaving it highly ionized. Under
these conditions it will initially lose energy primarily through elec-
tronic interactions, but as it slows down it will make frequent colli-
sions with lattice atoms, the freguency increasing as the energy of the
knock-on decreases.

Caleulation of the total number of displaced atoms produced is
obviously a complex problem. To illustrate the order of magnitude of
the number we will follow the treatment of Kinchin and Pease.® They
assume that the knock-on loses energy entirely by ionization above some
cutoff energy approximately equal to the mass number of the struck atom
in thousands of electron volts and entirely by elastic collisions with
lattice atoms below this cutoff energy.

The number of additional displaced atoms produced per primary

knock-on atom is approximately

E
N, = EEE for 2B, < E<E, , (3)
and
Ei
Nd = 35 for E > Ei s (4)
d
where

E = the energy of the primary knock-on,
E. = the threshold displacement energy, approximately 25 ev
for metals,

E. = the energy of the primary above which it is assumed that

only ionization and no displacements are produced.
For example, if an iron atom (M = 56) is struck by a 1-Mev neutron

the maximum energy transmitted to the primary is [by Eq. (2)]

4 X 1

Emax T 56

 

~ (0,07 Mev .

This is above the ionization energy, so the number of displacements per

primary is [by Ea. (4)]

56,000 3 ..
— Sl A
Nd = 5% 5% 10° displacements.

It is important to realize that the displaced atoms are not produced
homogeneously throughout the material. For an individual collision the
defects reside in a small volume around the track of the primary knock-on,
which typically extends a few tens or perhaps hundreds of angstroms.

This volume is termed a displacement cascade, but in reality it may be
composed of subcascades produced by secondary knock-ons. Note too that
the distribution of vacancies and interstitial atoms within a cascade is
not uniform. In general, the interstitials are displaced outward, leaving
a vacancy-rich core in the center of the cascade.

Such regions are generally unstable and some dynamic recovery
occurs. The amount of recovery and the final configuration of the defects
depend critically on the irradiation temperature. At temperatures of
interest for normal reactor operation, both the interstitials and vacan-
cies have sufficient thermal energy to migrate through the lattice.

Many of the original defects are destroyed by recombination, trapping
at impurities, or absorption by dislocations and grain boundaries. Those
which survive cluster together to form stable configurations. At tempera-

tures in excess of approximately one-half the absolute melting point
(0.5 Tm) vacancies have sufficient thermal energy to overcome the
binding energy of clusters and to migrate freely through the lattice.
Thus at sufficiently high irradiation temperatures defects are
annihilated continucusly without cluster formation.

Transmutation reactions, in particular those which produce gaseous
species, may also have important effects on properties. Table 1 lists
the reactions and their approximate cross sections for a number of
important cases. We see that helium and hydrogen may be produced in
metals through neutron reactions both with impurities in the metals
and with the major alloying elements. Alter and Weber? have made calcu-
lations of the amounts of hydrogen and helium produced in various mate-
rials and concluded that for the iron- or nickel-base alloys used as
fuel cladding, approximately 100 ppm He and a few thousand parts-per-

million hydrogen would be produced in a fast reactor in a few years'

Table 1. Transmutation Reactions in Metals

 

Cross Section Neutron Energy

 

Nucleus Reaction a, Associated with
(barns) Cross Section
Lew (n,a) 41 Fission
10
B (n,x) 3800 Thermal
(n,o) 635 Fission
56 - .
Fe (n,oa) 0.35 Fission
(n,p) 0.87 Fission
8 (n,a) 0.5 Fission
(n,p) 111 Fission

 

al barn = 1072 cm?.
operation. In addition to these transmutation reactions producing
gaseous preducts, other possibilities exist in which solid impurities

are produced.

EFFECTS OF IRRADIATION ON MECHANICAL PRCPERTIES

Changes in mechanical properties produced by neutron irradiation
are a sensitive function of both irradiation and test variables. Impor-
tant irradiation variables include irradiation temperature, thermal
neutron fluence, fast neutron fluence, and possibly fast neutron flux.
Important test variables include test temperature and strain rate.

Other factors such as preirradiation heat treatment (in order to control

grain size, dislocation structure, and precipitate distribution) and time

at temperature (thermal aging) either before or following irradiation
have been shown to be important. Because of the large number of
variables and the vast amount of information which has been published
in this area we will not attempt a complete literature review. Rather
we will restrict our discussion to a general class cof metals and alloys
(those having a face-centered cubic crystal structure) and will be
concerned primarily with the mechanical properties at test temperatures
near the irradiation temperature. For convenience we define the following
temperature ranges: low temperatures, T < 0.40 Tm (where Tm is the
melting point of the alloy in degrees absolute); intermediate tempera-
tures, 0.40 T, <Tc< 0.55 im; and high temperatures, T > 0.55 Tm' Our
approach to the subject will be to summarize the observed changes in
properties, point out the important variables, illustrate changes in

microstructure and where possible correlate these changes with specific

mechanisms.
Low Temperatures

 

Tensile deformation of face-centered cubic metals at low tempera-

tures is usually terminated by a plastic instability, termed necking,

which leads to the development of a local reduced diameter region

followed by a shear fracture in this necked region.

limits the elongation of the material.

This local necking

The conditions under which this

instability occurs can be represented analytically.3 Assuming constant

volume and a power=-law relationship between true stress (E) and true

strain (¢) of the form

- —n
o = ke 3

(5)

where k is a constant, it can be shown that the plastic instability

occurs when the work-hardening exponent

i fn o
d in €

equals the true strain,

al{o|
oo
mIIQI

€ .

(6)

(7)

Figure 1 shows that Eq. (7) is reasonably well obeyed for type 304 stain-

less steel, but that n is not constant over the entire test.

When austenitic stainless steels are irradiated and tensile tested

in this low-temperature range, there is a large increase in yield stress

and large decreases in true uniform strain and work-hardening exponent.

by5

Figure 2 shows the room-temperature yield stress of type 304 stainless

steel after irradiation to 7 X 102° neutrons/cm?® (E > 1 Mev) and

9 x 10%0 neutrons/cm2 (thermal) at various temperatures. For irradiation

at temperatures between 93 and 300°C (approximately 0.35 Tm) the yield
(x103)

240

200

160

120

80

STRESS OR d5/d¢ (psi)

40

0.6

0.4

0.2

1, WORK HARDENING COEFFICIENT

 

   
     
 

 

T
/_——T
TRUE STRESS- !

ORNL-DWG 67-10786
1 T

T STRAIN

- . - ——————

~d5/de vs @

o

 

 

S
{~-ENGINEERING STRESS-

 

 

 

N

STRAIN

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

0.2

Fig. 1.

0.8
€, STRAIN

1.0 1.2 1.4

The Stress-Strain Characteristics of
Type 304 Stainless Steel at Room Temperature. The
arrows 1lndicate the strain at which the plastic
instability was observed to develop.

ORNL-DWG 66-5707

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

120
I /\ I ULTIMATE TENSILE STRENGTH
e . : (IRRADIATED)
110
100 . \
/\ \.\._____.
oo / \'
/
ap
70
YIELD STRESS
{IRRADIATED)
60
50 :
2
o /\.
o qo .
=1
W)
2
X 30
o
20 YIELD STRESS
{UNIRRADIATED)
10
o
o 100 200 300 400 500
IRRADIATION TEMPERATURE °¢C
Fig. 2. Room-Temperature Tensile Properties of

Irradiated Type 304 Stainless Steel.
10

stress was increased by approximately a factor of 3.

strain curves from this investigation are replotted in Fig. 3.

Typical stress-

For irra-

diation temperatures of 93 and 300°C the true fracture stresses and true

strains were approximately the same as the unirradiated specimen.

Values of engineering elongation were somewhat less for the irradiated

specimens.

After irradiation at 454°C the elongation has increased

again, but the fracture stress and strain were somewhat lower than in

the other tests, indicating that a different mechanism is operating at

454°C than at the lower temperatures.

Figure 4 shows that the work-

hardening exponents in the plastic range are consistent with the uniform

and total elongation values as predicted by Egs. (5) through (7).

(x 10°)

240

200

160

STRESS (psi)

80

40

120

ORNL-DWG 67-107892

 

 

Fig. 3.

 

I l

ENGINEERING STRESS-STRAIN

TRUE STRESS-STRAIN

— — UNCERTAIN PORTION TRUE

 

 

e
-
-

- ——

 

 

 

 

P -r
b -
STRESS-STRAIN : //‘,/
e
| T /4' e
| | AR
i i // J” - |
” — —
i ’ —_— “
S ,4{¢/ T |
o 7 1 j
2 | |
’ !
v/’/:l’ ”:}?/ !
- /S s »

 

   

UNIRRADIATED

 

| IRRADIATED AT 300°C

I !
— [RRADIATED AT 93°C

 

 

| 1\ ;
|
IRRADIATED AT 454°C

 

 

 

 

 

 

 

0.2 0.4 086

08
STRAIN

The Engineering and True Stress-Strain Curves for Type 304
Stainless Steel at Room Temperature, Tested in the Unirradiated Condition
and after Irrsediation at Various Temperatures.
11

ORNL-DWG 67-10790

 

0.8

T~

0.6 /'
T~ UNIRRADIATED

¥
/""“ ~ IRRADIATED AT 454°C
0.4 / l

! !
// ¥ />~ L IRRADIATED AT 93°C

/ =1 IRRADIATED AT 300°C

o2t

 

 

 

 

 

n, WORK HARDENING COEFFICIENT

 

 

 

 

 

 

 

0 0.2 0.4 0.6 0.8 1.0
€, TRUE STRAIN

Fig. 4. The Work-Hardening Cheracteristics Associated with the
Stress-Strain Curves Shown in Fig. 3., The arrows indicate uniform
strain as determined by point of maximum load.

Before examining the effects of neutron fluence, test temperature
et cetera, we should first consider the behavior in terms of microstruc-
tural changes and the interaction of dislocations with the irradiation-
produced defect clusters. At irradiation temperatures of approximately
350°C and lower "black spots" on the order of a few tens of angstroms in
diameter are observed in the microstructure of irradiated specimens. An
example of this type of damage for irradiation at 93°C is shown in Fig. 5.
At higher irradiation temperatures the spots have a larger size and

decreased density, as shown in Fig. 6. After irradiation at 371°C both

the spot density and yield stress (see Fig. 2) are decreased markedly.
 

 

 

 

 

 

 

 

 

 

 

 

Fig. 5.

 

 

Transmission Electron Micrograph of Type 304 Stainless

Steel Irradiated at 93°C. The black spots are defect clusters produced
by the irradiation.

 

 

- Fig. 6.

YE-9197

 

 

 Trénsmissicn'Eleéfif@n}flfifirdgraph of Type 304 Stainless

Steel Irradiated at 177°C. The spots are larger and more widely
distributed than those in the specimen irradiated at 93°C (Fig. 4).

 
 

 

 

 

 

 

 

v}

)

[ 3]

wy

13

Irregularly shaped pigggr-defects, probablyjfi}ecipitates, developed, but
these were widely enough spaced that they did not affect the yield stress.
At an irradiation temperature of 454°C the dot-like défect clusters were
completely absent. As shown in Fig. 7, there was extensive precipitation
at this temperature, including a heavy precipitate layer and an associated

denuded zone at the grain boundaries.

E

 

Fig. 7. Transmission Electron Micrograph Showing Precipitate
Particles Formed in Type 304 Stainless Steel During Irradiation at
454°C, Note the denuded zone adjacent to the boundary and the
extensive precipitation on the boundary.

“' These observationsiare~infgd¢& agreement with those of Armijo et al.®

who detected a dot-likeidgmagédlstructure in the same material irradiated

‘at 43 and 343°C to-fastunefit?éfiiflfienceé of 10?0 and 10%* neutrons/em?,

respectively. These authors report that the defects were considerably
larger in the specimen irféfiiatéd*to the higher fluence at thé‘higher

temperature.

 
14

Recent quantitative electron microscopy studies of irradiated face-
centered cubic metals have at various times claimed the dot defects to
be exclusively vacancy clusters and 1oops,7’8 interstitial clusters and
J_oops.,gilO or mixtures composed of small vacancy clusters and larger,
resolvable interstitial loops.llflz As these differences still have
not been resolved, we must at this point conclude that all can probably
be formed but that experimental circumstances (irradiation temperature,
flux, and fluence) determine the proportions in which each occur.

Transmission electron microscopy>> —©

of postirradiation deformed
single crystals of copper and molybdenum has shown channels in which the
radiation-induced defect structure has been eliminated. The interpreta-
tion is that glide dislocations sweep out or in some manner remove the
radiation-induced defects. The channels are generally clean except for
deformation-induced tangles and dipoles. The radiation defects are
completely eliminated from the channels®® and not simply pushed to the
edge of the channel, as was originally suggested.l3 Sharpl6 examined
annealed specimens containing channels and found no development of struc-
ture within the channels, as would be expected if they contained a high
density of point defects or peoint-defect clusters below the resolution
1imit of the microscope. The mechanism by which the moving dislocations
destroy the radiation-produced defects has not been determined. The
slip associated with the channels, determined by measuring the slip line
off'sets, corresponds to the passage of two or three dislocations on each

plane within the channel, so ample opportunity exists for dislocations

to remove all the defects present.
15

The channels gradually fill with tangles and deformation-induced
debris, through normal work-hardening processes, and this ultimately
halts deformation in the channels. Sharp16 observed a higher density
of debris existing on a smaller scale in the channels than in unirra-
diated material, but attributed this to the higher stress at which the
Slip band developed. During the latter stages of deformation the slip
line pattern of irradiated crystals appears similar to that of unirra-
diated materials.

Seegerl7 suggested that the defect clusters harden the lattice by
providing obstacles which moving dislocations must cut with the combined
aid of the applied stress and thermal fluctuations. As a result of this
chopping, the defects are gradually reduced in strength and ultimately
destroyed or eliminated by the dislocations, leading to the channels
that are observed.

Makin and Sharpl8 pointed out that in irradiated materials rela-
tively few slip lines are observed, indicating that few sources are
activated, that full-grown slip lines form dynamically in times of the
order of a millisecond, and that partially formed slip lines are not
observed. They proposed on the basis of elimination of the defects by
moving dislocations that the critical stress to form a slip band is the
stress required to operate a source in the environment of the defect
structure. ©Subsequent loops can be formed more easily, since the first
one clears a path for them. A pileup then forms and expands, creating
the cleared channel very rapidly at the high stress levels necessary to
generate the first dislocation. The result is creation of a soft zone
in a hardened material in which extensive localized shear occurs in a

short time until work hardening halts the deformation.
These observations provide a qualitative explanation for the
reduced work-hardening coefficients, increased yield stress, and low
uniform elongations in irradiated materials. The channeling produces
a soft zone in & very hard material, zones in which extensive slip
occurs. Because of the limited number of sources or slip systems the
dislocation tangling and interactions which normally lead to work
hardening occur more slowly and result in a reduced rate of hardening.
Figure 8 illustrates the narrow regions to which slip is confined in
stainless steel irradiated at 121°C and deformed lO% by rolling at room
temperature. The defect structure is still clearly visible in the
regions between slip bands. The magnification is not high enough to
reveal defect-free slip channels.

Within the low-temperature range changes in mechanical properties
are 'a function of fast neutron fluence and irradiation temperature.4519_2o
Figure 9 shows the effects of fast neutron fluence on the yield stress
and elongation for various irradiation temperatures. Note that the
increase in yield stress and reduction in elongation are greatest for
irradiation temperatures in the range of 160 to 290°C, but that differ-
ences do not develop until the material has recelved fast neutron
fluences of approximately 1 X 1020 neutrons/cmz. This suggests that
the defect clusters grow more complex with increasing neutron fluence.
Without further direct evidence one can only state in qualitative terms
that the importance of irradiation temperature stems from its influence
on the mobility of various defects. At the lowest temperatures vacancy
mobility is insufficient to allow the formation of vacancy clusters.

This is supported by the observations of Wilsdorf and Kuhlmann—Wilsdorle
 

 

 

 

 

 

 

 

 

 

n

17

 

Fig. 8. Transmission Electron Micrograph of Type 304 Stainless
Steel Irradiated at 121°C and Deformed 10% by Rolling at Room
Temperature. All the deformation has been confined to the dark bands;
the radiation-induced defect clusters can still be seen between the bands.
that no detectable defect clusters formed in type 304 stainless steel
irradiated at ambient reactor temperature to 10+ neutrons/cm2 and by
the observations of Bloom et al.’ that for irradiation at
93°C the defect clusters were small and showed extremely weak
contrast while at 121°C their size and contrast had increased
significantly.

The hature of the damage in'thé l6w-témperature‘range is apparently

unchanged at very hlgh fast neutron fluences. Cawthorne and Fulton??

,report that for an austenltlc stalnless steel 1rrad1ated to fast neutron

fluences of up to 5 X 1022 neutrons/cm at temperatures between 270 and

approx1mately 350°C "black spot" defects are present in- the mlcrostruc-

"ture. On postlrradlatlon anneallng the defects grow 1nto dlslocatlon

loops;‘ These loops flnally dlsappear on anneallng at about 700°C.

 
18

ORNL—DWG 67-10788R
1000

YS 160 °C

YS 290°C

@ 100
23 YS < 100°C
=z < YS 400°C
o = % £ <100°C
’_
< 0
o
55
oo

- 10

> + % £ 290 °C

— TOBIN, REF.19
o MARTIN AND WEIR, REF.4
O MURR e/ o/, REF. 20

}

 

108 10'® 10%° 10°! 1022 102>

NEUTRON EXPOSURE [neutrons Zcm2 (>1 Mev)]

Fig. 9. Room Temperature Properties of Annealed Stainless Steel
after Irradiation at Various Temperatures.

Intermediate Temperatures

 

Temperatures in the range of approximately 0.40 to 0.55 Tm (380 to
550°C for austenitic stainless steels) are particularly important to the
first generation fast breeder reactors. It is also in this temperature
range that radiation-damage phenomena are least understcod. Two
separate effects have been observed. The first involves precipitatiocn
and thus will be dependent on the alloy system. The second effect is
related to displacement processes and appears to be important at high
fast neutron fluences.

As discussed in the previous section, irradiation of type 304 stain-
less steel at 454°C to 7 X 10%C neutrons/em® (E > 1 Mev) resulted in

an increase of the room-temperature yield stress from approximately
 

 

 

 

 

)

)

.

-t

19

30,000 to approximately 43,000 psi and smaLl.reductions in the fracture
stress and strain.? EXamination of the microstructure of this specimen
revealed extensive precipitation, including a heavy layer along grain
boundaries. Unlike the defect clusters formed at lower temperatures,
such precipitates are not removed by dislocations but rather provide
permanent obstacles and sites for tangling. Deformation thus leads to
the tangled dislocation configurations shown in Fig. 10.

Arkell and Preil?3 showed that precipitate structures in a niobium-
stabilized stainless steel irradiated at temperatures between 450 and
750°C were significantly different than those present in unirradiated
samples with identical thermal histories. Irradiated samples exhibited

enhanced precipitation within the grains.

YE-9209

 

Flg 10. Transm1551ofi“Electron Micrograph of Type 304 Stainless

Steel Irradiated at 454°C and Deformed 10% by Rolling at Room
Temperature. Compare the unlform distribution of tangled dislocations
with the localized slip bands produced in specimens 1rrad1ated at a

lower temperature (Fig. 8).

 
20

“ reported the effects of irradiation temperature

Martin and Weir
on the postirradiation stress-strain behavior of types 304 and
347 stainless steel irradiated to 7 X 1029 neutrons/cm2 (E > 1 Mev) and
9 x 1020 neutrons/cm2 (thermal). For an irradiation temperature of
400°C an increased yield stress was observed for test temperatures up
to approximately 600°C. The strength increase for type 347 stainless
steel which contains approximately 1% Nb was significantly larger than
that which occurred in type 304 stainless steel (unstabilized). Since
niobium is a strong carbide former, it might be postulated that precipi-
tation processes are involved in the hardening mechanisms. |
More recently it has been observed??s 24727 tnat irradiation of
austenitic stainless steels at temperatures between 350 and 600°C to
high fast neutron fluences (> 1022 neutrons/cmz) results in large
changes in both properties and microstructures. Cawthorne and Fulton?2,23
used transmission electron microsceopy to examine the fuel cladding from
experimental fuel pins and tensile specimens irradiated in the

022

Dounreay fast reactor to neutron fluences up to 6 X 1 neutrons/cm2

at temperatures between 270 and 600°C. At irradiation temperatures

above approximately 350°C voids which varied in size from the smallest
resolvable to approximately 500 A were present. Voids constituted
1 to 2% of the volume of the material and could be eliminated by
annealing at 900°C.

Data obtained by Murphy and Strolm?® and Holmes Efi.fll°24 have
demonstrated that this type of damage causes large increases in the
yield strength and large reductions in ductility parameters.

Holmes gfi_g;.24 have correlated the changes in yield strength of
21

type 304 stainless steel irradiated at approximately 530°C to

1.4 X 10?2 neutrons/ecm? (E > 0.18 Mev) with the irradiation-produced
defect structure. The as-irradiated structure consisted of Frank ses-
sile dislocation loops, about 400 A in diameter and with a density of
3.7 X 10% loops/em®, and polyhedral cavities approximately 150 A in

ot4 cavities/cm3 in number. Figure 11 is a

diameter and about 2 X 1
plot of the yield stress (corrected for temperature dependence of the
shear modulus) as a function of test temperature. At test temperatures
less than 380°C the yield stress shows a thermally activated temperature
dependence. The athermal yield stress component is attributed to the
strengthening expected from the Frank sessile loops. Above approximately
538°C the sessile Frank loops transform to glissile perfect loops which
interact to form a dislocation network upon annealing at 593°C. Above
648°C the cavities or a combination of cavities and dislccation network
account for the athermal strength increases that persist to 760°C.

Full recovery of the yield strength was observed at 816°C where neither

the cavities nor the dislocation network was detected.

Murphy and Strohm®® have conducted tube burst tests on irradiated
EBR-II type 304L stainless steel fuel cladding following irradiation to
spproximately 1 X 10%% neutrons/cm® (fast). Over the length of the
cladding tube there is a temperature gradient such that the temperature
ranges from 370°C at the bottom to 500°C at the top. In addition, there
is a gradient in the neutron flux that ranges from about
1 X 10%% neutrons ecm™2 sec™t at the top and bottom to

1

2.5 X 10%° neutrons cm~? sec™  at the midplane. In tests at 500°C the

irradiated tubes exhibited a large increase in burst strength and large
22

(x10%) Y-86413
60

 

 

 

50 - - - -

 

|

40 . O |IRRADIATED

A UNIRRADIATED
o o |
30

20

 

 

G
G ROOM TEMPERATURE

 

 

 

 

 

 

Bt
b

YIELD STRESS (psi) x
(

 

 

10

 

[ D__

 

 

 

 

 

 

 

 

 

0 100 200 300 200 500 600
TEST TEMPERATURE (°C)

Fig. 11. Yield Strength (Proportional Elastic Limit) of AISI
Type 304 Stainless Steel after Irradiation in EBR-IT to
1.7 X 20?2 neutrons/cm® at 0.49 T,. (Ref. J. J. Holmes, R. E. Robins,
J. L. Brirhall, and B. Mastel, "Elevated Temperature Irradiation
Hardening in Austenitic Stainless Steels," accepted for publication
‘n Acta Metallurgica.)

 

reduction in ductility as measured by diameter increase at the edge of
the fracture. Figure 12 is a plot of ductility as a function of test
temperature. Between room temperature and approximately 600°C the
ductility is reduced to extremely low values, on the order of 1 to 2%.
At 700°C and above there is some recovery of ductility but the values
remain much lower than the unirradiated values. Irradiated specimens
which were given a pretest anneal at 900°C and then tested at 500°C
recovered all the preirradiation ductility and the strength was reduced

to that of the unirradiated tubes.
23

 

 

 

 

Y -86412
28
0
24— O UNIRRADIATED o
X [RRADIATED ( 0.5-1.4 x 10%2neutrons /cm2 RANGE ) o
20 O~
o 0
0
16— o
e ° 0
12— 0 0
0 0o 0O 0
a 0
g gl — X
X
X X
4 — % § g
X ¥ f
X 11 1 k%, 3 ¥
00

0
0 100 200 300 400 500 600 700 800 900 4,000

TEST TEMPERATURE (°C)

Fig. 12. Ductility of EBR-II Type 304L Stainless Steel Fuel
Cladding after Irradiation at Temperatures Between 375 and 500°C.
(Ref. W. F. Murphy and H. E. Strohm, "Tube Burst Tests on
Irradiated EBR-II Type 304L Stainless Steel Fuel Cladding,”" to be
published in Nuclear Applications, April 1968.)

 

Stiegler gfi_gl.27 have examined the type 304L stainless steel
cladding from a similar EBR-II fuel element. Two structural features,
namely voids and dislocation loops, were present in all specimens.

Table 2 lists the approximate irradiation temperatures, neutron fluences,
and void densities for each section examined. A comparison of the
results for sections 1 and 5 and 2 and 4 indicates that for the condi-
tions examined the void density decreases with increasing irradiation

temperature for a constant fluence.
24

Table 2. Irradiation Conditions and Void Density
Measurements for EBR-II Fuel Cladding

 

 

. Irradiation Fast Neutron Voids per
Section .
Number Temperature Fluence Cubic

(°c) (neutrons/cm®) Centimeter
X 1032 X 1075
1 370 0.8 1.4
2 398 1.2 1.3
3 438 1.4 1.3
4 465 1.3 0.9
5 472 0.9 0.4

 

Figure 13 shows a histogram of the void sizes observed in section 3.
On the basis of this void size distribution and the number of voids per
unit veolume listed in Table 2, it was calculated that the cladding den-
sity was decreased 0.17% by irradiation. Figure 14 shows examples of
voilds observed for three different irradiation conditions. It is readily
apparent that void size increases with increasing irradiation temperature.

The distribution of the voids was remarkably homogeneous. Varia-
tions observed between different micrographs probably reflect differences
in foil thickness. It is significant, however, that no volids were present
in the grain boundaries. In fact, the void density within about 0.1 B
of the boundary was reduced, probably by annihilation of volds contacting
the boundary or the influence of the boundary on the void-formation
process.

A very complex dislccation substructure was present in each of the

five sections., At the lower irradiation temperatures the structure was
25

Y-85139

 

 

40

 

 

(%)

 

 

30

 

 

 

 

25 L .1 Lo . oL - -

 

 

 

20 =

 

 

 

 

FREQUENCY OF OBSERVATION

 

 

10 e

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

I
0 50 100 150 200 250 300

0
VOID DIAM. (A)

 

Fig. 13. Void Size Distributions in EBR-IT Cladding Irradiated
at 438°C to 1.4 X 1022 neutrons/cm®.
 

 

 

 

26

YE-9150

 

. Fig. l4. Void Formation in Type 304L Stainless Steel Fuel Claddin
from EBR-II. (a) 0.8 x 1022 neutrons/cm® at 370°C, 1.4 X 1015 voids/c

§

m”,
(b) 1.4 x 102 neutrons/cm® at 438°C, 1.3 x 1017 voids/cm3.

(c¢) 0.9 x 10?2 neutrons/cm?® at 472°C, 0.4 x 10% voids/cm3.

W

wl

C.

4)

 
 

 

 

27

so complicated that irdividual loops could not be observed. At 472°C,
however, well-defined loops were resolved as shown in Fig. 15. These
loops lie on {111} and appear faulted, suggesting that they are Frank
sessile loops formed by the precipitation of interstitial atoms. The
loops ranged in diameter from 200 to 900 A and were present to a density
of about 2 X 10%% /em?.

Changes in microstructure as a result of postirradiation annealing
were examined for specimen 3. After 1 hr at 600°C the dislocation loops
disappeared and were replaced by a dislocation network. At progressively
higher annealing temperatures, the disloqation density decreased. After
1 hr at 900°C the dislocation density was comparable to that of an
unirradiated annealed specimen. Concurrent with changes in loop and

dislocation structure, the void density decreased. Measurements of

YE-9453

 

 
28

vold size distribution after annealing indicated that the smaller voids
annealed more rapidly. All voids were removed after annealing for 1 hr
at 900°C.

These observations allow a qualitative interpretation of the data
of Murphy and Strohm.?® The as-irradiated tubing contained voids and
dislocation loops which cause an increase in strength, possibly through

the mechanism as discussed by Holmes et al.?*

The recovery of properties
at 500°C as a result of postirradiation annealing at 900°C is a result

of the complete recovery of the damage. At 700°C and above the as-
irradiated structure recovers very rapidly; thus a partial return of
strength and ductility to unirradiated values is observed. It is impor-
tant to note that ductility is not completely recovered at test tempera-
tures in the range of 700 to 1000°C. The reasons for this will be
discussed in the next section.

The formation of voids and dislocation loops as a result of irradia-
tion to high fast neutron fluences not only causes large effects on
mechanical properties but also leads to swelling or a decrease in the
density of the material. Figure 1€ shows the correlation between the
density decrease and the fast neutron fluence for austenitic stainless
steels irradiated at temperatures between 370 and 560°C. It should be
noted that some of these data were obtained by direct density measure-
ments and some by calculations from void density and size measurements.
Several of the results were obtained from specimens removed from actual

fuel cladding and thus the material was subjected to stress during

irradiation and there can be litftle doubt that this will influence void
29

Y -B569C

;, 420 A | A440
! A400

A400 I
@438

i
#‘{r_
|
i

®— ORNL
A—PNL
—— B—Dounreay

FAST NEUTRON FLUENCE (neutrons/cm?)

IRRADIATION TEMP. ARE
—— SHOWN IN °C.

 

0.4 1.0 100
DENSITY DECREASE ( %)

Fig. 16. Summary of Stainless Steel Density Data.

growth. Interpretation of the data in terms of mechanisms is thus diffi-

cult. Figure 16, which is a summary of the available swelling data,22’25’27”28
does illustrate, however, that for some combination of temperature, stress,

and fluences in excess of 10°3 neutrons/cm2 volume increases greater than

10% may occur.

High Temperatures

 

At temperatures above 0.55 to 0.60 Tm’ the irradiation-produced
vacancics and interstitials are sufficiently mobile to allow continuous
recovery of defects during irradiation. It is still observed, however,
that when the iron- and nickel-base alloys are irradiated and then tested
at these high temperatures, there are severe changes in mechanical proper-
ties. These changes are characteristically different from those observed
at lower temperatures. In tensile tests the stress necessary to produce

a given amount of strain is unchanged; but irradiated specimens fail at
30

a straln much smaller than that at which an unirradiated specimen fails,
In creep tests the strain-time relationship is approximately the same for
irradiated and unirradiated specimens. Because of the reduced ductility,
however, the rupture life is significantly reduced. Examples of the
reduction in ductility and rupture 1ife?® in type 304 stainless steel
are shown in Figs. 17 and 18, respectively.

There are several important experimental observations which indi-
cate the nature and cause of the damage, Since neither the yield nor
35,30

ultimate tensile strengths are affecte and the ductility cannot

% it can

be reccovered by high-temperature postirradiation annealing,3
be concluded that neither displacement damage nor precipitation reactions
are the primary cause. Secondly, the loss of ductility is associated
with the grain-boundary fracture process and becomes more severe as the
test temperature is increased and the strain rate decreased. For ther-
mal reactor irradiations the postirradiation ductility is related to
the initial 9B content of the alloy and the thermal neutron fluence.>%,33
Boron-10 has a large cross section (3800 barns) for the “°B(n,a)7Li
reaction with thermal neutrons — each reaction producing a helium and
lithium atom. By cyclotron injection of helium and lithium ions into
an austenitic alloy Higgins and Roberts®* demonstrated that of these two
transmutation-producea isotopes, only helium had a large deleterious
effect on elevated-temperature ductility.

The most widely accepted model for the loss of elevated-temperature
ductility stems primarily from the work of Hyam and Sumner,35 Rimmer

and Cottrell,>® Cottrell,?” and Barnes’® and is summarized as follows.

The helium, which is produced from “9B(n,)7Li reactions with thermal
31

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

3
(x40 )50 Y -B46389
1 T
@ — UNIRRADIATED
® — IRRADIATED AT 650°C,
40 2.0 x10'? neutrons/cm? ( E>2.9Mev’)
2.0 x40%%neutrons/cm? { THERMAL)

35—————————~—«3>h~;;:::—-——————m——m~u-~_——Tw e
~ 30 e
‘»
Q.
- 25
[7)]
o
x
— 20
w

15

0 |

{0 100 1,000 10,000

RUPTURE LIFE (hr)

Fig. 17. Effect of Neutron Irradiation on the Rupture Life of
Type 304 Stainless Steel at 650°C. [Ref. E. E. Bloom, In-Reactor
and Postirradiation Creep-Rupture Properties of Type 304 Stainless

Steel at 650°C, ORNL-TM-2130 {March 1968).]
32

 

 

(%)

 

 

 

 

TOTAL ELONGATION

 

 

 

 

 

Y -86388
| ® — UNIRRADIATED
¢ — |IRRADIATED AT 650°C.
2.0 x10'®neutrons/cm? (E>2.9 Mev)
a0l 2.0 x10%heutrons/cm? { THERMAL)
30—
20 -
/AT
10
0

10 100 1,000 10,000
RUPTURE LIFE ( hr)

Fig. 18. Effect of Neutron Irradiation on the Total Elongation
&t Fracture of Type 304 Stainless Steel at 650°C. ([Ref. E. E. Bloom,
In-Rezsctor and Postirradistion Creep-Rupture Properties of Type 304
Smenless Steel st 650°C, ORNL-TM-2130 (Merch 1968).]
33

neutrons and (n,0) reactions between fast neutrons and most alloy
constituents, has a very low solubility in the matrix and precipitates
to form bubbles. When a normal stress (o) is applied to a bubble having

an initial radius (y) larger than a critical radius (rc) given by

r, = 0.76 y/o (8)

where ¥ is the surface energy, the bubble will become unstable and expand
indefinitely. Those bubbles which are located at grain boundaries can
lead to fracture initiation for several reasons:

1. Due to higher grain boundary diffusivities helium is supplied
to these bubbles and they can grow much faster than bubbles located in
the matrix.

2. As a result of grain boundary sliding, stresses may be concen-
trated at grain boundary jogs and triple grain Junctions; thus, a bubble
located in such a region will be subjected to a normal stress several
times the applied stress.

3. Once a grain boundary crack is formed, its rate of propagation
may be increased by the presence of grain boundary bubbles.

Figure 19 shows that the elevated-temperature tensile ductility of
type 304 stainless steel is a sensitive function of the total helium

72 These data were obtained from alloys containing various

concentration.
amounts of boron and irradiated to various neutron fluences.

Helium bubbles have been observed®9:%9% in both the matrix and the
grain boundaries after high-temperature irradiation. Figure 20 shows

helium bubbles in type 304L stainless steel irradiated at 700°C and

containing approximately 35 X 107° atom fraction helium. Rowcliffe et g;.BS
34

ORNL-DWG 67-10792

 

 

 

 

   
 
 

 

 

 

 

 

 

 

60 T e ! l
Y l L l | |
| | ||| 20.015ppm B 4 ‘
50 L | Jr' ° 0.1t ppm B -
| | 4045 ppm B l |

|

 

3.9 ppmB

|

T
{

 

 

 

 

 

 

 

  

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

1 1

10”7 Tol 10

 

TOTAL ELONGATION AT FRACTURE (%)

CALCULATED ATOM FRACTION OF HELIUM

Fig. 19. The Elongstion at Fracture of Type 304 Stainless Steel
Containing Varicus Amounts of Boron and Exposed to Four Radiation
Doses, Ranging from 1 X 1018 to 5 x 1020 neutrons/cm?. The elongation
is shown to kte a functicon of totzal concentration of hellum produced
oy both high-energy and thermal neutrons. The specimens were tested
at 700°C after irrediation at 50°C in the Oak Ridge Research Reactor.

cron concentrations (in parts per million): open triangle, 0.015;
oper. cirele, O.11; solid triangle, 0.15; solid circle, 3.9.
 

35

YE-9438

 

3

 

 

 

Fig. 20. Helium Bubbles in Type 304L Stainless Steel Irradiated
at 700°C Containing Approximately 35 X 10”6 Atom Fraction Helium.

|
|
|
|
|
|
I
|
|

 

 

 
36

have observed the growth of helium bubbles under stress at 750°C as would
be predicted by Eq. (8).

For material irradiated in thermal reactors, the distribution of
helium bubbles is controlled primarily by the initial boron distribution.
Woodford gz_gi.4l have observed halos of bubbles around precipitate
particles in a precipitation-hardening austenitic stainless steel,
indicating that boron is contained within these precipitates. In this
case there was a reduction in both ductility and creep rate., The
reduced creep rate resulted from the pinning of dislocations by bubbles.
The same effect could also be responsible for the reduced ductility.

It is important to note that for irradiations conducted in fast
reactors in which the thermal flux is essentially zero most of the
helium will be produced as a result of (n,a) reactions between fast
neutrons and nearly all alloy constituents. Under these conditions the
initial helium distribution will be nearly homogeneous. It has been

4 2

shown by King and Weir and Kramer g£'§£.43 that homogeneous helium

distributions produced by injecting & particles into type 304 stainless
Steel cause reductions in elevated-temperature ductility similar to
those observed after irradiation in thermal reactors.

The observations by Murphy and Strohm®® that even at high test tem-
peratures the ductility of irradiated EBR-II fuel cladding is not
recovered suggests that helium is responsible. This is consistent with
the fact that strength properties are essentially the same as those of
unirradiated tubing.

The elevated-temperature embrittlement problem has been found to be

a function of structural and compositional variations. Decreasing the
37

grain size or producing grain boundary precipitates by preirradiation
aging treatments give significant improvements in the postirradiation
tensile and creep-rupture ductility of type 304 stainless steel.30

These effects are believed to be due to the decreased tendency for inter-
granular fracture as a result of the increased stress necessary to
nucleate and propagate grain boundary cracks.

Roberts and Harries** found that the postirradiation tensile ductil-
ity of a 20% Cr—20% Ni niobium-stabilized austenitic stainless steel was
significantly improved by aging 100 hr at 750°C before irradiation.
Again, the results were interpreted in terms of the effects of grain
boundary precipitates on the formation of "wedge" type cracks during
testing. It was also shown in this investigation that the magnitudes
of the postirradiation ductility in a 18% Cr—10% Ni niobium-stabilized
alloy decreased with increasing boron content up to 50 to 70 ppm (weight)
and are then partially recovered in alloys containing higher boron
contents,

Titanium additions of approximately 0.2 wt % give significant
improvement in the postirradiation tensile and creep-rupture ductility
of types 304 and 304L stainless steel.

Figure 21 compares the postirradiation creep-rupture ductility of
types 304, 304L, and 304L + 0.2% Ti stainless steels at test temperatures
of 650 and 700°C. This effect is believed to be a result of (1) the
decreased tendency of the titanium-mcdified alloy to fracture inter-
granularly, possibly as a result of the redistribution of elements such

as nitrogen, oxygen, et cetera, (2) the segregation of boron into
38

 

 

 

 

 

 

 

 

/o

 

 

 

 

 

 

Y -8411%
60
.
—
50 /
/
40 ///
B
MATERIAL ANN.TEMP. TEST TEMP.
c B—M00.304L 900°C 700°C
o 30 A —M0D.304 L 1038°C 700°C
° a O— 304 1038°C 650°C
g ®— 304 1038°C 700°¢C
o A ¢®— 304L 1038°C 700°¢C
20
A/,
10
O @
—_— ¢
o N
l o I g
|
10 50 100 500 {000 5000 10000

 

 

 

Rupture Life Hours
Fig. 21. Cormpzarison of the Postirradiation Creep-Rupture Ductility
of Types 304, 304L and 304L + C.2% Ti Stainless Steels. ALl specimens
were irrsdiated at 650°C to 1029 to 10°% neutrons/cm® (thermal).

precipitates, thus reducing the amount of helium produced in the grain

boundaries, and (3) a refinement in grain size.

SUMMARY

Materials selected for use as cladding and structural components
in a fast reactor system will operate over a wide range of temperature,
neutron flux, and stress conditions. The changes in mechanical and
physical properties which occur as a result of neutron irradiation are
a. function of many variables, the most important of which appear to be
irradiation temperature and neutron fluence. With regard to the

austenitic stainless steels it appears that all known forms of damage
39

may occur. In components which operate at the lower end of the tempera-
ture range {below approximately 380°C) the work-hardening coefficients
and uniform elongations will be reduced. How severe these effects will

0%2 neutrons/cm2 is unknown.

be at fast neutron fluences in excess of 1

Damage may take on several forms in the temperature range 380 to
approximately 600°C. Changes in the precipitate distribution and
morphology have been observed. The ways in wnich these changes affect
mechanical properties are not entirely understood. Very recently the
formation of veids and dislocation loops as a result of irradiation in
this temperature range to fast neutron fluences in excess of approxi-
mately 107% neutrons/c.m2 has been observed. This damage causes drastic
reductions in ductility parameters and a density decrease or swelling
of the material. Available data suggest that this form of damage is
most severe at temperatures near 550°C and that for fast neutron
fluences of 1 X 1027 neutrons/cm2 density decreases as large as 10% may
occur.

At temperatures above 600°C cne would expect displacement damage
to be unstable and to recover in short times after it is created. Under
these conditions changes in strength properties are small. Reductions
in ductility which become more severe at higher temperatures and lower
strain rates are, however, observed. These effects are a result of

helium which is produced by various (n,a) transmutation reactions during

irradiation.
40

ACKNOWLEDGMENTS

The authors are happy to acknowledge the assistance of several
members of the Metals and Ceramics Division in the preparation of this
document: F. W. Wiffen and C. J. McHargue for technical review of the

manuscript, C.K.H. DuBose for the electron photomicrographs, K. W. Boling

for the graphic work, and M. R. Hill for preparation of the manuscript.
41

REFERENCES

1. G. H. Kinchin and R. S. Pease, "The Displacement of Atoms in

Solids by Radiation," Rept. Progr. Phys. 18, 1 (1955).

 

2. H. Alter and C. E. Weber, "The Production of Hydrogen and
Helium in Metals During Reactor Irradiation,” J. Nucl. Mater. 16, 68 (1965).
3. E. W. Hart, "Theory of the Tensile Test," Acta Met. 15, 351 (1967).
4, W. R. Martin and J. R. Weir, "Effect of Irradiation Temperature
on the Post-Irradiation Stress-Strain Behavior of Stainless Steel,”
p. 251 in Flow and Fracture of Metals and Alloys in Nuclear Environments

Spec. Tech. Publ, No. 380, American Society for Testing and Materials,

 

Philadelphia, Pa., 1965.

5. E. E. Bloom, W. R. Martin, J. 0. Stiegler, and J. R. Weir,
"The Effect of Irradiation Temperature on the Strength and Microstructure
of Stainless Steel,” J. Nucl. Mater. 22, 68 (1967).

6. J. 8. Armijo, J. R. Low, Jr., and V. E. Wolff, Radiation Effects

 

on the Mechanical Properties and Microstructure of Type 304 Stainless
Steel, USAEC Report APED-4542, General Electric Company, Vallecitos
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7. M. Wilkens and M. Ruhle, "Observations of Point Defect Agglomer-

ations in Copper by Means of Electron Microscopy,” The Nature of Small

 

Defect Clusters, AERE-R 5269, p. 365 (1966).

 

8. M. Ruhle, "Electronenmikros kopie kleiner fehlstellenagglomerate
in bestrahlten metallen," Phys. Status Solidi 19, 279 (1967).

9. K. G. McIntyre and L. M. Brown, "Diffraction Contrast from Small
Defect Clusters in Irradiated Copper,’ The Nature of Small Defect

Clusters, AERE-R 5269, p. 351 (1966).
42

10. K. G. McIntyre, "Interstitial Loops in Neutron Irradiated
Copper,” Phil. Mag. 15, 205 (1967).

11. M. J. Mekin, F. J. Minter, and S. A. Manthorpe, "The Correla-
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(1966).

12. G. P. Scheidler, M, J. Makin, F. J. Minter, and W. F., Schilling,
"The Effect of Irradiation Temperafiure on the Formation of Clusters in

Neutron Irradiated Copper,” The Nature of Small Defect Clusters,

 

AERE-R 5269, p. 405 (1966).

13. TI. G. Greenfield and H.G.F. Wilsdorf, "Effect of Neutron
Irradiation on the Plastic Deformation of Copper Single Crystals,’

J. Appl. Phys. 32, 827 (1961).

14, B. Mastel, H. E. Kissinger, J. J. Laidler, and T. K. Bierlein,
"Dislocation Channeling in Neutron Irradiated Molybdenum," J. Appl. Phys.
34, 2637 (1963).

15. J. L. Brimhall, "The Effect of Neutron Irradiation on Slip
Lines in Molybdenum," Trans. Met. Scc. AIME 322, 1737 (1965).

16. J. V. Sharp, "Deformation of Neutron Irradiated Copper Single
Crystals,” Phil. Mag. 16, 77 (1967).

17. A. Seeger, "On the Theory of Radiation Damage and Radiation
Hardening,"” Proc. U.N. Intern. Conf. Peaceful Uses At. Energy, 2nd,
Geneva 1958 6, 250 (1958).

18, M, J. Makin and J. V. Sharp, "A Model of *'Lattice Hardening®
in Irradiated Copper Crystals with the External Characteristics of

'Source! Hardening,"” Phys. Status Solidi 9, 109 (1965).
43

19. J. C. Tobin, M. S. Wechsler, and A. D. Rossin, "Radiation
Induced Changes in the Properties of Non-Fuel Reactor Materials,” Proc.
U.N. Intern. Conf. Peaceful Uses At. Energy, 3rd, Geneva 1964 9, 23 (1965).

20. W. E. Murr, F. R. Shober, R. Lieberman, and R. F. Dickerson,
Effects of Large Neutron Doses and Elevated Temperature on Type 347
Stainless, Battelle Memorial Institute, BMI-1609 (Jan. 23, 1963); also
F. R. Shober and W. E. Murr, "Radiation-Induced Property Changes in AIST
Type 347 Stainless,”" p. 325 in Symposium on Radiation Effects on Metals

and Neutron Dosimetry Spec. Tech. Publ. No. 341, American Society for

 

Testing and Materials, Philadelphia, Pa., 1963.

21. H.G.F. Wilsdorf and D. Kuhlmann-Wilsdorf, "Dislocation
Behavior in Quenched and Neutron Irradiated Stainless Steel,” J. Nucl.
Meter. 5, 178 (1962).

22. C. Cawthorne and E. J. Fulton, "The Influence of Irradiation
Temperature on the Defect Structure in Stainless Steel," The Nature of

Small Defect Clusters, AERE-R 5269, p. 446 (1966).

 

23. D. R. Arkell and P.C.L. Pfeil, "Transmission Electron
Microscopical Examination of Irradiated Austenitic Steel Tensile
Specimens,” J. Nucl. Mater. 12, 145 (1964 ).

24, J. J. Holmes, R. E. Robins, J. L. Brimhall, and B. Mastel,
"Elevated Temperature Irradiation Hardening in Austenitic Stainless

Steels," accepted for publication in Acta Metallurgica.

 

25. C. Cawthorne and E. J. Fulton, "Voids in Irradiated Stainless

Steel," Nature 216, 575 (1967).
44

26. W. F. Murphy and H. E. Strohm, "Tube Burst Tests on Irradiated
EBR-II Type 304L Stainless Steel Fuel Cladding,” to be published in

Nuclear Applications, April 1968,

 

27. J. 0. Stiegler, E. E. Bloom, and J. R. Weir, "Electron
Microscopy of Irradiated EBR~II Fuel Cladding," to be published.

28. J. J. Holmes, Battelle Northwest Laboratories, personal
communication (1968).

29. E. E. Bloom, In-Reactor and Postirradiation Creep-Rupture

Properties of Type 304 Stainless Steel at 650°C, Oak Ridge National

 

Laboratory, ORNL-TM-2130 (March 1968).

30. W. R. Martin and J. R. Weir, "Solutions to the Problems of
High-Temperature Irradiation Embrittlement," Spec. Tech. Publ. No. 426,
American Society for Testing and Materials, Philadelphia, Pa.

31. W. R. Martin and J. R. Weir, "Effect of Post-Irradiation Heat
Treatment on the Elevated Temperature Embrittlement of Irradiated
Stainless Steel,” Nature 202, 997 (1964).

32. A. C. Roberts and D. R. Harries, "Elevated Temperature
Embrittlement Induced in a 20 per cent Chromium: 25 per cent Nickel:
Nicobium Stabilized Austenitic Steel by Irradiation with Thermal Neutrons,"
Nature 200, 772 (1963).

23. W. R. Martin, J. R. Weir, R. E. McDonald, and J. C. Franklin,
"Irradiation Embrittlement of fiow-Boron Type 304 Stainless Steel,”
Nature 208, 73 (1965).

34. P.R.B. Higgins and A. C. Roberts, "Reduction in Ductility of

Austenitic Stainless Steel after Irradiation,” Nature 206, 1249 (1965).
45

35. E. D. Hyam and G. Sumner, "Irradiation Damage to Beryllium,"

Symposium held in Venice May 7—11, 1962, International Atomic Energy

 

Agency, Vienna, 1962, p. 321.

36. D. E. Rimmer and A. H. Cottrell, "The Solution of Inert Gas
Atoms in Metals,” Phil. Mag. 2, 1345 (1957).

37. A. H. Cottrell, "Effects of Neutron Irradiation on Metals and
Alloys," Met. Revs. 1, 479 (1956).

38. R. S. Barnes, "Embrittlement of Stainless Steels and Nickel-
Based Alloys at High Temperature Induced by Neutron Irradiation,' Nature
206, 1307 (1965).

39. A. F. Rowcliffe, G.J.C. Carpenter, H. F. Merrick, and
R. B. Nicholson, "An Electron Microscope Investigation of the High
Temperature Embrittlement of Irradiated Stainless Steels,” Spec. Tech.
Publ. No. 426, American Society for Testing and Materials, Philadelphia,
Pa.

40. D. R. Harries, "Neutron Irradiation Embrittlement of Austenitic
Stainlegs Steels and Nickel Base Alloys," J. Brit. Nucl. Energy Soc. 2,
74 (1966).

41. D. A. Woodford, J. D. Smith, and J. Moteff, "Distribution of
Boron in an Austenitic Steel Inferred from the Observation of Helium Gas
Bubbles after Neutron Irradiation,” J. Nucl. Mater. 24, 118 (1967).

42. R. T. King and J. R. Weir, "Effects of Cyclotron-Injected
Helium on the Mechanical Properties of Stainless Steel," paper presented
at the 13th Annual Meeting of the American Nuclear Society, San Diego,

California, June 11-15, 1966.
46

43. D. Kramer, H. R. Brager, C. G. Rhodes, and A. G. Dard,

Helium Embrittlement in Type 304 Stainless Steel, Atomics International

 

Report, NAA-SR-12601,

44, A, C. Roberts and D. R. Harries, "Effects of Irradiation and
Heat Treatment on the High Temperature Tensile Properties of Austenitic
Stainless Steels," Spec. Tech. Publ. No. 426, American Society for

Testing and Materials, Philadelphia, Pa.
1-3.
4=5.

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