ORNL-3661.txt

 

 

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ORNL=-3661
UC=-25 — Metals, Ceramics, and Materials
TID-4500 (31st ed.)

 

INFLUENCE OF SEVERAL METALLURGICAL VARIABLES
ON THE TENSILE PROPERTIES OF HASTELLOY N

H. E. McCoy

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ORNL-3661

Contract No. W-7405-eng-26
METALS AND CERAMICS DIVISION

TNFLUENCE OF SEVERAL METALLURGICAL VARTABLES ON

THE TENSILE PROPERTIES OF HASTELLOY N

H. E. McCoy

AUGUST 1964

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

IONAL LABORATORY LIBRARIES

M

3 445k 0548242 1

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 
   

  

 

 

  

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CONTENTS

Page

ADSETACE « v v v v e e e e e e e e e e e e e e e e e e e D
Introduction . . ¢« « ¢ ¢ ¢ ¢ o 0 e e e e e e e e 4 e e 1
Fxperimental Details . . . . « « « « ¢« ¢« o 0 0 e . 2
Experimental Results . . . « . « « « + « « « o o 4
Influence of Solution Annealing Treatment . 4
Orient8tion . « « « v o v v v 0 o o v e v w40
Tnfluence of Cooling Rate . v « « « v « « « v o » . o 40
Influence of Cold WOrking . « v « o v o o o o o o o+ - 42
Influence of Carbon Content . . . . « « « « « « o .+ . 42
Influence Of AZINE « « + o o o o o o o o o o 0 o . . b4
Discussion of ReSUItS v v v v v « o o o « o o o o o« « 20
Summary and Conclusions . . « « & « o o o o o o o 0 . . 60
Acknowledgments . . « v v v e e e e e e e e e e e e e 60

RETETENCES v & v o o o o o o o o o o o o o v o 2 e e 61
 

INFLUENCE OF SEVERAI, METATLURGICAL VARIABLES ON
THE TENSILE PROPERTIES OF HASTELIOY N

 

H. E. McCoy

ABSTRACT

The tensile properties of Hastelloy N have been evalua-
ted after various heat treatments. One vacuum-melted and
four air-melted heats were studied. It was found that the
vacuum-melted material exhibited good ductility after all
heat treatments. Annealing the air-melted material to tem-
peratures in excess of 2150°F brought about significant
reductions in the minimum fracture strain exhibited by the
alloy. Holding at temperatures of about 1600°F for an ex-
tended period recovered the fracture ductility. Aging
material in the 1100 to 1200°F range that had been previous-
ly annealed at 2150°F brought about a significant reduction
in the ductility. These changes in ductility occurred with
very small changes in tensile strength.

Tt is felt that these effects can be explained in terms
of the formation of a brittle grain boundary layer along which
a crack can propagate easily at elevated temperatures. Inter-
rupting the continuity of this layer by overaging or cold
working recovers good fracture ductility. The formation of this
layer is associated with the presence of trace alloying elements.

 

TNTRODUCTION

Modern technology depends in many ways upon materials which have the
capability of sustaining loads at elevated temperatures in corrosive
environments. Because of the variety of service conditions of interest,

a number of alloys have been developed with various properties and capa-
bilities. However, because of the complexity of most of these alloys, it
is necessary that extensive metallurgical investigations be carried out to
determine whether the properties of the alloy remain suitable under the
proposed service conditions. Although these studies are costly and time=-
consuming, it is only after such studies that the alloy can be safely and

efficiently utilized.
e 2 L T TR e R T e R R T

   

One such alloy that was developed for a specific application is
Hastelloy N. This alloy is nickel-base and was chosen for use in the
MSR because it offered good resistance to corrosion by molten-fluoride
salts and possessed moderate mechanical strength.l The basic properties
of this alloy have been investigated and reported previously.2 However,
the use of numerous heats of this alloy under various sets of ciicum-
stances has revealed potential problem areas. In order that this alloy
might be better utilized for its intended purposes and for other future
applications, further studies have been conducted. These studies have -
been concerned with the influence of the following variables on the prop-
erties: solution annealing temperature, specimen orientation, aging, cold -
working, and carbon content. Tensile tests and metallographic studies
have been the principal techniques used to evaluate the influence of the

above variables.

EXPERTMENTAL DETATLS

The chemical analyses of several hsats of the Hastelloy N used in
this investigation are indicated in Table 1. Four of the heats of material
were produced by the Stellite Division of Union Carbides Corporation and
were received in the form of l/2—in.—thick prlates. One heat of material
was obtained from the Allvac Metals Company in the form of l/2-in.-diam

rod. The geometry of the test specimen used is shown in Fig. 1.

UNCL ASSIFIED
ORNL-DWG 63-6434

 

 

D
=
1

 

 

 

 

 

 

 

‘\A‘\A\XA‘ \ ‘ ):w
WYY 2 DIAM L 4/ DIAM f'
L 716 4 5/5» RADIUS

yb—i3'THREAD

ALL DIMENSIONS ARE IN INGHES

Fig. 1. Geometry of Test Specimen.
 

Table 1. Chemical Composition (by Weight Percent)
of Several Heats of Hastelloy N

 

 

 

Heat Vendor C S Mn Si Cr Mo Co Ti

2477 Allvac 0.057 0.003 0.04 0.015 7.05 16.32 0.14 0.10

5073 Stellite 0.06 0.008 0.47 0.59 6.73 16.09 0.07 0.01
Division

5074 Stellite 0.06 0.006 0.45 0.58 6.76 16.28 0.07 0.01
Division

5075 Stellite 0.07 0. 007 0.50 0.62 6.87 15.95 0.06 0.01
Division

SP-25 Stellite 0.05 0.011 0.30 0.21 6.81 16. 58 0.49 0.03
Division

Al B Fe Cu P W v

2477 Allvac 0.055 0.0008 4.25 0.10 0.008 0.47

5073 Stellite 0.01 0.006 3.89 0.01 0. 004 0.045 0.30
Division

5074 Stellite 0.02 0.001 4,05 0.01 0.003 0.04 0.28
Division

5075 Stellite 0.01 0. 008 3.84 0.01 0.002 0.04 0. 26
Division

SP-25 Stellite 0,01 0.003 4,10 0.011
Division

 

Most of the tensile tests were run in a hydraulic Baldwin testing
machine at a crosshead speed of 0,05 in./min or a strain rate of 2,5%/min.
A limited number of tests were run in an Instron testing machine at various
strain rates. All specimens were tested in air. The furnace used was of
the clamshell type and was preheated before being closed around the test
specimen. A standard equilibrating time of 1/2 hr was used for all speci-
mens to reach the desired test temperature.

All heat treatments prior to testing were carried out in an argon
atmosphere., Unless otherwise indicated, the specimens were cooled from

the annealing temperature by pulling them from the hot zone into the
water-cooled end of the furnace tube. Thermocouples were attached to
several of the specimens, and the average cooling rate down to 500°F was

200 to 500°F/min.

FXPERIMENTATL, RESULTS
Influence of Solution Annealing Treatmernt

The tensile properties of heat 5075, after annealing 1 hr at 2150
and 2300°F, are given in Table 2 and depicted graphically in Figs. 2 and
3. Both the tensile and yield strengths were lower for the material
annealed at 2300°F than for the material annealed at 2150°F except at test
temperatures of 1800°F where the reverse was noted. After annealing at
2150°F, the material exhibits a ductility minimum in the temperature range
of 1200 to 1400°F. A 2300°F anneal shifts the ductility minimum to 1600°F
and reduces the minimum ductility significantly. However, the ductility
below 1200°F is not very different for material annealed at 2150 and 2300°F.

The tensile properties of heat 5074 after annealing at 2150 and
2300°F are given in Table 3., Duplicate sets of specimens were run on the
Baldwin and Instron testing machines. The strength values obtained on
the Instron (hard) machine were consistently higher than those obtained on
the Baldwin (soft) machine for the same heat of material. However, the
differences are very small and are probably because of slight variations
in strain rate. The fracture ductilities did not show any consistent
variations. The tensile and yield strengths of heat 5074 are comparable
with those of heat 5075. However, the fracture ductility exhibits a dif-
ferent behavior. After annealing at 2150°F, the ductility minimum of
heat 5074 occurs between 1400 and 1600°F. After annealing at 2300°F, the
fracture ductility is still decreasing with temperature at 1800°F, However,
the duectility below 1200°F is not significantly different for material
annealed at 2150 and 2300°F.

The properties of heat 5073 in longitudinal and transverse orienta-
tions (with respect to rolling direction) are given in Table 4. The
strength variations with test temperature and heat treatment are comparable
with those just discussed for heats 5074 and 5075. However, the fracture
ductility shows a significant difference. After annealing at 2150°F, the

e e R i e b kg st o B b b ke ST s S tak @ e s et s n B e e

 
 

Table 2. Tensile Properties of Hastelloy N
Heat 5075
(Strain Rate: 2.5%/min)

 

 

Ultimate
Test Yield Tensile Reduction
Heat Temperature Strength  Strength  Elongation in Area
Treatment (°F) (psi) (psi) (%) (%)
a 75 43,500 113,800 50.0 53.05
a 800 31,300 99,000 55.0 51.00
a 1200 31,500 79,100 30.0 23,00
a 1400 30,400 61,600 25.0 23.36
a 1600 29,700 37,300 30.0 32.22
a 1800 21,400 21,600 41.0 36.50
b 75 40,100 109,900 58.0 51.77
b 800 26,900 93,600 59.5 54,12
b 1200 25,000 71,100 29.5 31.32
b 1400 25,100 50,300 16.0 13.15
b 1600 24,300 38,100 7.5 3.97
b 1800 22,600 22,800 26.0 35.25

 

ZAnnealed 1 hr at 2150°F in argon, fast cooled.
Pannealed 1 hr at 2300°F in argon, fast cooled.

ductility continues to decrease with test temperature through 1800°F.
After annealing at 2300°F, the fracture ductility at lower temperatures
is not changed significantly, but it decreases very rapidly above 1400°F.
The ductility continues to decrease through the highest temperature in-
vestigated, 1800°F, where a reduction in area of only 1.6% was obtained.
The tensile properties of heat SP-25 after various heat treatments
are given in Table 5. The strength of this heat is slightly less than
that of heats 5073, 5074, and 5075, Heat treatments at 2000 and 2300°F
resulted in comparable low-temperature rupture ductilities, Both heat
treatments also resulted in decreasing fracture ductilities with

increasing test temperature with the 2300°F heat treatment yielding lower
UNCLASSIFIED

 

ORNL—-DWG 64-2604

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fig. 2.
at Various Test Temperatures.

TEMPERATURE (°F)

 

(x 10°)
O~
-\-\
100 \—-< ——TENSILE STRENGTH —~—
N\{
80 ™~
2 HE AT 5075 N
= o ANNEALED 1hr AT 2150°F \ |
¢» 60 — @ANNEALED 1hr AT 2300°F
ac
|_
N o l \\\\\\\
40 o t
~NOTe—
— T — — — PN
20 /
YIELD STRENGTH
o I
0 2 4 6 8 10 12 14 16 (x10

Tensile and Yield Strengths of Hastelloy N (Heat 5075)

2)
REDUCTION IN AREA (%)

60

50

40

30

20

10

 

UNCLASSIFIED
ORNL-DWG 64-2605

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

:8‘_:::=-——=:='_:::::- \
\ /'i.
.
\ " g
o\
\
HEAT 5075 \\\
o ANNEALED 1 hr AT 2450°F .
-—— @ ANNEALED 1 hr AT 2300°F \\\\‘//
0 2 4 6 8 10 2 14 6 (x10°)

TEMPERATURE (°F)

Fig. 3. Ductility of Hastelloy N (Heat 5075) at Various Test
Temperatures.
 

 

 

 

Table 3. Tensile Properties of Hastelloy N

 

 

Heat 5074
Ultimate
Test Yield Tensile Reduction
Heat Temperature Strength Strength Elongation in Area
Treatment (°F) (psi) (psi) (%) (%)
Baldwin Machine
a 75 47,900 117,100 50.0 48,16
a 800 32,600 98,600 50.0 48.75
a 1200 27,900 75,200 32.5 34,08
a 1400 25,600 56,400 26.0 29. 44
a 1600 25,900 36,000 26.0 20,52
a 1800 20,800 20,800 50. 42,25
b 75 39,500 106,100 02, 48.16
b 800 29,100 91,700 62. 51.01
b 1200 25,700 70, 200 33. 41.77
b 1400 25,500 54,800 25.5 23.98
b 1600 24,700 36,100 .5 4,77
b 1800 21,900 21,900 .5 3.97
Instron Machine
a 75 45,100 124,100 45,4 48.90
a 800 37,600 108,600 48,5 49,47
a 1200 31,700 79,900 28.2 37.29
a 1400 33,700 63,900 32.9 32.39
a 1600 26,300 37,400 36.0 34.88
a 1800 21,700 21,800 57.9 51.38
b 75 40,800 110,200 54.7 50.37
b 800 30,400 96,300 58.1 49,55
b 1200 26,400 76,300 32.9 34.19
b 1400 26,500 61,100 21.7 22.62
b 1600 27,300 37,500 .27 7,11
b 1800 23,700 24,100 .64 4.7

 

8Annealed 1 hr at 2150°F in argon, fast cooled.
PpAnnealed 1 hr at 2300°F in argon, fast cooled.

 
 

Table 4, Tensile Properties of Hastelloy N

 

 

Heat 5073
Ultimate
Test Yield Tensile Reduction
Heat Temperature Strength Strength ZElongation in Area
Orientation Treatment (°F) (psi) (psi) (%) (%)
Longitudinal a 75 43,100 112,100 50.0 44,79
Longitudinal a 800 31,800 100,000 56.0 39.93
Longitudinal a 1200 27,900 76,000 36.5 34.20
Longitudinal a 1400 27,500 60,000 26.0 27.50
Longitudinal a 1600 27,700 36,300 21.0 20,50
Longitudinal a 1800 19,800 19,800 10.0 9.41
Longitudinal b 75 41,100 107,000 61.0 50.05
Longltudinal b 800 30,400 91,800 61.5 51.18
Longitudinal b 1200 24,600 71,100 40.0 39.93
Longitudinal b 1400 24,800 52,300 21.5 24.75
Longitudinal b 1600 26,100  34.400 .5 6.30
Longitudinal b 1800 21,200 21,200 .0 1.60
Transverse a 75 43,500 113,800 53.5 50.60
Transverse a 800 32,000 100,400 54.0 50.60
Transverse a 1200 27,500 79,700 37.5 36.00
Transverse a 1400 29,000 62,000 27.5 27,50
Transverse a 1600 27,500 36,700 27.5 26. 14
Transverse a 1800 20,600 20,600 20.0 13.19
Transverse b 75 41,100 107,400 54.5 50.60
Transverse b 800 29,000 20,700 53.5 51.90
Transverse b 1200 25,500 73,100 37.5 40. 54
Transverse b 1400 26,700 53,800 19.0 24.75
Transverse b 1600 26,600 35,700 5.8 5.55
Transverse b 1800 22,400 22,400 .0 2,40

 

“Annealed 1 hr at 2150°F in argon, fast cooled.
Pannealed 1 hr at 2300°F in argon, fast cooled.
 

 

10

 

 

 

Table 5. Tensile Properties of Hastelloy N
Heat SP-25
Ultimate
Test Yield Tensile Reduction
Heat Temperature Strength Strength Elongation in Area
Treatment (°F) (psi) (psi) (%) (%)
a 75 42,900 116,700 56.8 48,63
a 800 28,000 99,700 58.0 45,73
a 1200 26,400 73,300 31.6 32.12
a 1400 26,100 61,000 24,0 29.18
a 1600 23,500 36,500 28.0 24.79
a 1800 18,100 18,700 24.8 20. 24
b 75 37,600 101,000 64.8 47,56
b 800 24,000 85,700 65.6 43,96
b 1200 20,400 66,800 50.8 37.50
b 1400 21,000 57,100 32.0 25.45
b 1600 24,400 34,800 6.4 6.60
b 1800 20,000 20,000 .8 3.14
c 75 40,400 112,200 60.0 46,90
c 800 28,000 98,700 62.0 45.45
c 1200 27,200 77,500 40.0 34.08
c 1400 25,700 62,300 50.8 38.51
c 1600 25,200 35,000 60.8 44,55
c 1800 19,500 19,500 28.8 23.73
d 75 38,900 96,600 59.6 36.63
a 800 25,000 79,900 56.8 40.67
d 1200 22,500 61,900 44, 8 36.94
d 1400 21,000 49,900 32.0 30.51
d 1600 25,500 35,600 9.60 17.02
d 1800 19,600 19,600 3.20 7,37
ZAnnealed 1 hr at 2000°F in argon, fast cooled.
Pannealed 1 hr at 2300°F in argon, fast cooled.
CAnnealed 100 hr at 2000°F in argon, fast cooled.
dAnnealed 100 hr at 2300°F in argon, fast cooled.
 

11
ductility values. Prolonged heating at 2000 and 2300°F significantly
increased the fracture ductility at test temperatures above 800°F. These
treatments resulted in very minor changes in strength.

The tensile properties of a heat of vacuum-melted Hastelloy N,
heat 2477, are given in Table 6. The first two tests listed in the table
indicate that the mill anneal probably did not actually reach 2150°F.
This material has strength comparable with that observed for the other
heats of material melted in air. The ductility of the mill-annealed ma-
terial goes through several fluctuations as a function of test temperature
A l1-hr anneal at 2300°F

slightly reduces the ductility at high temperatures but the minimum reduc-

but the reduction in area never goes below 35%.

tion in area observed is 26%

 

 

Table 6. Tensile Properties of Hastelloy N
Heat 2477
Ultimate
Test Yield Tensile Reduction
Heat Temperature Strength  Strength  Elongation in Area
Treatment (°F) (psi) (psi) (%) (%)

a 75 38,400 109,100 55.5 62,25
b 75 61,900 121,100 45,0 59.42
b 800 46,400 103, 500 49.0 54,66
b 1200 43,300 84,500 29.0 35.28
b 1400 42,300 68,600 41,5 39.93
b 1600 31,200 37,100 83.0 78.51
b 1800 20,200 20,900 45,0 50.60
c 75 38,700 104,300 75.0 61.99
C 800 28,000 20,900 70.0 60.93
c 1200 24,400 73,300 50.0 41.03
c 1400 23,400 60,800 29.5 29. 54
c 1600 29,300 38,800 45.0 38.95
C 1800 21,200 21,400 33.5 26.23

 

“Annealed 1 hr at 2150°F in argon, fast cooled.
PMill annealed.
CAnnealed 1 hr at 2300°F in argon, fast cooled.
   

12

The individual stress-strain curves for Hastelloy N show several
interesting features. Tensile curves at 75°F are quite smooth and appear
normal. Curves at 800 to 1600°F exhibit serrations equivalent to as much
as 4000 psi and occur at frequencies as rapid as 1 serration/0.025% strain.
These serrations vary in magnitude and frequency in the course of a single
test or from one test temperature to another. However, the characteristics
of the serrations at a given strain rate and temperature are quite repro-
ducible. At 1800°F, the curves are usually quite smooth; the material
shows some strain hardening, reaches its maximum load at a very low strain,
and then the load continues to decrease during the rest of the test.

Since air-melted Hastelloy N exhibits good ductility after annealing
at 2150°F and much lower ductility after annealing at 2300°F, several
tests were run to determine how rapidly the ductility decreased with in-
creasing annealing temperature. The results of these tests are given in
Table 7. Specimens were annealed 1 hr at temperatures between 2150 and
2400°F and were tested at 1600°F, a temperature in the minimum ductility
range. As shown in Table 7, the ductility is significantly reduced by an
anneal at 2200°F over that observed after a 2150°F anneal. Annealing at
2250°F brings about a further reduction in ductility, but increasing the
annealing temperature to 2300 and 2400°F does not result in further embrit-
tlement. These changes in ductility occur with only small changes in

tensile strength.

Table 7. Influence of Annealing Temperature on the
Tensile Properties of Hastelloy N2

 

 

Heat 5075
Annealing Yield Ultimate Reduction
Temperature Strength Tensile Strength Elongation in Area
(°F) (psi) (psi) (%) (%)
2150 29,700 37,300 30.0 32.20
2200 28,100 36,800 11.0 9.33
2250 26,900 35,800 7.5 5.54
2300 27,300 38,100 7.5 3.97
2400 26,300 39,500 6.0 4.78

 

8Annealed 1 hr at indicated temperatures, rapidly cooled; test
temperature: 1600°F,

T oo ST
 

Since the fracture ductility of the air-melted Hastellloy N was

rendered quite low by pretest annealing at 2300°F, several notched speci-

mens were tested to determine the influence of a sharp notch on the frac-

ture ductility.

The notch was 0.030-in. deep and had a root radius of 0.001 to 0.0015 in.
and an included angle of 30° at the base.

the photomicrograph in Fig. 4, indicate that, after the material has been

These test results, as well as

The results of these tests are swmarized in Table &.

annealed at 2300°F, fracture can occur at a notch with no measurable

plastic strain.

 

 

Table 8. Influence of Notching on the Properties of Hastelloy N
Heat 5073
(Notch Radius is 0.001 to 0.0015 in.)
Ultimate
Test Yield Tensile Flonga- Reduction
Heat Temperature Strength Strength tion in Area
Treatment Notched (°F) (psi) (psi) (%) (%)
a No 75 43,500 113,800 53.5 50.6
a No 1400 29,000 62,000 27.5 27.5
a No 1600 27,500 36,700 27.5 26.1
a Yes 75 57,800 124,400 15.3
a Yes 1400 42,700 68,200 '
a Yes 1600 41,400 59,100 .2
b No 75 41,100 107,400 54.5 50.6
b No 1400 26,700 53,800 19.0 24.8
b No 1600 26,600 35,700 5.8 5.6
b Yes 75 49,400 115,400 22.8
b Yes 1400 42,200 58,600 6.4
b Yes 1600 41,500 44,400 0.0

 

SAnnealed 1 hr at 2150°F in argon, fast cooled.
PArmealed 1 hr at 2300°F in argon, fast cooled.
14

UNCLASSIFIED
Y-55702

 

Fig. 4. Fracture of Notched Hastelloy N Specimen Tested at
1600°F, Annealed 1 hr at 2300°F Prior to Testing. Heat 5073. Etched

in glyceria regia. 100x

 
et S5 PR T S S T T

 

15

In an effort to determine whether the changes in ductility brought
about by solution annealing were associated with microstructural changes,
several specimens were examined metallographically. Figures 5, 6, 7, g,
9, and 10 illustrate that microstructure of heat 5075 after annealing 1 hr
at 2050, 2200, 2250, 2300, 2400, and 2500°F, respectively. The stringers
of precipitate have been tentatively identified by electrolytic extraction
as carbides of the MgC type. These stringers do not appear to dissolve
at an appreciable rate at temperatures less than 2400°F. At 2500°F, the
discrete precipitate particles have dissolved, but a lamellar product is
present in the grain boundaries. There is also some evidence of melting
during the 2500°F anneal and the inhomogeneous distribution of the molten
areas illustrates the inhomogeniety of this material. The grain sizes
are equivalent after anneals at 2150 and 2200°F, The grain size is in-
creased by an anneal at 2250°F although the grain growth is reduced sig-
nificantly in areas where the stringers are present. Annealing at 2300°F
does not result in a significantly larger grain size than that obtained
at 2250°F. The influence of the stringers on the grain growth has disap-
peared at 2400°F and the grain size is quite large. However, numerous
individual precipitate particles are present which retard the moticn of
the grain boundaries and cause them to be quite irregular (see Fig. 9b).
The boundaries also etch rapidly and appear quite broad, indicating that,
possibly, an impurity layer is present. Annealing at 2500°F does not
result in additional grain growth, but the grain boundary layer and pre-
cipitate particles appear to be converted to an intergranular lamellar
product (Fig. 10).

The diamond-pyramid hardness (DPH) of the material after each anneal
ig listed with the photomicrographs in Figs. 5 through 10. It is quite
surprising that the hardness decreases with increasing annealing temperature.

Heats 5074, 5073, and SP-25 were also annealed at various temperatures.
The resulting structures were quite gimilar to those just described for
heat 5075. However, the structure of heat 2477 was quite different.
Figures 11, 12, 13, 14, and 15 illustrate the structure of this material
after annealing 1 hr at 2150, 2200, 2300, 2400, and 2500°F, respectively.
After annealing at 2150°F, the structure contains randomly dispersed

precipitates rather than the stringers found in heats 5073, 5074, 5075,
16

“ UNCLASSIFIED
Y-49313

g g

 

Fig. 5. Photomicrograph of Hastelloy N After Annealing 1 hr at
2150°F in Argon and Rapidly Cooling, Heat 5075. Hardness (DPH)
average 178; range 162-193, Etched in glyceria regia. 100X

 
17

"
1

D v PN NETED Ll e
= - " 3 \s E
=8 4t i jx
wis ] -
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/'\ i

 

  

Hardness (DPH) average 185;
100x

"

Heat 5075.
range 174-213. Etched in glyceria regia,

Photomicrograph of Hastelloy N Annealed 1 hr at 2200°F

6.
in Argon and Rapidly Cooled.

Fig.

 
c

UNC LASSIFIED
Y-49005

 

   
 

Sl e
e o .w.. v &awf..ut.__ rn

Hardness (DPH) average 179;

Photomicrograph of Hastelloy N Annealed 1 hr at 2250°F
Heat 5075.

Elg. 7.
in Argon and Rapidly Cooled.

range 166—196.

100x

Etched in glyceria regia.
 

Y-49007

o
o
w
0
W
35
O
=
=

 

Hardness (DPH) average 175;
100x

Heat 5075.

FPhotomicrograph of Hastelloy N Annealed 1 hr at 2300°F
Etched in glyceria regla,

Fig. B&.
in Argon and Rapidly Cooled.

range 172=182.
20

e ) X 0 - (¥ UNC LASSIFIED

:L&\\ = il - o ¥ 49426

 

 

- :\) UMCLASSIFIED
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X C‘JtC:J X

Fig. 9. Photomicrographs of Hastelloy N Annealed 1 hr at 2400°F
in Argon and Rapidly Cooled. Heat 5075. Hardness (DPH) average 168;
range 153-175. Etched in glyceria regia. (a) 100X, (b) 500x. Reduced 19%.

 
<l

UNCLASSIFIED
Y-49699

 

Fig. 10. Photomicrograph of Hastelloy N Annealed 1 hr at 2500°F
in Argon and Rapidly Cocled. Heat 5075. Photograph taken with specimen
in the as-polished condition. Hardness (DPH) average 149; range 140-159.
100x

 
22

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Fig. 11. Photomicrograph of Hastelloy N Annealed 1 hr at 2150°F
in Argon and Rapidly Cooled. Heat 2477. Etched in glyceria regia. 100X

 
23

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Fig. 12. Photomicrograph of Hastelloy N Annealed 1 hr at 2200°F
in Argon and Rapidly Cooled. Heat 2477. ZEtched in glyceria regia. 100X

 
24

: - ~ *  UNCLASSIFIED
: | -~ Y-5239]
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Fig. 13. Photomiecrograph of Hastelloy N Annealed 1 hr at 2300°F
in Argon and Rapidly Cooled. Heat 2477. Etched in glyceria regia. 100X

 

 
25

UNCLASSIFIED
Y-52394

 
   

Fig. l4. Photomicrograph of Hastelloy N Annealed 1 hr at 2400°F
in Argon and Rapidly Cooled. Heat 2477. ©Etched in glyceria regia, 100X

 
26

UNCLASSIFIED
Y-52398

 

Fig. 15. Photomicrograph of Hastelloy N Annealed 1 hr at 2500°F
in Argon and Rapidly Cooled. Heat 2477. Etched in glyceria regia. 100X
 

27

and SP-25. Annealing at 2200°F causes some grain growth and solution of

a significant number of the precipitates. A 2300°F anneal causes some
additional grain growth and further solution of precipitate, A 2400°F
anneal produces a structure which appears to be single phase with a few
inclusions present. The grains are equiaxed and the grain boundaries
regular, contrasted with the irregular boundaries shown in Fig. 9 for

heat 5075. Annealing at 2500°F produces further grain growth but no inter-
granular voids indicative of melting are produced. Faint traces of a
lamellar intergranular product are visible.

several of the tested specimens were examined metallographically.
Figure 16 shows the fracture of a specimen annealed at 2150°F and tested
at 75°F, The fracture was transgranular and most of the small cracks
away from the fracture appeared to have been initiated by cracks in the
precipitates. The structure resulting from a test at 800°F is quite
similar with the exception that more cracks are present. Testing at 1200°F
resulted in failure by combined intergranular and transgranular modes.
Numerous intergranular cracks are visible away from the fracture, and
cracking of the precipitate continues to occur at this temperature. Testing
at 1400°F produces an intergranular fracture and numercus intergranular
cracks. The precipitate particles do not crack very frequently at this
temperature but cracks are initiated by separation at the precipitate-
matrix interface. The structure resulting from testing at 1600°F is shown
in Fig. 17. It is quite striking that the intergranular cracks are pre-
dominantly normal to the applied stress and cover the entire grain width.
The structure obtained by testing at 1800°F 1s shown in Fig. 18. The
specimen is heavily cracked and recrystallization has occurred. This
specimen exhibited high rupture ductility.

Specimens tested at 1600°F after annealing 1 hr at 2200, 2250, and
5300°F are shown in Figs. 19, 20, and 21, respectively. Although the
ductility showed a large reduction as a result of increasing the annealing
temperature, there are no striking differences 1in microstructure except
the increase in grain size. However, since the intergranular cracks extend
the entire grain width, the change in grain size cannot be ignored. A

specimen annealed 100 hr at 2300°F and tested at 1600°F is shown in Fig. 22.
0}

  
   

2 { L UNCLASSIFIED
{i 3 Y-48946
1
!: - J
r \L .

S

g
oy
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bon v iy _,, & s | ¥ R o i .
; I I SR oy R TR
| it 3 . '

Fig. 16. Fracture of a Hastelloy N Specimen Tested at 75°F and
Annealed 1 hr at 2150°F Prior to Testing. Heat 5075. Etched in glyceria
regia. 100x

 

 
UNC LASSIFIED
Y-48975

UMC LASSIFIED
Y-48974

 

Fig. 18. Hastelloy N Specimen Tested at 1800°F and Annealed 1 hr
at 2150°F Prior to Testing. Heat 5075. Etched in glyceria regia.
(a) 100x, (b) 500%. Reduced 17%.

 
31

® o« UNCLASSIFIED
9 Y -48992

® v

 

\
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& 1- g = v € 4
2 0t
3 . S r

Fig. 19. PFracture of & e
Annealed 1 hr at 2200°F Prior to T
regia, 100X

® H
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o

 
UNCLASSIFIED
Y-48%95

 

Fig. 20. Fracture of a Hastelloy N Specimen Tested at 1600°F and
Annealed 1 hr at 2250°F Prior to Testing. Heat 5075. Etched in glyceria
regia. 100X

 
     
 

.

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Fig. 21. PFracture of a Hastelloy
Annealed 1 hr at 2300°F Prior to Testing,
regia. 100X

N Specimen Te

1600°F and

]
N
ct
m
3
k
[
p

J

Heat 5075. BEtched in glyceria

 
  

UNCLASSIFIED
Y-408999

A

=
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g3

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3
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By s
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po UNCLASSIFIED
.‘ Y-4899R8
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=

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o'
S
—

 

Fig. 22. Hastelloy N Tested at 1600°F and Annealed 100 hr at 2300°F

Prior to Testing. Heat 5075. Etched in glyceria regia. (a) 100X,
(b) 500% . Reduced 16%.

 
35

This anneal has resulted in some solution of the precipitates and the
formation of patches of a lamellar grain boundary phase. A substructure
is also evident, probably having been decorated by carbon precipitation
during the 1600°F tensile test.

Although all of the microstructures used to illustrate the structure
of tested specimens have been of heat 5075, the structures of heats 5073,
5074, and SP-25 were quite similar with only one noteworthy exception.

The data in Tables 2, 3, 4, and 5 show that some heats under certain condi-
tions went through a ductility minimum and then exhibited good ductility
at 1600 to 1800°F. Tt was found by metallographic studies that this re-
covery in ductility was always associated with recrystallization. To
illustrate, heat 5075 exhibited good ductility at 1800°F after having been
annealed at 2150°F and the microstructure in Fig. 18 indicates that
recrystallization occurred. Heat 5073T, after being annealed at 2300°F,
continued to show decreasing ductility with temperature at 1800°F, and
Fig. 23 indicates that recrystallization did not occur during test.

Heat 5074 also exhibited poor ductility at 1800°F after a 2300°F anneal
and was found not to be recrystallized. However, after a 2150°F anneal,
heat 5074 was ductile at 1800°F and Fig. 24 shows that recrystallization
did occur during the test.

Tested specimens of heat 2477 (vacuum melted) exhibited structures
quite similar to those of the other heats at temperatures up to about
1400°F. TFigure 25 shows the microstructure of a specimen annealed at
2300°F and tested at 1400°F. The reduction in area of this specimen was
about 30%. TFigure 26 shows the microstructure of a specimen annealed at
2300°F and tested at 1600°F. The reduction in area was 39%. Instead of
recrystallization occurring, extensive polygonization has resulted. Hence,
it appears that the recovery of fracture ductility at elevated temperatures
occurs through recrystallization in the air-melted heats and through

polygonization in the vacuum-melted heat.
UNCLASSIFIED
¥-55701

 

Fig. 23. Microstructure of Hastelloy N Tested at 1800°F and
Annealed 1 hr at 2300°F Prior to Testing. Heat 5073T. Btched in glycerisa
regia. 100X
37

UNCLASSIFIED
Y=55703

 
  
  

-

1800°F and Annealed 1 hr
EMri

5074. Etched in glyceria regias. 100X

o

rig, 24 Fracture of Hastelloy N Tested at
 

UNCLASSIFIED
Y-55697

 
   

Fig. 25. Fracture of Hastelloy N Tested at 1400°F and Aunneeled 1 hr
at 2300°F Prior to Testing. Heat 2477. Etched in glyceria regia. 100X

 
I
-
=g T .-;f:—fi' 779 . F‘i‘ + 1..1 a [.i : r i“- '_|_:‘!_. ceria reg ia . '14' )2 -

UNCLASSIFIED
Y-55698

    
 
  
    

   

. Jr y 41, !
t ‘I:‘__,;" r =3 j ; .
_np'\‘, . I,i Vil

lmgtelloy N Annealed at 1600°

  

Fracture of

it

3
40

Orientation

Because of the stringers that were present in the air-melted heats,
it was thought probable that the properties might be quite different in
directions normal and parallel to the stringers. The tests summarized
in Table 4 for heat 5073 indicate that the tensile properties of this

meterial are not influenced by specimen orientation.

Influence of Cooling Rate

The normal cooling rate used after the pretest anneal was about
500°F/min down to about 500°F. Several tests were run to determine what
influence varying the cooling rate had on the tensile properties. The
results shown in Table 9 indicate that the fracture ductility is improved
by reducing the cooling rate from the pretest annealing temperature of
2300°F. The microstructure of the specimen cooled at a rate of 2°F/min
is shown in Fig. 27. This structure should be compared with that shown
in Fig. 21 where the cooling rate was about 500°F/min. The flow lines in
Fig. 27 are quite significant and indicate that the material was able to

absorb large amounts of plastic deformation,

Table 9. Effect of Cooling Rate on the Tensile Properties
of Hastelloy N

Heat 5075

(A1l specimens annealed 1 hr at 2300°F and cooled
at indicated rate; tested at 1600°F)

 

Ultimate
Cooling Yield Tensile Reduction
Rates Strength Strength FElongation in Area

(°F/win) - (psi) (pei) (%) (%)

 

200-500 27,300 38,100 7.5 3.97
8.1 22,800 37,700 32.0 26.32
4.2 22,400 36,700 39.0 34.08

2.0 22,800 37,700 32.0 26.32

 

 
     
  

UNCLASSIFIED
Y-51062

 

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% D - - T = . Le o
R > T e e R O
- & > ';' % e 5 » . - :
* - -~ 5
»

——

- e e
- L m -

  

b - ;-!r g - Yor
= A -.]-.----u-'r-'-"m'_h."-"-\-"“':;.'-._ ok —_—s ‘6_ s - ‘
‘ P Ve S Salis i &

l."-a'_';’-l-:"" __f,—.ol ‘j_:f.--r:fl;‘\l:i_' e L-_.- ‘

5 r N
b L e Jo T, % w DL L % . _-. ,-. ¥ - - P ns
fosoy PRI O e RN e o v
v \.'.‘n- - 5 ._'j_"-'-'-"'q'--; : Y i = = T i % T =iy -+
e ‘u' - 2 o i Mo
: g & s e - WS ER 2
- i - - T 0 e —r
=¥ ay ey -* = — L —_ o
e e = . 3 & - o > i)
- . 4 s PN A g
‘ o # - - — TN e
e oUW 0 ER S R TR e e
- 3 2 e
i ool t
-,
JH '
~ &
- e - o AN - - E )
e e by et e A,
—_ o & e = - -y 2 -
- CT am L > i -"-_T"—- = .'--._., r‘x\ '
’ = . ..fl.--.A_,_.:- e 2l £ ~ = - —

Fig. 27. Fracture of a Hastelloy N Specimen Tested at 1600°F;
Annealed 1 hr at 2300°F and Cooled 2t a Rate of 2°F/min Prior to Testing.

-

Heat 5075. Etched in glyceria regia. 100X

 
42

Influence of Cold Working

Since high-temperature pretest anneals which produce low fracture
ductilities at 1200 to 1800°F do not lower the ductility at 75°F, experi- .
ments were run to see what influence working at 75°F would have on the
ductility in the 1200 to 1800°F range. Typical results are shown in
Table 10. The improvement in fracture ductility brought about by pre-
straining at 75°F is quite significant. A photomicrograph of the specimen
strained 20% is shown in Fig., 28. The microstructure is quite similar to
that of a specimen which had not been prestrained (Fig. 21). The only
significant differences are the flow lines present in the prestrained

specimen.

Table 10. Influence of Cold Working on the Tensile Properties
of Hastelloy N

Heat 5075

(Annealed 1 hr at 2300°F in argon, rapidly cooled;
J
prestrained indicated amount at 75°F;
test temperature: 1600°F)

 

Ultimate
Yield Tensile Reduction
Prestrain Strength Strength Elongation in Area

(%) (psi) (psi) (%) ®)

 

27,300 38,100 7.5 3,97
34,000 35,200 10. 5 10. 04
20 33,100 34,100 26.5 26. 14

 

Influence of Carbon Content

Several specimens were decarburized in Hp-Hp0 gas to determine the
influence of carbon content on the fracture ductility. Typical results
are shown in Table 11. A reduction in carbon content from 0.068 to
0.015 wt % resulted in very small increases in the fracture ductility at
both 75 and 1600°F; the strength was also reduced slightly. The micro-

structure was apparently not modified by the loss of carbon. Hence, it
     

UNCLASSIFIED
Y«51057

 

Fig. 28, Fracture of Hastelloy N Specimen Tested at 1600°F, Annealed
1 hr at 2300°F, and Strained 20% at 75°F Prior to Testing. Heat 5075.
Etched in glyceria regia, 100X

 
 

 

4ty

Table 11. Influence of Carbon Content on the Tensile Properties
of Hastelloy N

Heat 5075
(Annealed 1 hr at 2300°F in argon, rapidly cooled)

 

 

Ultimate

Carbon Test Yield Tensile Reduction
Content Temperature Strength Strength Elongation in Area
(wt %) (°F) (psi) (psi) (%) (%)
0.068 75 40,100 110,000 58.0 51.77
0.015 75 40, 200 104,000 61.5 59.75
0.068 1600 27,300 38,100 7.5 3,97
0.015 1600 24,400 34,900 8.5 13.27

 

seems that the carbon concentration would have to be reduced to very low

levels in order to improve the ductility significantly.

Influence of Aging

Since annealing studies had shown that all of the precipitate could
be dissolved in heat 2477, the possibility of embrittlement as a result
of reprecipitation was considered. Specimens were annealed at 2400°F and
aged for 100 hr at 1500, 1700, and 1900°F. At 1500°F, the structure shown
in Fig. 29 resulted. At 1700°F, some precipitation occurred along twin
boundaries, and at 1900°F no visible reprecipitation occurred.

Several tensile tests were run to determine whether this reprecipi=-
tation altered the tensile properties. The results of these tests are
summarized in Table 12. Aging for 100 hr at 1200°F had very little, if
any, effect on the tensile properties. However, cold working the speci-
mens prior to aging at 1200°F increased the strength by an amount which
increased with increasing cold working. The resulting structure after
annealing at 2400°F, cold working 10%, and testing at 1400°F is shown in
Fig. 30. The precipitate formed under these conditions is too fine to be

resolved at a magnification of 1000 diameters, Annealing for 100 hr at
45

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- . L

Fig., 29. Photomicrograph of Hastelloy N Annealed 1 hr at 2400°F
and Aged 100 hr at 1500°F. Heat 2477. Etched in glyceria regia. 100X

 
03

Fig, 30,
Annealed 1 hr
to Testing.

Sl

Heat

 

UNCLASSIFIED
Y -55699

 
 

47

Table 12. Tensile Properties of Hastelloy N After
Various Aging Treatments

 

 

Heat 2477
Test Ultimate
Treatment Tempera- Yield Tensile Elonga- Reduction
Prior ture Strength Strength tion in Area
to Testing (°F) (psi) (psi) (%) (%)
75 38,400 109,000 55.5 62.2
1400 42,300 68,600 41,5 39.9
c 75 38,700 104,000 75.0 62.0
c 1400 23,400 60,800 29.5 29.5
100 hr at 75 36,300 101,000 69.5 58.2
1200°Fd
100 hr at 1400 60,600 33.5 38.7
1200°F4
Cold Worked 10% 1400 44,300 70, 500 28.0 22.7
and 100 hr at
1200°F¢
Cold Worked 20% 1400 60,500 77,500 33.0 20.6
and 100 hr at
1200°Fd
100 hr at 75 50,800 112,000 45,0 45,2
1500°Fd
100 hr at 1400 31,400 64,700 23,5 30.9
1500°F<

 

9Annealed 1 hr at 2150°F in argon, fast cooled.
PAs-Received; mill annealed at 2150°F.

CAnnealed 1 hr at 2300°F in argon, fast cooled.
dAnnealed 1 hr at 2400°F in argon, fast cooled.

1500°F does not result in strengths as high as observed after aging at
1200°F., The microstructure of a specimen aged at 1500°F and tested at
1400°F 1s shown in Fig. 31. The most important fact is that the changes
in strength produced by aging this alloy do not result in fracture duc-
tilities low enough to be of concern.

Further aging studies were carried out on air-melted heats 5074 and

5075. The results of these studies are summarized in Table 13. All
o
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R
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10V

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of Ha

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Fract

Gehed

 
 

 

49

 

 

Teble 13. Influence of Aging on the Tensile Properties of Hastelloy U
Test Temperature 1600°F
Ultimate
Yield Tensile Reduction
Heat Strength Strength Elongation in Area
Treatment (psi) (psi) (%) (%)
Heat 5074
a 25,900 36,000 26.0 20.5
10 hr at 1100°F" 26,700 36,600 17.0 19.8
100 hr at 1100°F* 26,300 35,700 12.5 13.2
1000 hr at 1100°F* 30,600 36,200 14.0 11.7
10 hr at 1200°F2 29,700 35,800 14.5 14.6
100 hr at 1200°F% 28,300 37,000 13.0 12.4
1000 hr at 1200°F 29,700 35,800 39.5 39.8
b 24,700 36,100 8.5 4.8
1000 hr at 1100°FP 26,200 32,900 .0 4.0
1000 hr at 1200°FP 23,800 36,400 47.5 47.0
Heat 5075
a 29,700 37,300 30.0 32.22
100 hr at 1100°F 27,500 34,900 16.5 13.18
1000 hr at 1100°F& 28,200 33,600 12.5 10.25
100 hr at 1400°F* 27,600 37,600 22.5 25.63
100 hr at 1600°F* 27,700 37,200 19.0 26.23
100 hr at 1800°F 26,600 36,700 25.5 27.30
b 27,300 38,100 7.5 3.97
1 hr at 1600°F° 26,500 37,800 11.5 9. 44
4 tr at 1600°F° 25,800 38,100 23.5 18.43
16 hr at 1600°F° 24,400 38,700 33.0 26.05
100 hr at 1600°F° 24,100 38,100 29.0 05.71

 

8Annealed 1 hr at 2150°F in argon, fast cooled.
Pannealed 1 hr at 2300°F in argon, fast cooled.
50

specimens were tested at 1600°F after aging. Aging at 1100°F after
annealed at 2150°F reduces the ductility by almost a factor of 2 over s
period of 1000 hr., The strength shows a slight increase over this same
time period. Figure 32 shows the microstructure of a specimen annealed
at 2150°F, aged 100C hr at 1100°F, and tested at 1600°F. The structure
locks quite similar to that shown in Fig. 17 for an unaged specimen,
Aging at 1200°F produces a decrease in the ductility by about a factor of
2, but the required aging times are much less than at 1100°F. After
1000 hr at 1200°F, the material has apparently overaged and the fracture
ductility is quite high. Figure 33 is a photomicrograph of a specimen
annealed at 2150°F, aged 1000 hr at 1200°F, and tested at 1600°F. The
specimen has recrystallized during the test. If the material is initially
annealed at 2300°F and aged at 1100°F for 1000 hr, no apparent change in
properties occurs. However, aging at 1200°F for 1000 hr results in a very
high fracture ductility.

Heat 5075 gave results very similar to those of heat 5074. After
annealing at 2150°F, aging at 1100°F reduced the fracture ductility.
Aging for 100 hr at 1400, 1600, and 1800°F did not significantly alter
The fracture ductility. After annealing at 2300°F, it was possible to
improve the ductility by aging at 1600°F prior to testing at 1600°F. As
can be seen from the data in Table 13, the ductility becomes progressively
better as the aging time at 1600°F increases. Figure 34 shows the fracture
of a specimen annealed at 2300°F, aged for 100 hr at 1600°F, and tested
at 1600°F. The structure is quite similar to that of the unaged specimen
(Fig. 21). The only significant difference is that the cracks are shorter

and more numerous in the aged speckmen than in the unaged one.

DISCUSSION OF RESULTS

The results that have been presented indicate that the tensile prop-
erties of Hastelloy N can be altered significantly by various pretest
anneals. Annealing at temperatures in excess of 2150°F lowers the frac-
ture ductility at test temperatures in excess of 1200°F. The changes
brought about at lower test temperatures do not appear to be of concern,

Some of the heats of material exhibit a ductility minimum with temperature

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» -i =k s > -.' ? v -. “I ._ 0 f ,._L' # ! an ‘:
BT ey X XSy Yo W T
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A 1%‘__ U :lf A I o = Al - ' = Ir_, - § - r = _!_ .4

Fig. 33. Photomicrograph of a Hastelloy N Specimen Tested at 1600°F,
Annealed 1 hr at 2300°F, and Aged t

1000 hr at 1200°F Prior to Testing.
Heat 5074. BEtched in glyceria regia, 100X

 
53

UNCLASSIFIED
Y-48990

 
   

T - TV il P - . if_ -

Fig. 34. Photomicrograph of Hastelloy N Specimen Tested at 1600°F,
Arrieas] an 7 1} =54 eFalatslrn L ~ A 3 M + = o =
Annesled 1 hr at 2300°F, and Aged 100 hr at 1600°F Prior to Testing.

Hano BN 5 - mlhad : i
Heat '75. Btched in glyceria regia. 100X

 
54

and then good ductility at higher temperatures, whereas the ductility of
other heats continues to decrease with increasing temperature through
1800°F, the maximum temperature investigated. The air-melted heats showed
reductions in area at fractures of less than 2% under scme conditions.

The vacuum-melted material was more ductile with g minimum reduction in
area of 26% being obtained.

Metallographic studies revealed that the microstructure of the air-
melted material consisted of large precipitate particles dispersed in
stringers in a solid solution matrix. These particles were identified as
MgC type carbides. As the pretest annealing temperature was increased
above 2150°F, some grain growth occurred, the precipitate particles began
to dissolved, and the grain boundaries became broader and etched more
readily. At 2500°F the precipitate particles were completely dissolved
and a lamellar grain boundary phase appeared. To determine whether the
lamellar phase formed directly from the MgC Precipitates, the end of a
wire was induction melted and cooled very rapidly. The photomicrograph
in Fig. 35 shows the transformation taking place. The lamellar phase was
not isolated and identified. Hence, it is not known whether it has a
stolchiometric ratio different from M¢C or whether it simply
represents another distribution of the same phase. Neither is it known
whether the formation of this lamellar grain boundary phase is indicative
of grain boundary melting or whether this is a solid-state transformation.
The vacuum-melted material appeared quite similar to that of the air-
melted material after annealing at 2150°F. However, the carbides appear
to dissolve more rapidly in the vacuum-melted material and a single-phase
microstructure is obtained by annealing at 2400°F. Metallographic studies
have also shown that the recovery of fracture ductility at elevated tem-
peratures is associated with recrystallization in the air-melted material
and with polygonization in the vacuum-melted heat,

Further studies have shown that after Hastelloy N has been annealed
at 2300°F, which should produce a very low minimum fracture ductility, it
is possible to recover the ductility by several technigues. Cooling the
material very slowly will assure good ductility upon reheating. Also,

aging at a temperature in the 1600 to 1800°F range for a few hours seems
55

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Fig. 35. Hastelloy N Wire Melted on One End. Heat 5075. Etched

in glyceria regia, (a) 100X, (b) 500x. Reduced 15%.

 
56

adequate. Working at a low temperature where the ductility is good before
testing at elevated temperatures is also quite effective.

Although Hastelloy N is not an age hardenable alloy in the same sense
as Inconel X and other materials which contain titanium and aluminum and
form intermetallic compounds with nicke1,3 it does possess some aging
tendencies. The ability to recover ductility after annealing at 2300°F
by aging at 1600°F has already been pointed out. There is also a tendency
to embrittle Hastelloy N after a 2150°F anneal by annealing in the 1100 to
1200°F range. The minimum ductility can be lowered by a factor of 2.

Although 1t is possible to cause wide variations in the fracture
ductility of this material, the changes in microstructure are quite small.
If a precipitate is formed during aging, it is too fine to be resolved at
1000%. It would also seem that such a fine precipitate would cause sig-
nificant changes in the strength as well as the ductility. The most
distinct microstructural features of the specimens tested at elevated
temperatures are the large intergranular cracks. If the ductility is low,
there are only a few very large intergranular cracks present, If the
fracture ductility is high, the cracks are more numerous and shorter.

In attempting to explain the behavior of Hastelloy N, one cannot help
but observe that numerous other metals and alloys have been known for some
time to exhibit ductility minima and that the attempts to explain these
results have been only partially successful.* Since experimenters have
not succeeded in this area with metals such as pure copper, it is indeed
a. large undertaking to attempt to rationalize the behavior of Hastelloy N,
a far from simple alloy. Before attempting this, a brief review of the
general thoughts in the area of fracture ductility are in order.

Ductility minima have often been accounted for in terms of strain
aging. Glen,” for example, has shown that various strain aging effects
occur during tensile testing of steels, each associated with the presence
of a particular alloying element. This minimum in ductility is accompanied
by a maximum in yield strength, and also, serrated stress-strain curves
are obtained. Hastelloy N exhibits serrated stress-strain curves over the
temperature range of 800 to 1600°F and often small maxima in yield strength
occur in the same temperature range. However, these serrations do not seem

to be influenced by the variables which alter the rupture ductility. Hence
57

although strain aging due to different impurities occurs over certain
temperature ranges, 1t is felt that strain aging 1is not responsible for
the low ductility of Hastelloy N.

Rhines and Wray4 have attributed brittleness in nickel to the occur-
rence of grain boundary shearing and associated intergranular rupture.
According to this theory, cavities or voids, nucleating at four-grain inter-
sections, grow intergranularly as long as they lie along shearing grain
boundaries, thus decreasing the effective cross section of the metal and
lowering its fracture ductility. The effectiveness of this process is
opposed by the ability of the grain boundaries to move either by grain
boundary migration or by recrystallizatiomn. Ductility minima occur under
combinations of temperature and strain rate which result in a critical
relation between the rates of void growth and the rate of grain boundary
migration, Smith® also looked at the ductility minimum in nickel, Although
he concurred with Rhines on several points, he felt that the intergranular
voids were nucleated at grain boundary voids by a process proposed by
Gifkins’ which involved the formation of grain boundary jogs due to slip
steps. Smith proposed that the opposing process wWas the ability of the
grain boundaries to migrate and to straighten out these jogs before grain
boundary shearing stresses sufficient to nucleate a void were developed.
However, the voids noted by both Rhines and Smith were gquite large and
were either round or elliptical in section. Such features were not
observed in the Hastelloy N specimens tested at elevated temperatures in
the present study; instead, the cracks were wedge-shaped and extended the
entire width of the grain boundary segment normal to the applied stress.
However, recrystallization and polygonization were effective in recovering
good ductility.

Grain boundary segregation of impurities has also been proposed to
account for poor ductility. Bieber and Decker® have investigated the
effects of a number of minor constituents on the ductility of nickel. Ad-
ditions of several elements in concentrations up to 0.5 wt % were made and
their influence on the rupture ductility evaluated. Noticeable loss of
ductility between 1000 and 1500°F was reported for additions of as little
as 0.0005 wt % of some of the impurities. The embrittlement was presumed

to result from the fact that these elements, because of thelr incongruous
 

58

slzes and valences, tend to preferentially locate at grain boundaries,
either in solution or as precipitatesor films. However, the authors

made no suggestion as to why the ductility was recovered at higher tem-
peratures. Olsen EE_El-g observed a minimum in the duetility vs heat-
treating temperature curve for cold-worked and recrystallized nickel
containing 0.0009 wt % sulfur. The suggested explanation was that nickel
sulfide formed early in the recrystallization Process, and as growth
proceeds, the sulfide is swept up and concentrated in the moving boundaries,
At stlll higher temperatures, the sulfide would dissolve and diffuse into
the matrix.

A slightly different impurity mechanism has been proposed to account
for the ductility of Nimonic 80A.1° During aging, a chromium-rich carbide
is formed in the grain boundary. This depletes the adjacent area in
chromium which increases the solubility for titanium and aluminum in these
areas and results in Ni3(Al1,Ti) not being precipitated in a band adjacent
To the grain boundary. This "soft" area is able to absorb large amounts
of strain, and very high ductilities are obtained. Although Hastelloy N
does not contain significant amounts of titanium or aluminum, such a
mechanism may be important with respect to carbide precipitation.

Although none of the mechanisms of embrittlement adequately describe
the behavior of Hastelloy N, it is felt that the embrittlement of this
alloy by impurity segregation is the general mechanism. The crack geometry
(wedge-shaped rather than circular), the width of grain boundaries, the
ease of etching of brittle specimens, and several other metallographic ob-
servations point toward the formation of a brittle grain boundary layer,

The influence of the several variables investigated can be rationalized
as follows. Heating to progressively higher temperatures results in the
migration of impurities to the grain boundaries. When the material is
tested at elevated temperatures where the grain boundaries must sustain
high shear stresses and crack propagation is quite rapid, the material
fractures at very low strains. Apparently, temperatures above 2150°F
result in the formation of this brittle grain boundary layer. This layer
becomes more pronounced as the annealing temperature is increased. If the
material is cooled very slowly or aged at a lower temperature, the grain

boundary layer agglomerates and a propagating crack cannot advance as

 

 
59

quickly. Cold working the material prior to testing would produce grain
boundary steps or jogs that would also interfere with the propagation of
a crack. Aging studies on material annealed at 2150°F indicate that this
grain boundary layer can also be developed by aging in the 1100 to 1200°F
range.

Since the ductility of the vacuum-melted material is much superior
to the arc-melted material and since there are varilations among the arc-
melted heats, it would appear that impurity elements are very important.
The fact that reducing the carbon content improves the ductility and the
metallographic observation that an intergranular carbide phase 1s formed
at elevated temperatures suggest that carbon is an important impurity.
The chemical analysis in Table 1 does not show any marked differences in
carbon content between the heats. However, it may be that the presence
of another impurity controls the behavior of the carbon. Significant
differences in the analyses of the air- and vacuum-melted heats exist for
sulfur, manganese, silicon, titantium, aluminum, and tungsten. The influ-
ence of these impurities on the fracture ductility should be investigated.

In considering the significance of results of this type, it is well
to consider their practical importance. Several factors seem important:

1. When this material is being used in the mill-annealed (2150°F)
state, prolonged use in the 1100 to 1200°F range will reduce the ability
of the material to deform.

2. During welding, the material will be subjected to temperatures
in the 2200 to 2800°F range. This will produce areas which are quite
brittle. Good ductility can be recovered by slow cooling, reannealing at
about 1600°F, or by cold working. The method used would be determined
from practical considerations.

3. Creep tests are needed to define the ductility as a function of
strain rate. However, in many structures, loading rates of the order of

those used in this study are encountered,
S A 30 0 AN o et SR e N B4~ 9 i A

60

SUMMARY AND CONCLUSIONS

It has been shown that the tensile properties of Hastelloy N can be
altered significantly by heat treatment. Although small variations in
strength are observed, it is felt that the influence on ductility is the
most important. A vacuum-melted alloy was found to exhibit good ductility
under all conditions studied. Four heats of air-melted material were
found to exhibit quite low fracture strains in the 1400 to 1800°F range
after annealing at temperatures in excess of 2200°F. Aging at 16C0°F,
very slow cooling rates, and cold working were effective in recovering
the ductility. The recovery of good ductility at 1800°F in some heats
was assoclated with recrystallization. After annealing at 2150°F, it was
found that the high-temperature ductility was reduced by aging in the
1100 to 1200°F range.

It 1s felt that these effects can be explained in terms of the forma-
tion of a brittle grain boundary layer along which a crack can propagate
easily at elevated temperatures. Breaking up the continuity of this layer
by aging or cold working recovers good fracture ductility. The formation

of this layer is associated with the presence of trace alloying elements.

ACKNOWLEDGMENTS

The author wishes to acknowledge the assistance of several Personnel
in the Metals and Ceramics Division. The anneals were carried out by
K. W. Boling, B. McNabb, and L. I. Heatherly. The tensile tests were per-
formed by C. W. Dollins. The metallographic work was performed by
M. D. Allen, E. D. Bolling, and H. R, Tinch. T. C. Smith was of consider-
able help in the early parts of this study when he was employed as a
summer student. Assistance is also acknowledged of the Graphic Arts
Department and the Metals and Ceramics Division Reports Office in the
preparation of this document. The author is also indebted to J. R. Weir,
H. Inouye, and D. A. Douglas who reviewed this manuscript and made several

helpful suggestions.

 
10.

 

6l

REFERENCES

T. K. Roche, The Influence of Composition Upon the 1500°F Creep-

 

Rupture Strength and Microstructure of Molybdenum-Chromium-Iron-
Nickel-Base Alloys, ORNI-2524 (June 24, 1958).

 

 

R. W. Swindeman, The Mechanical Properties of INOR-8, ORNL-2780
(Jan. 10, 1961).

 

Murray Kaufman, Trans. Met. Soc. AIME 227, 405 (1963).
F. N. Rhines and P. J. Wray, Trans. Am. Soc. Metals 54, 117 (1961).
J. Glen, J. Iron Steel Inst. (London) 190, 30 (1958).

T, C. Smith, Investigation of the Elevated Temperature Ductility

 

Minima in Nickel, Thesis, University of Tennessee (June 1963).

 

R. S. Gifkens, Acta Met. 4, 98 (1956).
C. G. Bieber and R. F. Decker, Trans Met. Soc. AIME 221, 629 (1961).

K. M. Olsen, C. F. Larkin, and P. H. Schmitt, Trans Am. Soc. Metals
53, 349 (1961).

M. J. Fleetwood, J. Inst. Metals 90, 429 (1961).
              

                

g
1-3.
5-6.

7—26.
27.
28.
29.
30.
31.
32.
33.
3.
35.
36.
37.

39.
4044,
45.
46,
4'7.

T 2 PR B AT R AT N AR et s o b ekt e i

63

ORNL-3661

UC-25 — Metals, Ceramics, and Materials
TID-4500 (31lst ed.)

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