ORNL-TM-2305.txt

 

OAK RIDGE NATIONAL LABORATORY
operated by

UNION CARBIDE CORPORATION
NUCLEAR Dl\”SION CARBIDE

for the
U.S. ATOMIC ENERGY COMMISSION

ORNL- TM- 2305

 

MASTER

 

POSTIRRADIATION TENSILE AND CREEP-RUPTURE PROPERTIES OF SEVERAL
EXPERIMENTAL HEATS OF INCOLOY 800 AT 700 AND 760°C

xy

NOTICE This document contains information of a preliminary nature
ond was prepared primarily for internal use at the Oak Ridge National
Laboratory. It is subject to revision or correction ond therefore does

not represent a final report.

WETRIUTION OF PéfS DOCUMERT & UM
v eem e mme e s oo LEGAL NOTICE - —mmmo e e e

This report was prepared as an account of Government sponsored work, Neither the United States,

not the Commission, nor any person acting on behalf of the Commission:

A, Makes any warranty or representotion, expressed or implied, with respect to the accuracy,
completeness, or usefulness of the information contained in this report, or that the use of
any information, apparatus, method, or process disclosed in this report may not infringe
privately owned rights; or '

B. Assumes ony liabilities with respsct 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 smpleyse or

contractor of the Commission, or employee of such contractor, to the extent that such employee

or controctor of the Commission, or employee of such contracter prepares, disseminates, or
provides access to, any information pursuant to his employment or contract with the Commission,

or his employment with such centracter.

 
ORNL~TM=-2305
Contract No. W-7405-eng-26

METALS AND CERAMICS DIVISION

POSTIRRADIATION TENSILE AND CREEP-RUPTURE PROPERTIES OF SEVERAL
EXPERIMENTAL HEATS OF INCOLOY 800 AT 700 AND 760°C

D. G. Harman

LEGAL NOTICE

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

A. Makes any warrunty 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 report 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 informaticn, 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 contracior prepares,
disseminates, or provides access to, any information pursuant to his employment or contract
with the Commission, or his employment with such contractor.

DECEMBER 1968

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

UFRIBULGY OF U KAAAALNS B UNUARERD

7
iii

CONTENTS

Page

Abstract & . . e e e e e e e e e e s e e e e e e e e e e .
Tntroduchtion .« & & v v v 6 e e e e 6 6 s 4 e s e e e e ..

Experimental Procedure . . . . « « c o s o e e e e e e e e e e

VoW

Results and Discussion . . . . . + + « &+ &
Effects on Ductility . .« & o o o v o« o v v o o o o o 0 . s 13
Alloy Composition . . .« & « ¢« v ¢ ¢ o« o « o o v . 13

Strain Rate . . . . ¢ v ¢ v v e e h e e e e 0 e e 17

Grain SizZe . .« v v 4 i 4 e e e e e e e e e e e e e 18
Preirradiation Aging . . . . « « « « « « o+ o . . . 21

Strength Considerations . . . . « « « « & « ¢ o « ¢ o o . 21
CONnClUSIONS + ¢« « ¢ o =+ o o o o o o o o o 2 s 4 s e e e e s 25

Acknowledgment . . .+ . « . 4 0 4 e e 4 e e e e e e e e e e e e 27
POSTIRRADIATION TENSILE AND CREEP-RUPTURE PROPERTIES OF SEVERAL
EXPERIMENTAL HEATS OF INCOLOY 800 AT 700 AND 760°C

 

D. G. Harman

ABSTRACT

Tensile and creep-rupture data have been obtained at
700 and 760°C for several experimental heats of Incoloy 8CO
that were irradiated in the ORR at elevated temperatures.
Effects of composition, grain size, and carbide morphology
were investigated.

Enhanced postirradiation ductility was achieved for
Tncoloy 800 containing about 0.1% Ti in creep-rupture tests.
The meximum ductility for this composition was obtained -
for the smaller grain sizes and at the lower creep stress
levels and appeared to be independent of carbon content.
Significant variations in properties {both control and
postirradiation tests) were noted for alloys within the
commercial Incoloy 800 composition specifications.

The ductility peak at about 0.1% Ti is not fully
understood; it might be best explained by two independent
mechanisms, one accounting for the increasing ductility
with increasing titanium at levels less than 0.1% and
the other explaining the decreasing ductility at higher
titanium levels, The grain size effect 1s thought to be
due to differences in either helium distribution or
stresses necessary for grain boundary fracture propagation.

 

INTRODUCTION

The elevated-temperature properties of Incoloy 800 make it an
attractive material for nuclear reactor application. The alloy has
been a backup material for fuel cladding for the BONUS (Boiling Nuclear
Superheat) Reactor® and is a prime candidate for various other reactor
systems. For example, Incoloy 800 is being considered for the LMFBR
(Liquid Metal Fast Breeder Reactor) fuel cladding, and Sweden is

 

1J. W. Arendt, BONUS Fuel Assemblies Progress Report No. 9,
Qak Ridge Gaseous Diffusion Plant, January 1966.
investigating vacuum-melted varieties for steam-cooled fast reactor
application.2

Incoloy 800 is nominally a 46% Fe—21% Cr—32% Ni ternary solid solution
alloy, but with important additions of carbon, aluminum, and titanium,
The commercial compositional specifications as listed in Table 1 allow
significant variations in the concentrations of these added elements.
Vendor recommendations for specific compositions depend upon the appli-

cation being considered.

 

M. Grounes, "Review of Swedish Work on Irradiation Effects in
Canning and Core Support Materials," pp. 200~223 in Effects of Radiation
on Structural Metals, Spec. Tech. Publ, 426, American Society for Testing
and Materials, Philadelphia, December 1967.

Table 1. Composition of Commercial Incoloy 800

 

Content, wt %

 

 

Element

Limiting Nominal
Iron Balance 46.0
Nickel 30-35 32.0
Chromium 19-23 20.5
Carbon 0.10 max 0.04
Manganese 1.50 max 0.75
Sulfur 0.015 max 0.007
Silicon 1.00 max 0.35
Copper 0.75 max 0, 30
Aluminum 0.15-0.60 .30
Titanium 0.15-0.60 .30

 
Other investigators3:4 have studied effects of high-temperature
neutron irradiation on the properties of Incoloy 800. Various titanium,
aluminum, and carbon levels were studied, but no comprehensive study of
compositional effects was undertaken. Also, postirradiation studies
included only short-time tensile testing. Little or no creep-rupture

data have been formally reported for irradiated Incoloy 800. Limited
preliminary creep data are currently available, however, from studies at
Studsvik, Sweden.” The present report shows that obtaining postirradiation
creep-rupture properties is essential to the evaluation of the Incoloy &CO

alloy system.

EXPERIMENTAL PROCEDURE

We tested several experimental 100-1b heats of Incoloy 800 as
listed in Table 2. Titanium contents range from less than 0.02 to 0.4%,
and two carbon levels are being studied — low carbon with 0.02 to 0.04%
and high carbon with 0.10 to 0.14%,

Control and irradiated specimens of the buttonhead design (See
Fig. 1.) used in previous ORNL experiments were tested at 700 and 760°C
under tensile and creep conditions. The tensile tests for both control
and irradiated specimens were conducted on a floor-model Instron testing
machine at crosshead speeds of 0.05 and 0.002 in./min. (Strain rates were
5 and 0.2%/min.) The creep-rupture tests for the control specimens were

conducted in air on dead-load and lever-arm creep frames., Irradiated

 

3C. N. Spalaris, Incoloy-800 for Nuclear Fuel Sheaths. A Monograph,
GEAP-4633 (July 1964).

“T. T. Claudson, "Effects of Neutron Irradiation on the Elevated-
Temperature Mechanical Properties of Nickel-Base and Refractory Metal
Alloys," pp. 67-94 in Effects of Radiation on Structural Metals, Spec.
Tech. Publ. 426, American Society for Testing and Materials, Philadelphia,
December 1967.

’Aktiebolaget Atomenergi, Stockholm, Sweden, private communications.

 
Table 2. Chemical Composition of Experimental 100-1b
Vacuum-Melted Heats of Ineoloy 8002

 

Element Content, wt %

 

 

Boron
Heat Ni or C Ti AL Mn Content

(ppm)
208 31 21 0.03 0.10 0.21 0.6 l
258 30 20  0.03 0.21 0.22 0.4 2
29¢ 29 19 0.03 0.28  0.28 0.5 2
33D 29 18 0.02 0.31 0.21 0.4 5
456 32 19 0.10 < 0.02  0.21 0.8 3
93H 30 21 0.12 0.10  0.24 0.8 6
411, 28 19 0.14  0.17 0.28 0.6 7
543 30 20 0.12 0.26  0.21 0.7 /s
61K 32 21 0.12 0.38  0.21 0.6 6

 

"Not listed: Si, 0.2%; V, Co, Nb, and Cu, < 0.05%; Fe, balance.

ORNL-DWG 67-3013

 

 

 

 

 

 

 

 

£e £ c &
88 8 on
g8 g8 g8
oo o0 QQ
+ 1 ¥ 90
@K f+s] !
o= s @
RE o 8%
S e 58
| | T
. | c]!
T I i o's
. . .4 3 LQ_S"
1 1 1 NIiD
. Sy
5 |
\, | !
5 % '
+0.0005 i
— 01875 R(TYP)
: |
e Yin ——— 1125 g —— Il
F
- — e T in, e

Fig. 1. Tensile Specimen.
specimens were creep-rupture tested in air on lever-arm creep machines
specially designed for hot-cell coperation. The cell during operation is
shown in Fig. 2.

Specimens tested at 760°C were irradiated in a poolside facility of
the ORR for one cycle (approx 1100 hr) at 760°C to a total fluence of 2 to
3 x 10°0 neutron/cm2 thermal and 1 to 2 x 1029 neutrons/em?® fast. From
cach of six compositions, eight specimens were irradiated — four having
been annealed at 1150°C for 10 min and four annealed and then aged 100 hr
at 800°C. This aging treatment was designed to agglomerate the grain
boundary carbides. To study the effect of strain rate, two tensile tests
and two creep-rupture tests were conducted for each of these two metal-
lurgical conditions for each composition.

Specimens tested at 700°C were irradiated at 650 or 700°C for two
cycles in a core position of the ORR to an average thermal and fast
fluence of about & x 1020 neutrons/cmz. Various metallurgical conditions
were investigated at this temperature. Grain diameters of 15 and 30 p
for the low-carbon alloys and 10 and 40 u for the high-carbon alloys were
included. The effect of the "carbide agglomeration" aging treatment on
cold worked and annealed specimens and the irradiation of cold worked
material were studied for both carbon levels. This aging treatment
actually recrystallized the ccld worked material to a grain diameter of
about & . These specimens were primarily creep-rupture tested, but a
limited amount of tensile testing was also conducted. Tests on control

specimens for this experiment are in progress and will be reported later.

RESULTS AND DISCUSSION

The results of the tensile and creep tests at 760°C are listed in
Tables 3 and 4. The tensile test results showed the expected loss in
ductility due to irradiation for all titanium levels for both the high-
and low-carbon alloys. Total tensile elongations of 40 to 70% were
reduced to 5.5 to 13.5% by irradiation.

With one exception similar ductility losses were noted for the creep-

rupture tests. Total creep elongations of 22 to 75% for the control tests
 

 

 

 

 

 

Fig. 2 Postirradia.tion Creep-Ru:pture Testing Facility. Pwelve lever=-arm creep machines within
this hot cell provide for remote long-time testing of irradiated materisls. Creep strain is monitored
manually from the dial gages and sutomatically with the LVDT transducers located at the lower end of
the stringer assembly.

 

 

 
Table 3. Tensile Properties of Experimental Incoloy 800 at 760°C™

 

 

 

 

 

 

 

 

Preirradiation Strain ?ate Strength, psi Strain, %

Condition (min=*) .

0.2% Yield Ultlgate Uniform Total

Tensile
x 103 x 10°

Heat 25B, 0.21% Ti, 0.03% C
Annealed 0.05 12.50 (17.52) 35,65 (34.71) 10.1 (16.7) 11.5 (52.1)
Annealed 0.002 12.80 (20.97) 22,49 (22.28) 5.2 (6.2) 8.2 (57.7)
Aged 0.05 12.56 (24.73) 32.00 (36.24) 7.4 (11.9) 7.9 (51.0)
Aged 0.002 13.65 (13.15) 21,91 (22.19) 2.1 (8.6) 5.8 (71.0)

Heat 33D, 0.31% Ti, 0.02% C
Annealed 0.05 8.02 (18.64) 38.02 (37.61) 9.7 (13.0) 11.5 (51.5)
Annealed 0.002 13.49 (17.25) 23.71 (22.24) 6.2 (6.9) 13.5 (61.0)
Aged 0.05 13.30 (16.82) 35,41 (37.22) 10.0 {12.7) 11.4 (54.8)
Aged 0.002 8.76 (12.55) 25.12 (21.49) 6.9 (7.3) 10.5 (57.8)

Heat 45G, <0.02% Ti, 0.10% C
Annealed 0.05 17.66 (18.93) 37.40 (39.51) 8.3 (13.4) 11.6 (38.9)
Annealed 0.002 18.44 (18.90) 23.03 (24.65) 8.0 (6.5) 10.2 (41.7)
Aged 0.05 22,10 (28.70) 36.85 (39.17) 7.7 (9.2) 11.1 (34.7)
Aged 0.002 18.52 (15.34) 23.34 (23.54) 6.5 (7.1) 9.5 (37.5)

Heat 93H, 0.10% Ti, 0.12% C
Annealed 0.05 19.60 (18.02) 35.50 (36.94) 9.9 (14.2) 12.8 (45.2)
Annealed 0.002 17.38 (17.88) 23.11 (24.94) 7.4 (7.6) 11.6 (55.9)
Aged 0.05 21.75 (18.27) 35.60 (38.44) 7.0 (10.8) 8.7 (47.7)
Aged 0.002 17.17 (17.84) 22.08 (24.53) 8.4 (10.0) 12.8 (~53)
Table 3. (continued)

 

 

 

 

 

 

Preirradiation Strain ?ate Strength, psi Strain, %
Condition (min™ ") .
0.2% Yield Ultimate Uniform Total
Tensile
x 107 X 103
Heat 41L, 0.17% Ti, 0.14% C
Annealed 0.05 17.59 (14.83) 37.30 (36.67) 8.1 (13.5) 10.3 (50.2)
Annealed 0.002 21.76 {25.59) 23.48 (25.84) 5.0 (2.3) 9.9 (50.4)
Aged 0.05 15.80 (16.35) 35.95 (39.25) 7.8 (7.0) 9.3 (50.0)
Aged 0.002 19.07 (21.75) 23.00 (23.71) 5.2 (6.5) 9.3 (41.7)
Heat 54J, 0.26% Ti, 0.12% C
Annealed 0.05 19.60 (14.91) 35.95 (34.64) 8.9 (12.1) 13.2 (54.1)
Annealed 0.002 17.51 (17.26) 23.32 (23.42) 6.9 (8.3) 11.2 (62.9)
Aged 0.05 16.89 (20.08) 35.95 {37.61) 7.7 (13.2) 10.2 (50.7)
Aged 0.002 19.75 (14.87) 23.11 (21.69) 5.7 (6.4) 9.2 (61.2)

 

aIrradiated one cycle in ORR poolside. Values in parentheses are

Specimens.

for unirradiated control
Table 4. Creep Properties of Experimental Incoloy 800 at 760°C%®

 

Hegt Contents, % Stress Time to Rupture Total Strain Preirradiation

 

(psi) (hr) (%) Condition
T4, C
x 10°
258 0.21 0.03 10.0 91 16.0 Annealed
8.5 481 17.0 Annealed
10.0 119 14.8 Aged
8.5 307 18.6 Aged
33D 0.31 0.02 10.0 179 14.8 Annealed
12.5 33 (113) 13.9 (63.9) Annealed
12.4 37 17.0 Aged
12.5 24 (101) 14.5 (39.2) Aged
45G <0.02 0.10 10.0 189 12.9 Annealed
10.0 260 15.2 Aged
93H 0.10 0.12 10.0 401 45 .4 Annealed
12.5 42 (132) 34,2 (41.3) Annealed
10.0 300 37 Aged
12.5 79 (199) 37  {(60) Aged
411 0.17 0.14 12.5 50 (200) 16.1 (69.8) Annealed
10.0 184 13.2 Annealed
12.5 32 (188) 8.7 (52.3) Aged
10.0 191 17.5 Aged
543 0.26 0.12 12.5 53.1 (181) 19.5 (55.7) Annealed
10.0 141 18.5 Annealed
12.5 46 (223) 13.6 (57.2) Aged
10.0 120 17.1 Aged

 

aIrradiated one cycle in ORR poolside. Values in parentheses are
for unirradiated control specimens.

were reduced to 9 to 20% by irradiation except for one heat of material
(heat 93H with 0.1% Ti), which showed only slight losses in creep
ductility.

Some loss in ultimate tensile strengths and an appreciable loss in
the rupture times-were observed; these were commensurate with the

ductility losses. Tensile yield strengths were relatively unaffected
10

by irradiation at this temperature. The results of both the control
tests and the postirradiation tests showed an insignificant effect of
the 100-hr aging treatment.

The results of tensile and creep tests at 700°C are listed in
Tables 5 and 6. Again, irradiation caused severe losses in ductility.
The control tests are still in progress and are expected to show

ductilities similar to those in the 760°C control tests.

Table 5. Postirradiation Tensile Properties of Experimental
Incoloy 800 at 700°Cc@

 

 

 

. ‘i Carbon Titanium Strength, psi Total
Material Condition Level Content Elongation
(%) 0.2% Yield Ultimate (%)
X 107 X 10°
Cold worked >50% Low 0.21 24,28 29.28 8.5
Cold worked, aged Low 0.21 22.02 25, 64 11.3
100 hr at 800°C
Cold worked, recrys- Low 0.21 15.26 32.24 6.7

tallized; 15-p
grain diameter

Cold worked, recrys- Low 0.21 13.49 33.61 8.4
tallized; 30-p
grain diameter

Cold worked, recrys- Low 0.10 20.69 28.54 10.2
tallized, aged 0.21 15.59 29.55 6.7
100 hr at 800°C; 0.28 16.35 31.56 6.7
30-u grain diameter 0.31 16.08 33.88 7.1

Cold worked, recrys- High <0.02 25 .47 32.29 7.1
tallized, aged 0.10 20.93 31.48 11.2
100 hr at 800°C; 0.17 19.22 30.99 11.1
40-p grain diameter 0.26 25.88 33.66 10.5

0.38 26.91 37.49 6.6

 

#Irradiated two cycles in the ORR core; strain rate 0.002/min.
11

Table 6. Postirradiation Creep Properties of Experimental
Incoloy 800 at 700°C?

 

 

Material Condition Carbon Titanium Rupture Total
Level Content ©Stress Time Elongation
(%) (psi)  (hr) (%)
x 107
Cold worked >50% Low 0.10 11.0 139 50,2
0.21 20.0 4 15.5
0.31 11.0 106 17.8
High <0.02 12.0 43 28.3
0.17 12.0 28 36.0
0.26 12.0 50 28.2
0.38 11.0 41 20.7
Cold worked, aged Low 0.10 12.0 70 50.7
100 hr at 800°C 0.21 15.0 17 26.2
0.31 12.0 48 11.5
High <0.02 15.0 12 17.7
0,10 12.0 24 49.0
0.17 15,0 11 23.7
0.38 15.0 14 14.3
Cold worked, recrys- Low 0.10 12.0 153 37.3
tallized; 15-u4 grain 0.21 15.0 32 11.9
diameter 0.31 12.0 300 4.8
Cold worked, recrys- High <0.02 12.0 20 19.0
tallized; 10-u grain 0.10 12.0 88 55.0
diameter 0.17 12.0 108 21.2
0.26 12.0 102 19.1
0.38 12.0 172 12.6
Cold worked, recrys- Low 0.10 12.0 117 33.7
tallized, aged 100 hr 0.21 15.0 29 12.0
at 800°C; 15~ grain 0.31 12.0 148 g.1
diameter
Cold worked, recrys- High <0.02 12.0 57 22.6
tallized, aged 100 hr 0.10 12.0 45 39.8
at 800°C; 10-m grain 0.17 11.0 63 23.9
diameter 0.26 12.0 61 25.2
0.38 12.0 21 18.1
Cold worked, recrys- Low 0.10 15.0 v 32.9
tallized; 30-i grain 0.21 15.0 99 15.3
diameter 0.31 15.0 123 9.2
12

Table 6. (Continued)

 

 

terial Conditi Carbon Titanium Rupture Total

Material Condltion Level = Content OStress Time  Elongation

(%) (psi)  (hr) (%)

x 10°

Cold worked, recrys- High <0.02 15.0 422 5.5

tallized; 40-u grain 0.10 15.0 167 22.7

diameter 0.17 15.0 67 10.1

0.26 15.0 326 5.4

0.38 15.0 696 5.5

Cold worked, recyrs- Low 0.21 20.0 13 12.0

tallized, aged 100 hr
at 800°C; 30-u grain
diameter

 

alrradiated two c¢ycles in the ORR core.

Tables 5 and 6 show that the postirradiation ductility is signifi-
cantly influenced by the grain size and composition of the material.
Increased ductility was noted for decreasing grain size. Enhanced
postirradiation ductility was observed for those heats having 0.1% Ti
(low-carbon heat 22A and high-carbon heat 93H). For example, the high-
carbon material with the 10-p grain diameter showed 54% postirradiation
creep elongation for 0.1% Ti but only 12% for 0.38% Ti heat. Similar
behavior was noted for the low-carbon heats. The preirradiation aging
treatment increased ductility except for the 0.1% Ti heats.

The results obtained so far illustrate the wide range in elevated-
temperature mechanical properties that can be obtained for material
‘within the Incoloy 800 specification. Significant differences were
noted in strength and ductility as measured in both irradiated and
unirradiated conditions. Also the preirradiation metallurgical condition
and the testing procedures (tensile or creep rupture) used may be
important considerations.

Although only a limited number of materials and testing conditions
have been investigated, some of the more important test results are
appropriate for discussion. The postirradiation ductility will be

treated first with the more important variables being discussed separately
13

as much as possible. The strength observations will then be discussed,
but to limited extent because of the scope of the report and the present

lack of 700°C control data.

Effects on Ductility

Alloy Composition

 

The test data collected on these 100-1b heats should indicate
desirable compositions for commercial large-scale heats of Incoloy 800
for nuclear applications. A similar scale-up approach has been used
for type 304 stainless steel® and is under way for type 316 stainless
steel” and the nickel-base Hastelloy N,® all modified by titanium
additions.

As has been the case for the other alloy systems, the specific
level of titanium has a significant effect on postirradiation properties
at elevated temperatures. The most dramatic effect is seen in the
ductility as measured in the creep-rupture test. The total elongations
are plotted against the titanium content in Figs. 3 and 4 for test
temperatures of 700 and 760°C, respectively. A decisive peak in creep
ductility is noted for both test temperatures at around 0.1% Ti, and
the position of this peak is not affected by the carbon level over the
range of 0.03 to 0.13% at 700°C.

A similar ductility peak was found by Weir and Martin®

for experi-
mental heats of type 304 stainless steel alloyed with titanium,as is
shown in Fig. 5. Since the formation of TiC is thermodynamically

favorable (the free energy of formation is —41,000 cal/mole at 1000°K),

 

6J. R. Weir, Jr., "Radiation Damage at High Temperatures,” Science
156, 1689—-1695 (June 1967).

7E. E. Bloom, Metals and Ceramics Division, ORNL, private communi-
cation.

8. E. McCoy and J. R. Weir, "Development of a Titanium-Modified
Hastelloy N with Improved Resistance to Radiation Damage,” Presented at
the Fourth International Symposium on Effects of Radiation on Structural
Metals, San Francisco, June 2628, 1968.

°W. R. Martin and J. R. Weir "Solutions to the Problems of High-
Temperature Irradiation Embrittlement,"” pp. 440457 in Effects of
Radiation on Structural Metals, Spec.Tech. Publ. 426, American Society
for Testing and Materials, Philadelphia, 1967. =

 
Y -20103
ORNL-DWG 68-10894

 

 
 
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

319 T T 60 T [
HIGH CARBON LOW CARBON
O ANNEALED AT 1038°C O ANNEALED AT 980°C
A ANNEALED AT 1150°C 50 /A ANNEALED AT 1038°C
— A ANNEALED AT 1150°C ] A ANNEALED AT 1038°C
AND AGED 100 HR AT AND AGED 100 HR AT
800°C 800°C
40 o
5 2
= = [I
o o 10 [
= ~ L -1
g g ,II CREEP
z 5 L
W w /1
20 ——-/—I
/1
\ /1
o jop ’,’
L IE—\—
40u TENSILE No
| |
| 0 |
0 04 0.2 03 04 0.5 0 04 0.2 0.3 04 05

Ti CONTENT (wt %)

Fig. 3. Postirradiation Ductility of E

Ti CONTENT (wt %)

xgerimental Heats of Incoloy 800 Irradiated in the ORR at

650 and 700°C t» About 0.8 x 10?! neutrons/cm? (Thermal) and Tested at 700°C.,

71
15

ORNL-DWG 68-10890

 

50 ‘
0 — 042 % CARBON

/ \ 0 —0.04% CARBON
40

CREEP

 

{%)

30

 

ELONGATION

 

 

 

\fl\i__ n/‘fl

 

 

 

 

 

 

 

 

 

o 04 0.2 0.3 c.4 0.5

TITANIUM CONTENT ( wt %)

Fig. 4. Postirradiation Ductility of Experimental
Incoloy 800 at 760°C. Alloys were annealed at 1150°C, then
irradiated at 760°C to 3 to 4 x 1020 neutrons/cm® thermal.

ORNL—-DWG 66— 239

 

 

   
  

 

 

80 T [ '
'STRAIN RATE OF 2%/ min E
70 e “ i ! : |
UI\EIRRADH-?TED ,
|
60 _IRRADIATED 1x102° neutrons/em®_|____ |

(THERMAL) |
15X 10'? neutrons/cm? (FAST)

a0 |- - :
1‘\‘

 

 

 

 

 

 

DUCTILITY, TOTAL ELONGATION (%)

 

 

P———s
30 ;
20 | — ]
)_—__H)\J>
10 r Ty T T

 

 

 

 

 

/RE'GULAR 304 SS
0 02 04 06 08 1.0 1.2 1.4
PERCENT Ti

0

 

Fig. 5. Ductility at 842°C of Irradiated Austenitic
Stainless Steels as a Function of Titanium Content.
16

it has been arguedlo that the ductility maximum occurs at the 1:1 atom
ratio of titanium and carbon (4:1 weight ratio). The argument is that
the formation of TiC utilizes all of the available carbon and prevents
the formation of the grain boundary embrittling M»3Cg. However, carbides
other than titanium carbide (e.g., M33Cq) have been seen'? in the
structure for all titanium levels. Indeed, the present study shows that
for Incoloy 800 the position of the ductility peak is independent of
carbon content over the range studied.

The radiation embrittlement of high~temperature alloys has been
attributed to the generation of helium by the 1°B(n,a) reaction.®
Titanium is a strong boride former — the free energy of formation of TiB
is ~38,400 cal/mole at 1000°K — and therefore may distribute the boron
and subsequent helium throughout the matrix (and away from the grain
boundaries) by the formation of TiB. This may account for the rise in
ductility with increasing titanium content at the lower titanium levels
but does not readily explain the decreasing ductility as the titanium
exceeds 0.1%.

Many compounds appear in the microstructure of Incoloy 800, as is
the case for most engineering alloys. Two competing processes could be
taking place. On cne hand, the formation of TiB could result in increased
ductility with increasing titanium; on the other hand, at higher titanium
levels the formation of other compounds could enhance grain boundary
embrittlement and decrease ductility. These could be titanium compounds
or other compounds whose formation is enhanced by the presence of
titanium.

This argument might account for the absence of a ductility peak for
Hastelloy N modified with titanium. A continuous increase in ductility
was noted® up to about 1.2% Ti. The thermodynamics of the Hastelloy N

alloy system might be such that at low titanium levels the grain boundary

 

10D, R. Harries, Atomic Energy Research Establishment, Harwell,
England, private communication.

¥, A. Comprelli and J. E. Lewis, Microstructural Evaluation of
Superheat Cladding Materials, GEAP-4751 (January 1965).
17

embrittling reaction described above would not occur. However, the post-
irradiation ductility of Hastelloy N seems to be insensitive to the
initial boron content up to 50 ppm, which questions the above effect of

TiB for this alloy.

Strain Rate

The enhanced postirradiation ductility at the 0.1% Ti level is shown
in the creep tests but not in the tensile tests, as can be seen in Figs. 3
and 4. Little or no compositional effect is shown for the tensile results,
whereas the creep-rupture data show a sizable effect on postirradiation
ductility.

Other studies'?s13 on the postirradiation ductility of both experi-
mental and commercial Incoloy 800, which used only tensile testing, did
not reveal any significant effects of composition. Our results show that
the reason for this was an unfortunate choice of postirradiation testing
conditions.

A model based on the stress-induced growth of helium bubbles that
are present in irradiated metals has been offered'” as an explanation
for the increasing postirradiation ductility at decreasing strain rates
(i.e., decreasing stress levels). A certain critical stress, given as
0.76 %, is required to cause the helium bubble present at a grain
boundary to grow and subsequently link with others to cause premature
intergranular fracture. 1In this expression r is the radius of the helium
bubble, which is assumed to be spherical, and y is the effective surface
energy. According to this model, the irradiated material would show the

same ductility as that unirradiated at sufficiently low stresses.

 

12¢. N. Spalaris, Incoloy-800 for Nuclear Fuel Sheaths. A
Monograph, GEAP-4633 (July 1964).

13p, T, Claudson, "Effects of Neutron Irradiation on the Elevated-
Temperature Mechanical Properties of Nickel-Base and Refractory Metal Alloys,
Alloys," pp. 67—94 in Effects of Radiation on Structural Metals, Spec.
Tech. Publ. 426, American Society for Testing and Materials, Philadelphia,
December 1967.

4R, S. Barnes, "Embrittlement of Stainless Steels and Nickel-Based
Alloys at High Temperature Induced by Neutron Irradiation,” Nature 206,
1307 (1965). =

 

=
18

Figure 6 shows this to be the case Tor the 0.1% Ti material at 760°C

(data from Tables 3 and 4). However, this does not explain why the

other heats of Incoloy 800 included in the figure responded only weakly

to decreasing stress levels., It could be that the surface energy term

in the critical stress relationship is less for these alloys than for those
with 0.1% Ti. Some very complex grain boundary reactions have been

found for Incoloy 800 and may be indicative of surface energy variability.
Recent findings by Snyderl5 on some of the alloys included in this report
show that these reactions are sensitive to titanium content. However,
Gehlbach*® found no titanium present in the grain boundary precipitates

of the one experimental alloy that he analyzed using the microprobe.

Grain Size

Grain size has been reportedl7 to be important to the postirradiation
ductility of type 304 stainless steel, with the finer grained material
more ductile. Figure 3 verified this for the experimental Incoloy 800
at 700°C., The same data are plotted as a function of grain diameter in
Fig. 7 for titanium levels of 0.1 and 0.2%. Under these testing conditions
the postirradiation creep ductility of Incoloy 800 with a given titanium
level is governed almost entirely by the grain size and seems to be
independent of carbon level. This is particularly important for
Incoloy 800, as laboratory tests show that the grain size can be
controlled over a wide size range for all compositions. Figures 8 and 9
show resulting grain sizes for representative heats of both low and high
carbon content annealed at various temperatures.

Martin and Weirl7 suggested that the effect of grain size on post-

irradiation ductility might result from differences in helium concentration

 

15y, B. Snyder, Graduate Student, the University of Tennessee,
unpublished work performed at ORNL in 1967.

Lér, m. Gehlbach, Metals and Ceramics Division,private communication.

7. R. Martin and J. R. Weir,"Solutions to the Problems of High-
Temperature Irradiation Embrittlement,"” pp. 440-457 in Effects of Radiation
on Structural Metals, Spec,Tech. Publ. 426, American Society for Testing
and Materials, Philadelphia, 1967. _

 

 
19

Y-85656
ORNL-DWG 68-10856

 

80

70 77T Tl
W ONTROL SPECIMENS
60— / %g /7\\
I > &9, |

 

 

 

 

 

 

   

 

 

 

 

 

TOTAL PLASTIC ELONGATION (%)

 

 

 

 

 

1072 10! 10° 0 402 102

STRAIN RATE (% /hr)

Fig. 6. Tensile and Creep Ductility of Experimental Incoloy 800 at
760°C. Alloys were solutzon annealed at 115C0°C. TIrradiation was conducted
at 760°C to 2 to 3 x 10°° neutrons/em?, while the control specimens were
soaked at 760°C for the duration of the irradiation.

ORNL-DWG 68-10889

 

60

0O-—HIGH CARBON
A—LOW CARBON

 

50 p— -

 

(%)

40 p—

30 e e meme e o m e

 

 

CREEP ELONGATION

20— -

 

 

 

 

 

 

 

0 10 20 30 40 50

APPROX. GRAIN DIAM. (p)

Fig. 7. 7Postirradiation Creep Ductility »f Experimental Incoloy 800
at 700°C. Alloys were annealed to obtain the indicated grain sizes and
irradiated at 650 and 700°C to about & x 1029 neutrons/ecm®. Creep tests
were conducted in air at 12,000 psi.
 

 

 

 

 

 

 

   

   
   

 

0.035 INCHES
A lw : im 100X

| Y-80471

 

  

Fig. 8. Grain Size of Experimental Incoloy 800. (a) Low-
~carbon material annealed at 980°C. (b) High-carbon material
annealed at 1040°C. ({c¢) Low-carbon material annealed at 1040°C.
(4) High-carbon material annealed at 1150°C. ‘

 

Y-84200

 

 

 

 

 

 

I = UL001/5 INCHES —
T e T T T io- T ~TgoGox T o [ T T T

. Fig. 9. Gréih'Size¥bfrExpérifienta1'Ihcoloy 800 Annesaled

at 800°C. (a) Low-carbon material. (b) High-carbon material.

 
21

on grain boundaries or from changes in grain boundary shear stresses that
cause fracture. ©OSince a given amount of helium is produced from boron
burnup and if this helium is present as bubbles in the grain boundaries,

a lower concentration at the boundaries would be expected if the total
available grain boundary area were increased (i.e., grain size decreased).
The lower helium bubble concentration would be expected to be less
embrittling to the grain boundary and thus improve overall ductility.
Even though not all of the helium may reach the grain boundaries the
overall conclusion of increasing ductility with decreasing grain size

would probably remain valid.

Preirradiation Aging

 

Some of the specimens received a heat treatment designed to
agglomerate the carbide that i1s present in these alloys. Carbide in
this form spaced along the grain boundary surface should be less detri-
mental to grain boundary ductility than the nearly continuous carbide
films that frequently result from conventional annealing treatments.
Figures 10 and 11 show the carbide agglomeration. Comparison of the
test results for aged and unaged specimens as listed in Tables 3 and 4
shows little difference in postirradiation properties for the 760°C
irradiation and test temperatures. EBEvidently the 1100 hr at 760°C during
the irradiation also provided sufficient carbide agglomeration. This
was also true for the control specimens, which also experienced the 1100 hr
at 760°C before testing.

Tables 5 and 6 list the 700°C test results, showing some effect of
preirradiation aging. The postirradiation ductility of the high-carbon
material was slightly lowered for the heat containing 0.1% Ti but was
increased for all other titanium levels., In effect this broadens the

ductility peak. The results were not conclusive for the low-carbon heats.

Strength Considerations

In general, the effects of irradiation on the strength of the

experimental Incoloy 800 were about those expected from date collected st

ORNL on similar materials and are in agreement with the Incoloy 800 findings
22

 

I-s

Iro

 

o

0.007 INCHER
15 500X

n

 

™

 

 

 

 

Fig. 10. Low-Carbon Incoloy 800 Annealed at 1040°C and Aged

100 hr at 800°C.

 

 

 

 

 

 

1000X

 

 

o

~hw

 

 

 

Fig. 11. High—Carbon'InCOloy 800 Annealed at 1150°C and Aged
100 hr at 800°C.

 

# Y-83690
23

reported in the literature. That is, the strength properties were
affected only through a curtailment of total ductility. Tables 3 and 4
list data for both the irradiated and the control specimens at 760°C.
Examination of the tensile results in Table 3 shows that the 0.2%
yield strength, the ultimate tensile strength, and the uniform elongation
are unaffected by irradiation. The stress-strain curve is affected by the
irradiation only in that portion beyond the peak load, where nonuniform
deformation and grain boundary fracturing are occurring. These occur
with less overall elongation for the irradiated specimens. Typical
tensile curves comparing the unirradiated and irradiated material at

760°C are shown in Fig. 12.

ORNL-DWG 68-10892

 

24

 

22

 

 

 

 

 

N
o
* 1
0
A
>
O
5
m
O
_.(’
m
w
—

 

o
!
i
|
|
i
i
1
|

 

o
! \
o

I
O
{
|
1
T

   

 

 

=
»
!

CONTROL TEST

 

D
|
|
|
E
l
1

 

STRESS (10°psi)

N
!

 

 

—
O

{

I

i
—

{

!

!

|

|

1

\

I

|

 

8 b—— &

 

 

 

 

 

 

 

 

 

 

|

| i X
| |

0 10 20 30 40 50 60 70

ELONGATION (%)

 

Fig. 12. Tensile Stress-Strain Curves from Irradiated and Control
Semples of Incoloy 800 containing 0.12% C and 0.26% Ti at 760°C and
0.002/min. After an 1150°C anneal, material was irradiated to
2—3 x 10°° neutrons/em® at 760°C or soaked at 760°C {control).
24,

Similarly, the initial stages of the creep curve are very nearly
the same for both irradiated and control specimens, as shown in Fig. 13,
Minimum creep rates were generally slightly higher for the irradiated
material. Bloom™® has attributed this to an early onset of third-stage
creep before a true minimum creep rate was established. However, the
second stage of creep lasted much longer for the unirradiated material

and more elongation occurred during third-stage creep.

 

18g, E. Bloom, In-Reactor and Postirradiation Creep-Rupture Properties
of Type 304 Stainless Steel at 650°C, ORNL-TM-2130 (March 1968).

ORNL-DWG 68-10891

60

 

 

 

 

ELONGATION (%)

 

 

 

 

 

 

 

 

 

{

|

|

| i
0 20 40 60 80 100 120 140 160 180
TIME (hr)

Fig. 13. Creep Curves from Control and Irradiated Samples of
Incoloy 800 containing 0.12% C and 0.26% Ti at 760°C and 12,500 psi.
After an 1130°C anneal, materisl was irradiated to 2-3 x 10éo neutrons /cm?
at 760°C or sosked at 760°C (controls).
25

The curtailment of creep ductility seriously affected the creep-
rupture life of Incoloy 800. Figure 14 illustrates the shorter rupture
times at 760°C for the irradiated material for the high-carbon heats
with 0.17 and 0.10% Ti. Included in these plots is the ultimate tensile
strength at two strain rates plotted against the total test time. The
higher titanium heat shows the nearly parallel rupture curves typical of
high-temperature alloys. However, the more ductile 0.1% Ti heat shows
the rupture curves converging at the lower stresses, This is expected
since the ductility curves for this heat also converge at these stresses
(see Fig. 6). It is also evident from the stress-rupture curves that
aging 100 hr at 800°C before irradiation had no effect on the 760°C
rupture life.

Any conclusions regarding rupture strength based on the 700°C data
must await the testing of the control specimens. However, we note that
the postirradiation rupture strength increases with grain size, as would
be expected at this temperature for unirradiated material. For example,
for the 0.1 and 0.2% Ti alloys stressed at 12,000 psi, increasing the
grain diameter from 10 to 40 u increased the rupture time from 100 to
500 hr, and this increase seemed to be independent of carbon content.
Aging the high-carbon alloys for 100 hr at 800°C before irradiation
reduced the rupture life at 700°C about in half. Also, the higher
titanium heats had longer rupture times, which could be due to an aging

effect.

CONCLUSIONS

The conclusions below are based on the test results contained in
this report and correlation with reported results on similar materials.
More meaningful conclusions regarding Incoloy 800 will undoubtedly be
forthcoming as additional material, irradiation, and testing variables
are investigated.

1. Enhanced postirradiation ductility can be achieved for
Incoloy 800. Maximum ductility is obtained for a concentration of
0.1% Ti, and this optimum titanium level is independent of carbon content

over the range of 0.03 to 0.12%.
ORNL-DWG 68-10893

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

40,000 = [ p——
® O ANNEALED AT 1150°C
30,000 - @ O ANNEALED AT 1150°C AND
= - AGED 100 HR AT 800°C
& m b 'T
m Lok . — DS SR
& 20,000 S . JT 'D TfiPOL t ‘f
o / _
t I S N R : €T§Q“ ~ \\\\‘\\q_ . ____MLH
\
10000 L1 T~
0.01 O 1.0 10 100 @ 1,080 5,000
RUPTURE TIME (hr)
40,000 [ — ORNL—]DWG 6l8-10895
— @ O ANNEALED AT 1150°C |
30,000 . @O0 ANNEALED AT 1150°C AND
?: | "‘\[@—— B AGED 100 HR ATY§OTO°C
Q. Iy .
; % | \\ C ; !
»n 20,000 - | > Omfig—mwu* <}
L — _ —... >l
° %D!\
10,000 L@ze--0-0
0.01 o' 1.0 10 100 1.000
RUPTURE TIME (hr)
Fig. 14. Creep Rupture of Experimental Incoloy 800 Containing 0.13% C and (a) 0.17% or

(v) 0.10% Ti.

Specimens were irradiated or soaked (controls)
Both tensile and creep data are plotted.

at 760°C and tested in air at 760°C.

9¢
27

2. The ratio of titanium to carbon content is unimportant to the
postirradiation ductility of Incoloy 8COC.

3. Grain size markedly affects the postirradiation strength and
ductility of Incoloy 800 at 700°C. Decreasing grain size improves
ductility and reduces strength. The increased ductility may be explained
by helium distribution or grain boundary fracture considerations.

4, Aging Incoloy 800 before irradiation to agglomerate grain
boundary carbide has no effect on the properties at 760°C but broadens
the ductility vs titanium concentration peak at 700°C. This might
allow a wider titanium specification for desired material behavior at
this temperature.

5. Creep-rupture testing after irradiation is essential to the
evaluation of Incoloy 800 for nuclear applications. Postirradiation
ductility of certain compositions increases with decreasing strain rates
(stress levels). All levels of titanium produce essentially the same
postirradiation tensile ductility.

6. Widely varying creep-rupture properties both before and after
irradiation are obtained for alloys within the Incoloy 800 composition
specification. A close control of titanium content and grain size is
imperative for elevated-temperature properties that are both desirable

and predictable.

ACKNOWLEDGMENT

I would like to express my appreciation to those of the Metals and
Ceramics Division responsible for conducting the experiments contained
in this report; J. W. Woods, V. R. Bullington, and C. K. Thomas for
construction and supervision of the ORR irradiation experiments;

B. C. Williams, C. W. Walker, L. G. Rardon, and T. T. Brightwell for
the mechanical testing; and H. R. Finch for the metallography.

I thank G. T. Newman, C. E. Sessions, and S. Peterson for their
review of the manuscript, Sharon Woods for the typing and preparation
of the report for reproduction and J. R. Weir, Jr. and C. R. Kennedy
for their helpful suggestions throughout the studies.

I am also pleased to acknowledge the support of the Division of
Space Nuclear Systems of the AEC for a portion of the work.
1-3.
4=5 .

6-15.
16.
17.
18.
19.
20.
21.
22.
23.
2 .
25.
26.
27.
28.
29.
30.
31.
32.
33.
34,
35.
36.
37.
38.
39.
40.
41,
42.
43,
bty
45,
46.
47.
48.
49.
50.
51.
52.
53.
54.
55.
56.
57.
58.

Central Research Library 59,
ORNL — Y-12 Technical Library 60.

Document Reference Section 61.
Laboratory Records Department 62.
Laboratory Records, ORNL RC 63,
ORNL Patent Office 64 .
G. M. Adamson, Jr. 65.
T. E. Banks 6685,
J. H. Barrett 86.
S. E. Beall 87—89,
R. J. Beaver aQ.
M. Bender 91,
R. G. Berggren 92,
J. O. Betterton, Jr. 03,
D. S. Billington O,
E. E. Bloom a5,
A. L. Boch 96.
E. S. Bomar Q7.
B. S. Borie 08.
G. E. Boyd 99,
R. A. Bradley 100.
R. B. Briggs 101.
R. E. Brooksbank 102.
W. E. Brundage 103.
D. A, Canonico 104.
R. M. Carroll 105,
J. V. Cathcart 106.
A. K. Chakraborty 107,
Ji Young Chang 108.
G. W. Clark 109.
K. V. Cock 110,
G. L. Copeland 111.
C. M. Cox 112.
F. L. Culler 113.
J. E. Cunningham 114,
H. L. Davis 115.
V. A. DeCarlo 116.
J. H. DeVan 117.
C. V. Dodd 118,
R. G. Donnelly 119.
K. Farrell 120.
J. 8. Faulkner 121.
J. I. Federer 122,
D. E. Ferguson 123.
R. B. Fitts 124,
B. E. Foster 125,
A, P. Fraas 126.

INTERNAL: DISTRIBUTION

29

P ORGP PSRN SR YRR PPN AR IR EUER Sy nE 3y

ORNL-TM-2305

H Frye, Jr.
Fulkerson

mEToHMmOQUURLM

Godfrey, Jr.
Gray

Grimes
Guberman
Hamner
Harman
Harms

Hill

Hinkle
Hobson

. Hoffman
. Horton

Huntley

Inouye

W.

Keilholtz

Komatsu

HEr=E2adaoaedagx GHEHBEQH §3;13:E1¢123p*u:$1?1$:F=

Lane
Leitnaker
Lindemer
Litman
Livesey
Lotts
Lyon
MacPherson
MacPherson
Martin
Martin
McClung
McCoy, Jr.
McCurdy
McDonald
McDuffie
McElroy
McHargue
McQuilkin
Miller
Miller
Moore

. Morgan
. Neill

Newman
Nilsson

. Nolan (K-25)

Ohr
Olsen
127.
128.
129,
130.
131.
132.
133,
134,
135.
13e6.
137.
138.
139.
140,
141.
142.
143,
144,
145,
146.
147.

187-190.
191~193.
194.
195.
196.
197.
198.
199.
200.

201,

202.
203.
204,
205.
206.

207.
208.
209-211.
212.

213.
214,

30

P. Patriarca 148, I. Spiewak

W. H. Pechin 149, W. J. Stelzman
A, M. Perry 150. J. 0. Stiegler
S. Peterson 151. D. B. Trauger
R. A. Potter 152. R. P. Tucker
R. B. Pratt 153. G. M. Watson
M. K. Preston 154, M. S. Wechsler
R. E. Reed 155, A. M. Weinberg
D. K. Reimann 156-175. J. R. Weir, dJr.
P. L. Rittenhouse 176. W. J. Werner
W. C. Robinson 177. H. L. Whaley
M. W. Rosenthal 178. G. D. Whitman
C. F. Sanders 179, J. M. Williams
G. Samuels 180. R. K. Williams
A. W. Savolainen 181. R, 0. Williams
J. L. Scott 182. J. C. Wilson
J. D. Sease 183. R. G. Wymer

O. Sisman 184. F. W. Young, Jr.
G. M. Slaughter 185. C. S. Yust

S. D. Snyder 186. A, F, Zulliger
K. E. Spear

EXTERNAL DISTRIBUTION

W. Albaugh, Battelle, PNL
J. Allio, Westinghouse Atomic Power Division
D. Baker, Los Alamos Scientific Laboratory
Baroch, Babcock and Wilcox
P. Calkins, General Electric, NMPO
Christopher, Combustion Engineering, Inc.
B. Coburn, Gulf, General Atomic
F. Cope, RDT, S8R, AEC, Oak Ridge National Laboratory
K. Dicker, Division of Reactor Develcopment and Technology,
AEC, Washington
D. E. Erb, Division of Reactor Development and Technology,
AEC, Washington
. Evans, General Electric, Vallecitos
. Francis, Idaho Nuclear Corporation
Goodjohn, Gulf, General Atomic
. Grant, Babcock and Wilcox
Griffo, Division of Space Nuclear Systems,
Washington
Grove, Mound Laboratory
. Gurinsky, Brookhaven National Laboratory
Kane, Lawrence Radiation Laboratory, Livermore
Haruo Kato, Albany Metallurgy Research Center, P.0. Box 70,
Albany, Oregon 97321
E. E. Kintner, Fuel Fabrication Branch, AEC, Washington
J. H. Kittel, Argonne National Laboratory

?).C’Uffl<if')wwfij

“n P
fd*fi*dc4c33>

& oo
w02
215.
216.
217.
218.
219.

220,
221.
222,
223.

224,
225,

226.
227 .

228.
229.

230,
231.
232—234.

235,
236.

R37—239.
240,
241,

42

243,
ey,
245,
246,

2477,
248,

249,
250~264 .

31

J. Kreih, Westinghouse, Bettis Atomic Power Laboratory
J. Larkin, AEC, Oak Ridge Operations

L. Larsen, Iowa State University, Ames Laboratory
. F. Leitten, Jr., Linde Division, Union Carbide Corp.
J. Levine, Westinghouse Advanced Reactor Division,
altz Mill Site, Box 158, Madison, Pa. 15663
J. J. Lombardo, NASA, Lewis Research Center

J. H. MacMillan, Babcock and Wilcox

C. L. Matthews, RDT, SR, AEC, Osk Ridge National Laboratory
A. Mullunzi, Division of Reactor Development and Technology,
AEC, Washington
M. Nevitt, Argonne National Laboratory

R. E. Pahler, Division of Reactor Development and Technology,
AEC, Washington

S. Paprocki, Battelle Memorial Institute
W. E. Ray, Westinghouse Advanced Reactor Division,

Waltz Mill Site, Box 158, Madison, Pa. 15663

B. Rubin, lLewrence Radiation Laboratory, Livermore

F. C. Schwenk, Division of Reactor Development and Technology,
AEC, Washington

F. Sheely, Division of Research, AEC, Washington

G. Shewmon, Argonne National Laboratory
M. Simmons, Division of Reactor Development and Technology,
AEC, Washington

* »

H
L

L. E. Steele, Naval Research Laboratory
R. H. Steele, Division of Reactor Development and Technology,
AEC, Washington

D. K. Stevens, Division of Research, AEC, Washington
A. Strasser, United Nuclear Corporation
A. Taboada, Division of Reactor Development and Technology,
AEC, Washington
A. Van Echo, Division of Reactor Development and Technology,
AEC, Washington
James Watson, Gulf, General Atomic
C. E. Weber, Atomics International
J. F. Weissenberg, RDT, OSR, General Electric, NMPO
G. W. Wensch, Division of Reactor Development and Technology,
AEC, Washington
G. A. Whitlow, Westinghouse Advanced Reactor Division,
Waltz Mill Site, Box 158, Madison, Pa. 15663
E. A. Wright, AEC, Washington
Laboratory and University Division, AEC, Oak Ridge Operations
Division of Technical Information Extension