ORNL-TM-2043.txt

LOCKHEED MARTIN ENERGY RESEARCH 185,

N

ARIES

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I,

 
e e, e LEGAL NOTICE oo e

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ORNL-TM-2043

Contract No. W-7405-eng-26

METAIS AND CERAMICS DIVISTION

FEFFECTS OF IRRADIATION ON THE MECHAWICAL PROPERTIES OF
TWO VACUUM-MELTED HEATS OF HASTELLOY N

H. E. McCoy, Jr.

JANUARY 1968

OAK RIDGE NATIONAL LABORATORY
O=k Ridge, Tennessee
operated by
UNIQON CARBIDE CORPORATION
for the
U.S. ATOMIC ENERGY COMMISSION

A

3 4456 051324y 2
Abgtraet . . . . . . . o .

Introduction . . . . . . .

Experimental Details . .
Test Materials . . . .
Heat Treatments . .

Test Specimen

Irradiation Conditions .

Testing Techniques . .

Experimental Results .
Discussion of Results . .
Summary and Conclusions

Acknowledgments . . . . .

TABLE

iii

OF CONTENTS

Page

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EFFECTS OF TRRADIATION ON THE MECHANICAL PROPERTIES OF
TWO VACUUM-MELTED HEATS OF HASTELLOY N

 

 

H. E. McCoy, Jr.

ABSTRACT

The mechanical behavior of two vacuum-melted heats
of Hastelloy N was evaluated at 650 and 76¢0°C. The
material was subjected to several thermal-mechanical
treatments and then irradiated at 650 and 760°C to a
thermal dose of 2.3 x 104 neutrons/cmz. The results
are compared with those for unirradiated specimens that
were given a similar thermal treatment. The various
thermal-mechanical treatments had some relatively small
effects on the unirradiated tensile properties, but the
creep properties were very similar. The primary effects
of irradiation were reductions in the creep-rupture
life and the rupture ductility in both creep and tensile
tests. These observations are explained on the basis of
helium production in the metal by the lOB(n,oz) transmutation.

 

INTRODUCTION

The potential use of Hastelloy N in several reactors has developed
considerable interest in how the properties of this material change
with neutron irradiation. Previous studies at ORNL'»? have shown that
the high-temperature mechanical properties of this alloy deteriorate
under neutron irradiation. This deterioration manifests itself as both
a reduction in the creep-rupture life and in the rupture ductility.
However, these studies involved air-melted material with B levels in the

range of 20 to 50 ppm.

 

. R. Mertin and J. R. Weir, Nucl. Appl. 1(2), 160-167 (1965).
*W. R. Martin and J. R. Weir, Nucl. Appl. 3(3), 167-177 (1967).
We have recently run two series of experiments that were aimed
primarily at characterizing this alloy for use in the SNAP-8 system.
This system will uvtilize thin-walled Hastelloy N tubing for fuel element
cladding and will operate over the temperature range of 650 to 760°C.
One series of experiments involved in-reactor tube-burst tests on cladding
material, and the details of this study have been reported.? The second
series of experiments involved postirradiation creep rupture and tensile
tegts on small bar samples, and these results are presented in this
report. Two vacuum-melted heats were used, and the properties were
evaluated after the material had been subjected to various thermal-
mechanical treatments. These treatments were dictated largely by the

steps used to process the fuel element cladding.

EXPERIMENTAT, DETATILS

Test Materials

The two lots of material used in this study were 12-in.-diam,
10,000-1b double vacuum-melted heats obtained from Allvac Metals Company.
The chemical analysis of each heat is given in Table 1. Heat 5911 was
obtained in two forms: forged to a bar 2 x 2 in. (designated 5911 AW)
and forged and machined to a tube shell 2-in. OD x 1 1/2-in. ID
(designated 5911 TH). Heat 6252 was obtained in the as-cast condition
(designated 6252 AC).

Heat Treastments

The materials were given several different mechanical and thermal
treatments prior to irradiation. These treatments are described in
Table 2 and will be referred to by number. All annealing was carried
out in an argon environment, and the specimens were cooled by pulling

them into a water-cooled section at the end of the furnace.

 

*H. E. McCoy, Jr., and J. R, Weir, In- and Ex-Reactor Stress-Rupture
Properties of Hastelloy N Tubing, ORNL-TM-1906 (September 1957).

 
Table 1. Chemical Analysis of Test Materials

 

Content (wt %)

 

 

Element
Heat Number 5911 Heat Number 6252

Fe 0.03 0.12
Cr 6.14 7.26
Mo 17.01 16.53
Ni bal bal

C 0.056 0.051
Mn 0.21 ' 0.20

B 0.0010 0.0003
S 0.0022 0.0022
P 0.0022 0.002%
Si 0.05 0.05
Cu < 0.01 < 0.01
Co 0.08% 0.11%
Al 0.15 0.20
Ti 0.067 0.13
W 0.018 0.02%
Zr < 0.01 < 0.01

0 0.0014 0.0008
N < 0.0005 0.0005

 

%Ladle analysis, Allvac Metals Company.
A1l other values obtained at ORNL on the finished
product.

Test Specimen

The small tensile specimen shown in Fig. 1 was used for in- and ex-
reactor tests. The small size made it possible to get several specimens
into a single experiment. Our work with this specimen has shown that it
yields data that are quite similar to those obtained from larger specimens.
Because of the stress concentration at the base of the filet, there is

some tendency for the brittle specimens to fail at this point.

Irradiation Conditionsg

The specimens were irradiated in a single-test capsule in the P-4
poolside position in the ORR. The peak thermal flux was
2 sec”l, and the peak fast (> 2.9 Mev) flux was

sec'l.

6 X 101% neutrons em”

2

5 % 10%? neutrons cm” The duration of the experiment was 1080 hr
Table 2. Description of Thermal-Mechanical Treatments

 

 

Designation Thermal ~-Mechanical Treatment
1 Annealed 1 hr at 1177°C in argon
2 Hot rolled 50% at 1150°C,

Annealed 1 hr at 1177°C in argon

3 Annealed 1 hr at 1177°C in argon,
Cold worked 25%,

Annealed 1 1/2 hr at 1066°C
in argon,

Annealed 2 hr at 1093°C in argon ,
Annealed 10 min at 1150°C in argon

3A Annealed 1 hr at 1260°C in argon

24 Hot rolled 50% at 1150 °C,
Annealed 1 hr at 1260°C in argon

 

ORNL-DWG 67-3013

 

 

 

 

 

 

 

 

 

 

 

 

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Fig. 1. Test Specimen
(time at temperature and full power), so the thermal and fast doses were
2.3 x 10%9 and 1.9 x 10%°? neutrons/cmz, respectively. Each specimen was
heated by a small furnace, and the temperature was controlled by a
proportioning controller which acted on response to a Chromel-P-Alumel
thermocouple attached to the specimen gage length. Some of the specimens
were controlled at 650°C, but most of them were held at 760°C. The
enviromment in the capsule was flowing He—1% 0. Ex-reactor control

specimens were given the same thermal exposure as the in-reactor sgspecimens.

Testing Techniques

The laboratory creep-rupture tests were run in conventional creep
machines of the dead load and lever arm types. The strain was measured
by a dial indicator that showed the total movement of the specimen and
part of the load train. The zero strain measurement was taken immediately
after the load was applied. The temperature accuracy was #0.75%, the
guaranteed accuracy of the Chromel-P-Alumel thermocouples used.

The postirradiation creep-rupture tests were run in lever arm
machines that were located in hot cells. The strain was measured by an
extensometer with rods attached to the upper and lower specimen grips.
The relative movement of these two rods was measured by a linear dif-
Terential transformer, and the transformer signal was recorded. The
accuracy of strain measurements is difficult to determine. The exten-
someter (mechanical and electrical portions) produced measurements that
could be read to about #0.02% strain. However, other factors (temperature
changes in the cell, mechanical vibrations, etc.) probably combined to
give an overall accuracy of #0.1% strain. This is considerably better
than the specimen-to-specimen reproducibility that one would expect for
relatively brittle materials. The temperature measuring and control
system was the same as that used in the laboratory with one exception.
In the laboratory, the control system was stabilized at the desired
temperature by use of a recorder with an expanded scale. 1In the tests
in the hot cells, the control point was established by setting the
controller without the aid of the expanded-scale recorder. This error
and the thermocouple accuracy combine to give a temperature uncertainty

of about =*1%.
The tensile tests were run on Instron Universal Testing Machines,
The strain measurements were taken from the crosshead travel.

The test environment was air in all cases. Metallographic exami-
nation showed that the depth of oxidation was small (< 0.002 in.), and
hence, we feel that the enviromment did not appreciably influence the

test regults.

Experimental Results

The resgults of tensile tests run on the materials in this study
are summarized in Tables 3 and 4. All of the unirradiated specimens
were subjected to a thermal aging treatment (1080 hr at 650 or 760°C)
equivalent to that of the irradiated specimens. The data in Table 3
indicate several important features of the tensile properties of the
unirradiated materials.

1. The yield stress decreases slightly with increasing test

temperature, whereas the tensile stress decreases by about
a factor of 2 over the temperature range of 550 to 760°C.
2. There appear to be some small variations in the yleld

and tensile strengths due to the various heat treatments.
For example, the data for heat 5911 AW-anneal 1 indlcate
that aging at 650°C results in lower yield stress and a
higher tensile stress than aging at 760°C. The anneals

at 1260°C (3A and 2A) cause slight strength reductions.

3. The fracture ductility decreases with increasing test

temperature for all materials.
4. The data for 5911 AW-anneal 1 indicate that aging at
650°C results in better ductility than aging at 760°C.

5. A comparison of the data for 5911 AW-anneals 1 and 3
indicate that the ductility of the materials receiving
the anneal 3 is lower when aged and tested at 650°C
and higher at 760°C.

6. Heat 6252 AC generally exhibited lower ductility than

heat 5911. This is probably due to the smaller amount
of working received by heat 6252 AC. This is supported
 

 

 

Table 3. Tensile Properties of Unirradiated Materials®
m . - \ »
. b Specimen Test Stress (psi) Elongation (%) R?ductlon Pretest
nneal Thamber Temperature in Area AzingC
(°c) Yield Tensile Uniform Total (%) ging
Heat 5911 AW
1 3113 550 32,600 87,800 57.1 58.2 41.6 2
1 3114 00 35,400 65,200 22.5 23.5 23.2 2
1 3109 760 32,200 47,700 8.0 13.4 14.4 2
1 3115 760 31,700 51,100 12.8 16.6 13.1 2
1 3117 550 31,400 96,500 61.7 64.8 47.5 1
1 3108 650 29,700 77,700 39.8 40.8 35.3 1
1 3121 760 28,900 50,900 12.5 25.9 27.3 1
34 3106 650 26,400 70,800 43.2 43.4 36.0 1
34 2854 760 31,700 46,700 8.0 13.6 12.5 2
3 2959 650 34,300 76,600 29.6 30.3 25.7 1
3 2953 760 32,200 45,800 8.2 22.9 22.1 2
Heat 5911 TH
1 3103 650 29,200 74 ,300 41.1 42.7 33.5 1
1 3091 760 32,400 47,800 7.7 16.2 13.9 2
Heat 6252 AC
2 2919 650 34,000 66,600 15.1 15.4 17.4 1
2 2921 760 32,000 45,000 6.1 8.6 7.3 2
2 + 3 2948 650 37,100 73,500 21.6 21.9 19.1 1
2+ 3 2944 760 33,500 44,700 8.7 20.2 17.6 2
24 3130 760 25,700 46,800 7.7 10.6 10.1 2

 

aAt a strain rate of 0.002 min

-1

bAnneal designation given in Table 2.
®1 — 1080 hr at 650°C; 2 — 1080 hr at 760°C-
Table 4.

Tensile Properties of Irradiated Materials®

 

Temperature (°C)

Stress {psi)

Elongation (%)

 

 

 

Reduction

 

Heat Specimen )

Number Anneal Number in Aresg
- Irradiation Test Yield Tensile  Uniform  Total (%)

5911 AW 1 2837 760 760 32,600 32,700 1.0 1.0 3.6
5911 AW 1 2839 760 760 33,700 33,700 1.1 1.4 2.0
5911 AW 1 2846 650 650 38,200 50,600 7.7 8.1 11.2
5911 AW 1 28477 650 650 36,400 52,600 10.5 11.2 14.6
6252 AC 2 2917 760 760 32,200 32,200 1.1 1.5 0.9
6252 AC 2 2918 760 760 31,800 31,800 1.2 1.5 0.0
6252 AC 2 2926 650 650 39,100 48,500 4.6 4.8 5.6
6252 AC 2 2927 650 650 40,200 48,900 4.6 5.6 8.0

 

"Thermal dose equals 2.3 X 1

020

neutrons/cng

straln rate equals 0.002 min™=.

i
by the fact that the additional working received in the
2 *+ 3 treatment improved the ductility over that obtained
after just the treatment 2.
The tensile properties of some of the materials were obtained at
650 and 760°C after irradiation, and these results are given in Table 4.
A comparison of these data with those for the unirradiated specimens in
Table 3 leads to several important observations:
1. The yield stress at 650°C is higher for the irradiated
specimens whereas the yield stress at 760°C is equivalent
for irradiated and unirradiated specimens.
2. The tensile stress is lower for the irradiated materials
at both 650 and 760°C.
3. The ductility at both temperatures is reduced severely
by irradiation, the reduction being much greater at 760°C.
The variations in the rupture ductility of Hastelloy N in the various

conditions investigated are summarized in Fig. 2. The spread in the rupture

ORNL-CWG &7-7252
70 . : i~
‘ © 5911 AW-1-AGED AT 550°C
6 ® 591 AW-1-AGED AT 760°C
2 591 AW-34
O 5911 AW -3
**** e LT g BANTH -
. 0 6252402
mG232AC-2+3
A 8252AC-2A
-+ IRRADIATED

 

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FRACTURE STRAIN (%}

 

 

 

 

 

 

 

 

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500 550 600 650 700 750 80C
TEST TEMPERATURE (°C)

Fig. 2. Variation of the Tensile Fracture Strain with Temperature.
(Unless designated otherwise, specimens were irradiated or aged at the
test temperature.)
10

strain at a given test temperature is quite large. At 650°C, for example,
the unirradiated material shows a range of about 15 to 44%, and the range
is further extended by values as low as 5% in the irradiated condition.

The main emphesis in this study was on the creep-rupture properties,
since the potential application involves service under creep conditions.
The results of tests on unirradiated and irradiated specimens are given
in Tables 5 and 6, respectively. The data on unirradiated materials are
of wvalue themselves, but we are more concerned with how the properties
change as a result of irradiation. Figure 3 shows that the rupture
ductility varies from about 15 to 30% for the unirradiated material at
650°C. However, many of the variations appear to be random rather than
due tc the effects of a particular pretest heat treatment. The general
trend seems to be for decreasing ductility with increasing rupture life.
Figure 4 compares the rupture lives of unirradiated and irradiated
specimens at €50°C. The line for the irradiated material is based cn
our results for several air-melted heats. The results on the present heat
at 650°C are inadequate to establish & stress-rupture curve, but the data
are in reasonable agreement with those for the alr melts. The rupture life
variations appear to be entlirely random with respect to material and
conditions of annealing. The rupture life is reduced by an order of
magnitude by irradiation, but there is some indication that this factor
decreases with increasing rupture life. The minimum creep rate is shown
in Fig. 5 as a function of the stress at 650°C for irradiated and unirra-
diated materials. The wvariation due to heat and anneal 1s again random
and irradiation does not have any appreciable effect.

The rupture strain at 760°C is shown in Fig. & as a function of
rupture life for unirradiated material. The rupture strain varies from
10 to 50% with most of the variation appesring to be independent of
heat and anncaling treatment. Most of the materials exhibit a trend
of increasing strain with increasing rupture life. The rupture lives
of the irradiated and unirradiated materials are compared in Fig. 7.
Again, the variations due to heat and anneal appear to be random. One
exception may be heat 6252 AC, which, after irradlation, has a greater
rupture life. At a stress level of 20,000 psi, the rupture life is

reduced about two corders of magnitude by irradiation. The curves converge
Table 5. Creep-Rupture Properties of Unirradiated Materials

 

 

_ Test Minimum Rupture Rupture Reduction
NE;%Er Anneal® N§;§Zr Temperature ?;:i?s Creep Rate Life Strain in Area irizegt
(°C) (%/hr) (br) (%) (%) sHne
5911 AW 1 6185 650 65,000 0.590 9.4 26.5 21.9 1
5911 AW 1 6014 650 55,000 0.140 49.6 27.4 21.2 1
5911 AW 1 6013 650 47,000 0.043 206.5 17.3 17.1 1
5911 AW 1 6012 650 40,000 0.022 413.7 17.3 16.7 1
5911 AW 1 6186 650 43,000 0.018 598.2 22.7 11.0 1
5911 AW 1 6059 650 32,400 0.0045 1828.1 21.8 19.8 1
5911 AW 1 6126 760 30,000 0.83 21.8 29.8 23.1 2
5911 AW 1 6187 760 25,000 0.36 53.7 31.9 13.9 2
5911 AW 1 6023 7¢0 20,000 0.12 147.7 35.5 20.6 2
5911 AW 1 6039 760 17,500 0.049 367.7 24.3 10.6 2
5911 AW 1 6024 760 15,000 0.035 e07.2 39.8 29.0 2
5911 AW 1 6188 760 13,000 0.017 1198.¢9 29.7 16.0 2
5911 AW 3A 6057 650 47,000 0.026 120.8 27.2 21.1 1
5911 AW 3A 6015 650 40,000 0.012 489.0 21.0 17.3 1
5911 AW 3A 6040 760 30,000 0.94 15.4 23.6 16.1 2
5911 AW 34 6025 760 20,000 0.14 114.7 R4 .6 16.1 2
5911 AW 3A 6026 760 15,000 0.025 402 .8 13.3 8.0 2
5911 AW 3 6058 650 47,000 0.067 181.0 23.5 20.7 1
5911 AW 3 6016 650 40,000 0.025 761.7 29.5 22.8 1
5911 AW 3 6127 760 30,000 1.18 24 .2 38.5 33.0 2
5911 AW 3 6189 760 25,000 0.42 47.0 34 .4 15.0 2
5911 AW 3 6027 760 20,000 0.19 73.7 17.6 21.7 2
5911 AW 3 0041 760 17,500 0.13 153.3 29.3 17.3 2
5911 AW 3 6028 760 15,000 0.048 419.2 32.2 16.4 2
5911 TH 1 6042 650 55,000 0.018 48.9 27.9 21.6 1
5911 TH 1 6018 650 47,000 0.038 189.1 26.6 21.9 1
5911 TH 1 6017 650 40,000 0.023 706.3 29.2 25.4 1
5911 TH 1 6056 650 32,400 0.0019 2082.1 18.7 13.4 1
5911 TH 1 6125 760 30,000 0.765 17.7 19.2 14.3 2

1T
Table 5 (continued)

 

 

Test Minimum Rupture Rupture Reduction
Ngeizr Anneala NE;EQT Temperature ?tz§§s Creep Rate Life Strain in Area irigegt
e (°c) P (%/hr) (hr) (%) (%) S1e
5911 TH 1 190 760 25,000 0.41 27.7 14.1 10.7 2
5911 TH 1 6029 760 20,000 0.13 113.5 25.3 16.2 2
5911 TH 1 6043 760 17,500 0.054 161.4 11.9 10.0 2
5911 TH 1 6030 760 15,000 0.026 6547 24.8 13.2 2
6252 AC 2 6022 650 47,000 0.054 9.3 .4 12.2 1
6252 AC 2 6019 650 40,000 0.026 64'7.3 21.2 15.6 1
6252 AC 2 6060 650 32,400 0.0041 1813.3 15.9 14.7 1
6252 AC 2 6124 760 30,C00 1.37 7.3 14.1 10.2 2
6252 AC 2 6191 760 25,000 0.51 48,9 31.2 14.3 2
6252 AC 2 6031 760 20,000 0.19 119.1 39.0 29.0 2
6252 AC 2 6044 760 17,500 0.13 130.8 22.5 15.7 2
6252 AC 2 6032 760 15,000 0.049 636.5 46.7 37.7 2
6252 AC 2A 6033 760 20,000 0.13 162.9 35.6 21.9 2
6252 AC 2A 6034 760 10,000 0.0047 3570.2 28.0 13.7 2
6252 AC 2 + 3 6021 650 47,000 0.073 123.6 17.5 13.9 1
6252 AC 2 + 3 6020 650 40,000 0.024 £22.2 14.9 14.0 1
6252 AC 2 + 3 6045 760 30,000 0.80 10.0 13.1 12.0 2
6252 AC 2+ 3 o047 760 20,000 0.24 114.3 53.1 31.6 2
6252 AC 2+ 3 6046 760 15,000 0.044 535.6 46.6 21.6 2

¢t

 

®Arneal designation given in Table 2.

Py _ 1080 hr at 650°C; 2 — 1080 hr at 760°C.
Table 6. Creep-Rupture Properties of Irradiated Materials®

 

 

Test and
— Minimum Rupture Rupture Reduction Specifi-
Ngzii:r AnnealP NE;@; é;;agizziiz S%rzjf’ Creep Rate Tife  Strain  in Area cation Comments
P (°C) r (%/nhr) (hr) (%) (%) Number

5911 AW 1 R-194 650 47,000 0.056 12.8 1.68 0.9 2851

5911 AW 1 R-182 650 40,000 G.015 43.0 0.83 8.2 2850

5911 AW 1 R-170C 650 32,400 0.0049 144.3 0.84 —4.0 2848 C
5911 AW 1 R-177 760 10,000 0.0086 104.6 0.98 0.3 2842

5911 AW 1 R-196 760 g,000 0.0011 834.8 1.77 G.0 2841

5911 AW 3A R-183 7¢0 15,000 0.22 2.2 0.57 C.8 2849

5911 AW 3A R-201 760 10,000 0.0041 179.4 1.08 1.1 2845

5911 AW 3 R-175 760 20,000 0.85 0.5 0.58 1.5 2950

5911 AW 3 R-180 760 10,000 C.011 55.2 1.34 4.5 2951 c
5911 AW 3 R-223 760 8,000 0.0024 825.7 4 .49 G.0 2949

5911 TH 1 R-176 760 20,000 0.28 1.0 0.67 3.1 2981

5911 TH 1 R-184 760 15,000 0.090 2.1 0.99 —0.3 2980

5911 TH 1 R-218 760 8,000 0.0024 365.2 1.45 2982

6252 AC 2 R-197 650 47,000 0.021 16.8 1.50 0.4 2937

6252 AC 2 R-181 650 40,000 0.029 23.5 0.93 1.1 2929

6252 AC 2 R-173 650 32,400 0.0C50 220.6 1.70 0.3 2928 C
0252 AC 2 R-185 760 15,000 0.18 3.5 C.95 0.0 2920

6252 AC 2 R-224 760 12,500 0.039 24.1 1.55 4.8 2923

6252 AC 2 R-178 760 10,000 0.0032 581.5 2.36 -1.8 2922

6252 AC 24 R-186 760 15,000 0.36 Q.7 C.32 —2.2 2930 c
6252 AC 2A R-217 760 10,000 0.0028 289.8 1.12 ~3.7 2932 c
6252 AC 2 + 3 R-187 760 15,000 0.17 5.0 1.49 2.5 2939

6252 AC 2 + 3 R-225 760 12,500 0.0067 284.3 3.18 2941

6252 AC 2+ 3 R-209 760 10,000 0.0046 504.6 3.58 3.5 2940

et

 

Dose equals 2.3 x 1020 neutrons/cm?.

“Anneal designation given in Table Z.

CSpecimen broke in the radius at the end of gage sectilon.
14

ORNL-DWG 67-7254

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

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& e 6252 AC-2+3 | | |
10 [ b |- 1
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1 10 100 1000 10,000

RUPTURE LIFE (hr)

Fig. 3. Variation of the Rupture Strain with Rupture Life Under
Creep at 650°C.

ORNL-DWG 67-7255R

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

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IRRADIATED
50 | ® 6252 AC-2+3 e AR MELTS N
IRRADIATION TEMPERATURE — 65C° C
THERMAL DOSE-2.3x40%° aut
N L
10 ‘ ‘
0 \ i ‘ . :
1 10 100 1000 10,000

RUPTURE LIFE (hr)

Fig., 4. A Comparison of the Creep-Rupture Properties of Irradiated
and Unirradiated Hastelloy N at 650°C.
15

ORNL. DWG 67-7256

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

70 : —
‘ | 11
| J A7
! ‘ -1 ‘
8O f-omom et b o
%
50 | % L AT
&l vofpoe |
= L1 1
oy | :
Q. : o \
g 40 246 ot L
e 4 ‘
-~ d |
4 v i ‘
& » o 5311 AW-{ ‘
g 30— //‘?G# & 59 AW-3a T |
x , o 5911 AW-3
0 r v 504 TH-1
| ¢ 6252 AC-2
20 | ‘ ® 6252 AC-2+3
| + IRRADIATED
! IRRSAD[ATION TEMPERATURE
— 0°
10 b ; THERMAL DOSE - 2.3 x40%° sy |11
I ! i
UL LI ey
o } -’ o ]
Q.00 0.0 oY i

MINIMUM CREEP RATE (% hr)

Fig. 5. A Comparison of the Creep Rates of Irradiated and
Unirradiated Hastelloy N at 650°C.

CRNL—DWG &7-72567

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

70 -1
| T LT |
© 594t AW ~1 !
60 A5 AW-3A & | |
0 5941 AW -3 |
v 5941 TH— !
0 6252 AC-2 ¢ ;
.50 ° 6252AC-2+3 [ T
¥® ) *0
=z | !
< 40 ‘ pepeeseeenees - 5t e
5 O cA
EE' 30 --10- LO gy a i -
D o o a
l...
o % ||
2 a o °
20 R R T et
{ |
R e ¥ v A !
10 g + ?
|
| :
0 ! il
{ 10 100 1000 10,000

RUPTURE LIFE (hr)

Fig. 6. The Variation of Rupture Strain with Creep-Rupture Life
for Hastelloy N at 650°C.
16

ORNL-DWG 67-7258

UNIRRADIATE

JIRRADIATED 1477 om

°

594 AW-1
5914 AW -3A

5941 AW-3

B TH-1

6252 AC-2

6252 AC-2+3

6252 AC-2A

IRRADIATION TEMPERATURE -760°C
THERMAL DOSE-2.3x1029 ayt

STRESS (1000 psi)

 

oA { 10 to0 1000 10,000
RUPTURE LIFE (hr)

Fig. 7. A Comparison of the Creep-Rupture Properties of Irradiated
and Unirradiated Hastelloy N at 760°C.

slightly as the stress level is reduced. The minimum creep rates of the
irradiated and unirradiated specimens are compared in Fig. 8. At a stress
level of 20,000 psi, the creep rate of the irradiated material is higher.
As the stiress is decreased, the curves converge to where the difference
in creep rate below 10,000 psi is negligible.
The rupture strains of the irradiated specimens are compared in
Fig. 9 as a function of strain rate. Although the tensile tests at a
strain rate of 0.002 min~t (12%/hr) indicated that the ductility was
much lower at 760°C than at 650°C, the ductility appears to be independent
of temperature at strain rates below about O.l%/hr. There is also a
distinct trend of inecreasing fracture strain with decreasing strain rate.
Although all of the creep curves were examined in detail, only a
few typical ones will be presented. The data were taken manually, but
were reduced and plotted by computer. Figures 10, 11, 12, and 13 are
direct photographs of the plotted data after the curves were drawn in
manually. Figures 10 and 11 show the results of tests on unirradiated
and irradiated specimens, respectively, from heat 5911 AW-anneal 1 at
650°C and 40,000 psi. The primary creep stage is very short, and the
accumulated strain at the beginning of secondary creep was only aboutl

0.1%. A comparison of points on the two plots indicates that the results
17

ORNL-DWG £€7-7259

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

40
-~
Yo ‘%‘ ‘C(
o 7"’}.’1{?
(27 |
20 /},&40- ’J—” =
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g,do-f Y -ea a-] ,
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UNlRRADIATfD/f/ ¢ e i
=10 i + LKJ : ..........................
o <+~
8 './". i ‘
g AT ; v E
n f ‘
@ i
E 5l o 5941 AW-1
@ & 5944 AW-3A
0 5944 AW-3
v 5944 TH-1 : ‘
0 6252 AC-2 1 |
O 6252 AC-2+3 ‘ 1
& 6252 AC-2A
o L CLOSED SYMBOLS IRRADIATED
IRRADIATION TEMPERATURE —760°C
THERMAL DOSE-2.3%1029 ny¢
i
, | |
0.00014 0.004 0.04 0 1 10

MINIMUM CREEP RATE (%/hr)

Fig. 8. A Comparison of the Creep Rates of Irradiated and
Unirradiated Hastelloy N at 760°C.

+ ORNL-DWG 67-3504R

‘ ?8.1 "2

 

Il

O 591 - AW-1—-650°C
& T A G252 - AC ~2-650°C
® 5541 —-AW — 1 — 750°C
A G252 -AC-2-760°C
® 591 —TH-1—~760°C

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

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¥ 6252~ AC-2+3 —760°C A
N ¢ 591 —AW-3 —760°C |
$ ® 59 - AW-3A~T760°C | :
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2 | |
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a
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n ' A 4 v ?
1 9.2 & & ¢ -
° !
o " 4 l
e :
o | i
0.001 0.04 oA 1 10 100

STRAIN RATE (% /hr)

Fig. 9. Fracture Ductilities for Hastelloy N Under Several Test
Conditions.
18

 

~8 » HEAT 5911RW » HERT TRERT NO. |

TEST NO 5186 » INOR

i e

- e g L]t
-l I e

=2

LT

 

 

 

 

 

 

 

 

 

 

 

 

 

f

 

 

 

 

 

 

 

 

200, D

TIMEs HR.

reep Curve for Test 6186,

C

Fig. 10.

T Fe e Ty

Tmunm..mflla ‘r.,thfH.uu,Lu‘i

s

LR R g Sl e I AT ey .,,_ L e ur,u,u.u,.ufl.nrllu.“.im..i,,.lfimu‘mm.“ft.w.‘A

 

 

 
19

TEST NO. R-182» INOR-Bs HEAT NO. 59118Ws HEAT TREATMENT NO 1.

Q
R

 
  
   

10
TIMEs HR.
Fig. 11. Creep Curve for Test R-182.
20

TEST NO. 6034s INOR-NT NO 2R

 

1800.0 L B O k.0

B00.0

TIMEs HR.

Creep Curve for Test €034,

12.

ig.

n
21

TEST NO R-217s INOR-8» HEAT B252RC 8NN, 2R

<
o~

 

.0 0 A0 ¥o.0
TIMEs HR,

Fig. 13. Creep Curve for Test R-217.
22

are quite similar up to the point where the irradiated specimen failed.
Figures 12 and 13 show the results of tests at 760°C and 10,000 psi on
unirradiated and irradiated specimens, respectively. There 1s no dis-
tinguishable primary creep stage, and the results are quite comparable
up to fracture of the irradiated specimen.

Since considerable data were obtained on heats 5911 AW and 5911 TH
with heat treatment 1, the creep properties of these materlials were
examined in deteil. Further details of the creep behavior are given in
Table 7. At 650°C the behavior of the two lots of material was quite
similar. The times for rupture and 1% strain for unirradiated and irra-
diated material at 650°C are compared in Fig. 14. The curve for rupture
of the irradiated material falls just short of the 1% strain curve for
the unirradiated specimens. At 760°C materials 5911 AW and 5911 TH have
slightly different strength properties. Figure 15 shows the results for
heat 5911 AW-anneal 1. At 8000 psi, the time to 1% strain for the irra-
diated specimen agrees very well with that for the unirradiated material.
At 10,000 psi, the time to 1% strain for the irradiated material is
slightly less than that for the unirradiasted specimen. The data for
heat 5911 TH-anneal 1 are plotted in Fig. 16. The time to 1% strain for
the irradiated material is much shorter than that for the unirradiated
material at high stresses, but the two nearly converge at lower stresses.
This behavior agrees well with the minimum creep rate behavior shown
in Fig. 8. There is some uncertainty in Fig. 16 concerning the location
of the curve Tor the time to l% gstrain. Since an extensometer was not
used in the laboratory tests, the scatter in data for small strains is
understandable.

In an effort to determine the effects of gtrain rate on the rupture
ductility at 650 and 760°C, several specimeng were tested in the unaged
condition at various strain rates. The results of these tests are
given in Table 8. The tensile results from Tables 3 and 8 and the creep
results from Table 5 were used to obtain the plot of ductility versus

strain rate in Fig. 17. At 650°C the ductility decreases with decreasing
23

Table 7. Detailed Creep Properties on
Heats 5911 AW-1 and 5911 TH-1%

 

 

Tost Test Stress Time (hr) to Indicated Strain (%)
Number Temperature (psi)
(°C) 1 2 5 10 Rupture

™

Heat 5911 AW-1

6185 650 65,000 0.2 1.4 6.5 9.4 9.4
6014 650 55,000 17.5 46 49.6
6013 650 47,000 3.0 17 20 175 206.5
€012 650 40,000  31.5 &0 198 351 413.7
6186 650 40,000 62 113 260 445 598.2
6059 650 32,400 195 340 Ve 1322 1828.1
R-194 650 47,000 7.8 12.8
R-182 650 40,000 43.0
R-170 650 32,400 144.3
6126 760 30,000 1.0 2.1 5.7 11 21.8
6187 760 25,000 2.2 5 13 26 53,7
6023 760 20,000 9.5 17 41.5 73 147.7
6039 760 17,500 15 35 97 189 367.7
6024 760 15,000 30 63 152 277 607.2
6188 760 13,000 52 110 275 525 1198.9
R-177 760 10,000 ~105 104.6
R-196 760 8,000 645 834.8

Heat 5911 TH-1

6042 650 55,000 3 8.5 25 42 48.9
6018 650 47,000 14 40 99 166 189.1
6017 650 40,000 44 95 230 430 706.3
6056 650 32,400 260 495 1090 1750 2082.1
6125 760 30,000 1.0 2.4 6.3 12 17.7
6190 760 25,000 1.8 4.3 12 23 27.7
6029 760 20,000 g.7 16 38.5 72 113.5
6043 760 17,500 7.9 27 76 142 161.4
6030 760 15,000 52 20 202 372 654.7
R-176 760 20,000 1.0
R-184 760 15,000  ~2 2.1
R-218 760 8,000 250 365.2

 

aSee Tables 5 and & for other details.
70

60

» o
o o

STRESS (1000 psi)
o
o

20

24

 

ORNL-DWG 67~ 7260

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

z I ™ '
i {T\\Tx__ 5
T UNIRRADIATED RUPTURE
\\ UNIRRADIATED | ok | !
~. % STRAIN \\
4 4 fih‘ | \& 3 I ]
" y
\\\ . \\ |
] o do O T 41
~ \\ T
T:r\ ™
il k\\g\l Y
T “ IRRADIATED RUPTURE ™~ |~ 5
' | | |
. | | i :
B S ' ] I
i UNIRRADIATED IRRADIATED
1.1l RUPTURE 1% STRAIN RUPTURE 1% STRAIN
5911 AW~ { o o . o~ -
‘ 501 TH-1 a v |
el el L
1 10 100 1000 10,000

TIME (hr}

Fig. 14. A Comparison of the Unirradiated and Irradiated Creep
Rupture Behavior of Heats 5911 AW-1 and 5911 TH-1 at 650°C.

40

QRNL-DWG 67-7261

  

~ 20 591 AW -1

2

S ® o UNIRRADIATED

< IRRADIATED

o 4 1% STRAIN

W 10 [~ & RUPTURE

’-4

oy 8 |- p—ptferd b L b DAl
6
g - :
0.4 1 10 100 1000 10,000

TIME (hr)

Fig. 15. A Comparison of the Creep-Rupture Properties of
Heat 5911 AW-1 at 760°C in the Irradiated and Unirradiated Conditions.
ORNL-DWG 67-7262

40

n
Q

 

,’5
o
o
g RUPTURE
= 5oL~ !
» 0%
o
4 .
g ° SIUTH-1
® 8 T&o unirrapiaTED [T LUPTURE o
. IRRADIATED ) %
A 1% STRAIN
& RUPTURE
q
ol 1 10 100 1000 10,000
TIME (fr)

Fig. 16. A Comparison of the Creep-Rupture Properties of
Heat 5911 TH-1 at 760°C in the Irradisted and Unirradiated Conditions.

strain rate. At high strain rates, the ductility is lower at 760°C than
at 650°C and decreases with decreasing strain rate. A minimum is reached
at a strain rate of about lO%/hr and an increase in ductility occurs

as the strain rate is decreased further.

Cur metallographic studies were confined to the unirradiasted ten-
sile specimens that were aged and tested at 650 and 760°C at a strain
rate of 0.002 min~t. The fajilures were all intergranular at 760°C. At
650°C extensive intergranular cracking occurred, and the failures were
predominantly intergranular even though some areas failed by trans-
granular shear. The microstructure of a specimen from heat 5911 AW~
anneal 1 after aging and testing at 650°C is shown in Fig. 18. The
grain size is relatively coarse with stringers, probably of the MgC type.
The microstructure of heat 5911 AW-anneal 3 is shown in Fig. 19. The cold
working and recrystallization of this material has resulted in the forma-
tion of bands of small grains and a "ghost" structure in which fine
rrecipitates mark the boundaries of the original grains before recrystal-
lization. The coarse microstructure of heat 5911 AW-anneal 3A is shown
in Fig. 20. The "feathery" growth is probably fine carbides that have
reprecipitated. These were not visible in the specimen aged at 650°C.
Small amounts of the same product were visible in the cther heat 5911 AW
specimens that were aged at 760°C. The microstructure of heat 5911 TH-

anneal 1 after aging and testing at 650°C is shown in Fig. 21. The grain
Table 8. Tensile Properties of Heat 5911 AW-1 at Various Strain Rates

 

 

 

 

Specification Test Strain Stress (psi) Tlongation (%) Réduction
Number Temperature Rate in Ares
(°c) (min=1)  Yield Tensile Uniform Total (%)
5167 650 2.0 34,200 79,500 80.1 68.6 53.4
5168 650 0.5 37,000 81,700 65.5 67.2 53.6
5169 650 0.2 26,500 80,300 55.3 57.9 40.8
5170 650 0.05 26,700 75,600 41.0 42.6 36.1
5171 650 0.02 25,300 74 ,C00 38.9 40.5 32.4
5172 650 0.005 27,200 65,200 28.6 29.7 26.0
5173 650 0.002 26,900 61,900 26,5 27.6 23.6
5174 760 2.0 31,300 67,400 44 8 50.4 34.1
5175 760 0.5 25,900 68,300 36.5 38.4 30.6
5176 760 0.2 26,400 65,900 31.4 32.7 26.8
5177 760 .05 26,400 62,800 26.8 28.2 20.7
5179 760 0.02 256,600 62,500 11.2 11.7 24 .6
5180 760 0.005 26,0C0 55,600 15.6 21.9 18.4
5181 760 0.002 27,200 49,500 10.0 21.8 21.2

 

9¢
ORNL-DWG 67-7263

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

70 e : £
T T I LA
s 650°C-AS ANNEALED ; A
80 PTTTH - 850°C-AGED AT 650°C 11T 11T~ T s Bsoec |1l
'« 780°C-AS ANNEALED A |
£ 50 o TB0°C~AGED AT 760°C. 1111 T T L
£ 40 ) A e L L2 TB0°C
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GO oo 04 1 10 100 1000 10,000

STRAIN RATE (%/hr)

Fig. 17. Influence of Strain Rate on the Rupture Strain of
Heat 5911 AW-1 at 650 and 760 °C.

Y-79283 I

 

 

 

 

 

 

 

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c
X

 

 

 

103G INCHE

-
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lx

 

 

Fig. 18. Photomicrograph of the Fracture of =z Heat 5911 AW-Anneal 1
opecimen Tested at 650°C at a Strain Rate of 0.002 min~l. Btchant:
glyceria regia.
28

 

 

Y-79293 |

 

 

il

= 500X

3

 

 

(|7

Fig. 19. Photomicrographs of a Specimen from Heat 5911 AW-Anneal 3
Tested at 650°C at a Strain Rate of 0.002 min~*. Etchant: glyceria regia.
 

)

 

 

 

SIS

Fig. 20.
Tested at 760°C at a Strain Rate of 0.002 min~*+

 

 

 

 

 

 

 

 

 

 

Etchant:

 

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¥-79207

 

 

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Photomicrographs of a Specimen from Heat 5911 AW-Anneal 3

glyceria regisa.
i

 

D035 1T1IHES
!

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Fig. 21. Photomiecrograph of the Fracture of a Specimen from Hezat
5911 TH-Anneal 1 Tested at 650°C and at a Strain Rate of 0.002 min~*.
Etchant: glyceria regis.

size 1g quite similar to that of heat 5911 AW-anneal 1 {Fig. 18), but the
curved twin boundaries are indicative of the more complex and extensive
working of the tube hollow.

Heat 6252 AC was not homogenized adequately by any of the heat
Treatments used. This material was received as a section cut from an
as-cast 12-in.-diam billet. Figure 22 shows the microstructure after
anneal 2 and aging and testing at 650°C. A lamellar phase and the
typical carbide stringers have been "smeared" through the metal in a
very inhomogenous manner. Figure 23 shows that the working of treatments
2 + 3 dmproved the homogeneity but not sufficiently to yield a typical
microstructure. Figure 24 shows thal the higher temperature encountered
in anneal 24 increased the grain size, but only inereased the degree

of segregation.
31

| Y-79278

 

 

 

Y-79281

 

 

 

Fig. 22. Photomicrographs of a Specimen from Heat 6252 AC-~Anneal 2
Tested ‘at 650°C at a Strain Rate of 0.002 min~1. Etchant: glyceria regia.,.
 

Fig. 23. Photcomicrograph of the Fracture of a Specimen from
Heat 6252 AC-Anneal 2 + 3 Tested at 650°C and at a Strain Rate of
0.002 min~*. EBtchant: glyceria regis.

-/ Y-79290

Fig. 24. Photomlcrograph of The Fracture of a Specimen from Heat
€252 AC-Anneal 24 Tested at 760°C and at = Strain Rate of 0.002 min~*t.
Etchant: glyceria regia.

 

Tro 100X

 

 

 

100X

Trs

 

 
33

DISCUSSION OF RESULTS

The effects of several wvariables on the unirradiated properties of
Hastelloy N were investigated. The dramatic decrease in tensile ductility
with increasing test temperature is a characteristic of this alloy.4
However, the exact mechanism of the ductility minimum is not understoocd.
Qur results showed that the ductility is wvery strain rate dependent over
certain ranges of temperature and strain rate. TFor example, Fig. 17 shows
that the rupture strain at high strain rates is greater at 650°C than at
760°C, and at low strain rates the converse is true. The changes in
tensile properties as a result of varying the anneal and the aging treat-
ment {650 or 760°C) are relatively small for such a complex alloy.
Although adequate work was not done to support this conclusion, these
changes are probably due to variations in grain size and the concentration
and morphology of grain-boundary carbides. Under creep conditions, all
the variables investigated seem to have very minor effects (Figs. 3-8).
Although the results obtained on heat 6252 AC generally agreed with those
for the other materials, several tests exhibited low ductility. The
photomicrographs shown in Figs. 22-24 demonstrate the inhomogeneous nature
of heat 6252 AC. None of the mechanical and thermal treatments weras
adequate to homogenize thes alloy. Thus, the scatter in test results is
as expected.

The main object of these experiments was to determine the properties
of Hastelloy N after irradistion. Irradiation changed the propertieg in
the following ways:

1. At 650°C the yield stress was increased and the tensile

stress was decreased.

2. At 760°C the yield stress was unaffected and the tensile

stress was reduced.

3. The stress-rupture life was reduced, the effect being

greater at 760°C than at 650°C.

4. The minimum creep rate was unaffected at 650°C but was

increased at 760°C.

 

“H. E. McCoy, Jr., Influence of Several Metallurgical Variables on
the Tensile Properties of Hagtelloy N, ORNL-3661 (August 1964).

 

 
3

5. The rupture strain was lowered in both tensile and creep-
rupture tests, with the minimum strain occurring at a
strain rate of 1 to 10%/hr.
We shall discuss each of these observations separately, after we have
discussed high-temperature deformation in general terms.
At high temperatures there are at least two types of intergranular
failures that are reported: (a) wedge or triple point fracture and
(b) cavitation. The wedge crack forms when the stress concentration at
the end of a sliding boundary is sufficient to locally exceed the fracture
stress. Stroh?,© developed an analytical expressiocn for the value of

the critical shear stress, 1, required to form a crack under these

 

conditions
_ 12vG
where
vy = The effective surface energy
G = the shear modulus,
I, = the length of the sliding interface (normally taken to be the

grain diameter at high temperatures).
Cavitation is a term used to describe the formstion of small inter-
granular veids and the linking of these wolds to form cracks. Numerous
mechanisms have been proposed for the nucleation of these voids, however,

they all require plastic deformation.”

Oance nucleated, the cavities grow
by the ingress of vacancies. Balluffi and Seigle8 developed an expression

for the stability of the void under an applied stress

S A (2)

r cos? 9’

 

 

 

°A. N. Stroh, Proc. Roy. Soc. 2234, 404 (1954).

®A. N. Stroh, Advan. in Phys. ;18 (1957) .

"P. W. Davies and J. P. Dennis;fi, J. Inst. Metals 87, 119 (1958-59).
®R. W. Balluffi and L. L. Seigle, Acta Met. 5, 44971957).
35

where
= the applied stress,
the surface tension,

= radius of the void,

D@ R < Q
it

= the angle between the applied stress and the normal to the plane
in which the void lies.

When the stress becomes less than that required to satisfy this equality,

the void will decrease in sgize. IT the stress is greater, the void will

grow as rapidly as the supply of vacancies will allow.

Hull and Rirmer? developed an expression for the kinetics of failure
by the growth of these voids due to the stress induced diffusion of
vacancies along the grain boundaries. Their model assumed a constant
number of void nuclei that grew as spheres. The time to rupture, tr,

was related to the applied stress, o, and other factors

kTa3
r  4(D O )(g~ P)
gz

2

4 o, (3)

where

k = Boltzman's constant,

T = the absolute temperature,

a = the void spacing,
= the atomic grain-boundary diffusion coefficient,
= the width of the grain boundary

0 = the atomic volume,

P = the externally applied hydrostatic pressure.
Hull and Rimmer found in their experimental work with impure copper that
it was necessary to allow the vold spacing, a, to decrease with increasing
stress in order to obtain reasonable agreement between Eg. (3) and their
experimental data. The activation energy was found to be less than that
for bulk diffusion, so grain-boundary diffusion was assumed to be the
rate controlling process.

Although the wedge and the cavitation fractures appear distinctly

different when viewed in the optical microscope, we recently performed

 

°D. Hull and D. E. Rimmer, Phil. Mag. 4, 673 (1959).
36

some work with tungstenlo that makes us question whether they result from
basically different mechanisms. The stressed tungsten specimens were
fractured intergranularly and the surfaces replicated to study the
details of the deformation. All of the cracks secemed to be nucleated by
intergranular voids. At high stresses and relatively low temperatures,
the voids linked together to form cracks that appeared wedge-shaped in
the optical microscope. At low stresses and high temperatures, the voids
grew to quite large sizes as voids and linked together primarily by
impingement. Hence, the different appearance of cracks may be influenced
more strongly by the local stress concentrations available to propagate
them than by differences in the mechanism of crack macleation. Thus,
even though the cracks in the Hastelloy N specimens appear to be wedge-
shaped, we should still be concerned with how voids nucleate and grow in
this alloy.

Let us now consider how neutron irradiation can alter the high-
temperature deformation processes in this material. When the alloy is
irradiated at temperatures in excess of about one-half the absclute
melting point, the atoms have high mobility and the atoms displaced by
fast neutrons are able to return to their normal lattice positions.
However, many experimenters have shown that structural materials irradiated
even at high temperatures exhibit reduced rupture ductility when tested
at elevated temperatures. This effect has been correlated with the thermal
neutron exposure and more specifically with the helium that is formed by

11

the lOB(n,QJ transmutation. The boron, because of its size and low

12 The recoil

solubility, is initially segregated at the grain boundaries.
range of the o-particle produced by transmutation is 2 (ref. 13)

so most of the helium will lie in the proximity of the grain boundaries.
The solubility of the helium in the metal is very low and most of it
will be precipitated as small bubbles along the grain boundaries. The

size of a spherical gas bubble mist satisfy the following equality

 

197, 0. Stiegler, K. Farrell, B.T.M. Loh, and H. E. McCoy, Jr.,
Trans. ASM 60, 494 (1967).

11p.C.L. Pfeil, P. J. Barton, D. R. Arkell, Trans. ANS &, 120 (1965).

2D, R. Harries, J. Brit. Nucl. Energy Soc. 5, 74 (1966) .

 

3H. P. Meyers, Aktiebolaget Atcmenergl (Sweden) , Report AE-53 (May 1961).

 
37

2
P =rl: (4)
0O

'-d
It

the internal gas pressure,

the radius.

=
i

If a tensile stress, o, is now applied perpendicular to the boundary
in which the bubble lies, the bubble will grow to a new radius, r, and

the internal gas pressure will be reduced to P.

g+ P = %} . (5)

It can be shown that the critical stress, O s that must be applied for

unlimited growth to occur is given by
o =0.38 P =0.76 1, (6)
o ry

By comparing Egs. (2) and (6), it is apparent that the applied stress to
allow a bubble to grow without bound is only about one-third that required
for the growth of an empty void of the same size.

Thus, the presence of helium results in the formation of small gas
bubbles that can serve as void nuclei. These nuclel can grow by the
stress induced migration of vacancies and the applied stress necesgsary
for their growth is reduced by the presence of helium. Russel and Velal%
have attempted to develop an equation similar to that of Hull and Rimmer
[Eq. (3)] for predicting the time to rupture with the complicating factors

of a gas in the void and lenticular growth of the voids. They obtained:

kTd
0

b =57 Dpd, Qlo + (2nkTpiao) —'QKYdOp] ’

(7)

 

 

14B. Russel and P. Vela, J. Nucl. Mater. 21, 32 (1967).
38

waere
d = the bubble width perpendicular to the grain boundary,
- the diffusion coefficient (bulk or grain boundary) ,

= the number of bubbles per unit surface area,

s D o O
i

= the number of atoms within each bubble.

Although several guestions have been raissd about some of the assumptions
on which this equaticn is based,l5 the factors being considersd seem quite
reasonable and the equation is probably wvalid for gqualitative arguments.
The temperature dependence shows up primarily in the diffusion coefficient
that varies exponentially with temperature. At a given temperature, the
factors that combine to determine whether a void can undergo growth are
the applied stress, the gas pressure in the void, and the surface energy.
These three factors appear in order in the bracketed portion of the
denominator of Eq. (7).

Tet us now discuss the variocus experimental observations that were
made. The increase in the yield stress at 650°C after irradiating at
£50°C was not anticipated on the basis of previous work.1® However, the
previous work involved air-melted materials irradiated at 700°C, and =
direct compariscon cannot bs made with the results of the present study.
This increase in yield stress is probably due to the formation of fine,
generally dispersed precipitates during irradiation. Since 650°C is
only approximately 0.55 of the absolute melting point of the alloy,
there is alsc a possibility that some displacement defects such ss dis-
location loops may be present. At 760°C the yield stress is unaffected.
The tensile stresses at both 650 and 760°C are reduced by irradiation.
This does not reflect a basic difference in the shape of the initial
portion of the stress-strain diagram for the irradiated and unirradiated
specimens, but is due simply to the premature failure of the irradiated
specimen before the normal tensile stress is reached.

'The reduced vupture life under creep conditiocns is due to the effect
of helium on the processes responsible for forming wedge-shaped cracks and

for cavitation that we have just discussed in detail. The reason that the

 

*°B. Russel and P. Vela, J. Nucl. Mater. 22, 234 (1967).
oW, R. Martin and J. R. Weir, Nuecl. Appl. 1(2), 160-167 (19€5).
39

effect is greater at 760°C than at 650°C is probably due to the greater
atomic mobility at 760°C. The motion of helium to the grain boundaries,
whether as atoms or small bubbles, will be quite sensitive to temperature.
For example, studies by Bloom and Stieglert? on type 304 stainless steel
have failed to disclose any bubbles in material irradiated at 650°C,

but have shown that the same material irradisted at 800°C contains
bubbles up to 700 A in diameter.

The convergence of the stress-rupture curves for the irradiated and
unirradiated materials at low stresses can be rationalized from Eq. (2).
The lower the stress, the larger a vold must be toc be stable. Thus, the
number of void nuclel of sufficlent size to undergo extensive growth is
quite small in both irradiated and unirradiated materials and the defor-
mation processes in both approach each other at low stresses.

The reduction in ductility under tensile and creep conditions is
due basically to the accumulation of helium at the grain boundaries
causing premature fracture. At high strain rates, such as those
encountered in tensile tests, bulk deformation accounts for a large
portion of the total strain. The bulk deformation characteristics do not
seem to be affected significantly by irradiation at elevated temperatures
and the fracture strains for irradiated materials are Tairly high. As the
strain rate is decreased, the conditions become more favorable for grain-

boundary deformation®®

and the ductility of the irradiated material
decreases rapidly. The minimum fracture strain observed in a creep

test with a creep rate of 1 to lO%/hr is likely due to the worst
combination of two factors: (1) low enough stress that the ratio of grain
boundary to bulk deformation is fairly high and (2) high enough stress to
cause many of the helium bubbles to grow. Decreasing the stress level
further causes the first factor to be more detrimental and the second
factor less detrimental; the rate of change of the second factor

apparently being less since the ductility improves slowly with decreasing

stress level.

 

17E. E. Bloom and J. O. Stiegler, private communication.

185 Garofalo, p. 36 in Fundamentals of Creep and Creep-Rupture in
Metals, Macmillan, New York, 1965,

 
40

The higher minimum creep rate of the irradiated material at 760°C
is difficult to explain, particularly since the rates converge at low
stresses. Since the fraction of strain due to bulk deformation increases
with stress and the grain boundary contribution decreases, it would appear
that bulk deformation can occur more easily in the irradiated material.
However, it would seem that this would result in a decrease in The yield
stress, which was not cbserved. Thus, this result is not presently

explained.

SUMMARY AND CONCIUSIONS

We evaluated the mechanical behavior of two heats of wvacuum-melted
Hastelloy N in several metallurgical conditions before and after irra-
diation. The various mechanical-thermal treatments studied had some
small effects on the tensile properties, bubt the creep-rupture properties
were very similar.

The specimens were irradiated to a thermal dose of
2.3 x 1020 neutrons/cm2 gt temperatures of 650 and 760°C. It was found
that:

1. At 650°C the yield stress was increased and the tensile

stress was decreased.

2. At 760°C the yield stress was unaffected and the tensile

stress was reduced.

3. The creep-rupture life was reduced, the effect being

greater at 760°C than at 650°C.

4. The minimum creep rate was unaffected at 650°C but was

increased at 760°C.

5. The rupture strain was lowered in both tensile and creep-

rupture tests, with the minimum occurring at a strain rate
of 1 to 10%/hr.
These results are consistent with our understanding of how helium

introduced by the 1'OB(n,cy) reaction should influence the properties.
41

ACKNOWLEDGMENTS

This study would not have been possible without the asgssistance of
several individuals and service groups. I gratefully acknowledge the
following: J. R. Weir, E. E. Bloom, and J. O. Stiegler for reviewing
this report and making several useful suggestions; B. C. Williams,

N. 0. Pleasant, and B. McNabb for running the tensile and creep tests;
V. G. Lane, E. M. Thomas, and J. C. Feltner for data processing;

H. R. Tinch for metallographic work; the Metals and Ceramics Division
Reports Office for preparation of the manuscript; and Graphic Arts for
preparation of the illustrations.

We also are pleased to acknowledge the assistance of Atomics
International in supplying the materials used in this program and the

support of the Division of Space Nuclear Systems of the AEC.
1-3.
5.

6—25.
26.
27.
28.
29.
30.
31.
32.
33.
34.
35.
36.
37.
38.
39.
40,
41,
42.
43,

bli—4 6,
47,

80.
81.
g2.
83.
84.
85.
86.
87.
g8g.

89.
20.
91.
92.

93.
94,
95.
96.

97 .
98.

43

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