ORNL-TM-3321.txt

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OAK RIDGE NATIONAL LABORATORY

operated by
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for the

U.S. ATOMIC ENERGY COMMISSION
ORNL- TM- 3321

 

THERMAL STABILITY OF TITANIUM~-MODIFIED
HASTELLOY N AT 650 AND 760° C

C. E. Sessions and E, E. Stansbury

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THS DO%UN%LASSW\ED 10N
DIVISION OFf LASSIFICA _
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NOTICE This document contains information of a preliminery nature
and was prepared primarily for internal use at the Oak Ridge National
Laboratory. It is subject to revision or correction and therefore does
not represent o final report.

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This report was prepared as an account of work sponscred by the United
States Government. Neither the United States nor the United States Atomic
Energy Commission, nor any of their employees, nor any of their contractors,
subcontractors, or their employees, makes any warranty, express or implied, or
assumes any legal liability or responsibility for the accuracy, completeness or
usefulness of any information, apparatus, product or process disclosed, or
represents that its use would not infringe privately owned rights,

 

 
 

 

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ORNL-TM-3321

" Contract No. W-7405-eng-26

METALS AND CERAMICS DIVISION

THERMAL STABILITY OF TITANIUM-MODIFIED o
HASTELLOY N AT 650 AND 760°C

C. E. Sessions and E. E. Stansbury

 

| ‘This - report’ was prepared as an account of work
-| ‘sponsored by the -United ‘States Government. Neither |-
1 the United States nor the United States Atomic Energy
| Ccommission, nor any of their employees, nor any qf _
-1 their contractors, subcontractors, or their -employees, |
| makes any warranty, express or implied, or assumes any
i .| legal Habllity -or responsibility for the sccuracy, com-
t | -pleteness of usefulness of any information, apparatus,
-1 product or process disclosed, or represents that its_usg
would not infringe privately owned rights. - -

 

 

 

-

- O JuLY 1971

OAK RIDGE NATIONAL LABORATORY
. Qak Ridge, Tennessee
.- - operated by :
_ UNION CARBIDE. CORPORATION
o for the
U.S. ATOMIC ENERGY COMMISSION

DTSTRIBUTION OF THIS DOCUMENT 1S

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u o . CONIENTS

 

v . Abstract . . . . . .. .. e e e e e e e e e,
Introduction . . . & v ¢ v v v e e e e e e e e e e e e e
Experimental PrOCEAUEE « + « + o o o o o o o o o o o o 0 e e
Results and Discussion e e e e e e e e e e e e e e e e e e e
Effect of Preaging on. the MEchanlcal Properties . . .-. e e e e
Effect of Aging on Ductlllty 4 e e e e e s e e e e e
Effect of Aging on the Strehgth .« e e e ; v e e e e e e

B OV N

Stress-Strain Curves . « v ¢« 4 ¢« « s 4 4 s s e s e e s e .. . 10
Comparison of Measures of Ductility . . « v v v ¢ v v . . ... 11
Room-Temperature Tensile ?ropertieS' c f e e e e e e
Creep Properties After 3000 hr of,Aglng Ce e

Effect of Preage Treatments s s e e e e e e e e e e e e

i

Effect of Aging Temperature s 4 s w b 4 e s 4 s 4 e s o e w
Effect Of Tltanlm . .. » - . . s . . - . - . .- . - . . . .

&)

Comparison of Tensile and Creep Results c e e e e e e ae e e
Results ofrUnaged Samples e e e e e ; e e d e e e
Results for Aged Samples . . . . . . . . . C S e e e e e

Phases Identified in Aged Alloys C e et e e e e e e

MEtallography e e e e e e e e e e e e e e e e e e e e .20
Initisl Microstruetures s 4 e s e 4 4 e s s s s e e e s« 20
Mlcrostructures Developed During Aglng C e e e e e e e e 20

STEERGGHGREKRK

 

;
Structures Developed in Prestrained Materlal'.f, . i e . . .23 . 3

Fracture of Samples Tested 1n Creep T o |

Summary and Conclu31ons S~

| | Appendlx C e e e e ':;lfi; ce e e e e e e .~. C e e .. .35

 

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'THERMAL STABILITY OF TITANIUM-MODIFIED
HASTELLOY N AT 650 AND 760°C

C. E. Sessions and E..E. Sté,nsbury1

ABSTRACT

We have investigated the influence of small titanium
~ additions on the thermal stability of Ni-12% Mo-7% Cr-0.07% C.

' The mechanical properties at 650°C (tensile tests at 0.002/min
strain rate and creep tests at 40,000 psi stress) were mea-
‘sured for four heats of this alloy'w1th titanium contents
from 0.15 to 1.2%. Solution annealing temperatures were 1177 -
“or 1260°C, and subsequent precipitation heat treatments were
conducted at 650 and 760°C. Titanium increases the stebility
of a complex MC-type carbide. At low titanium levels the
MC carbide is stable at 650°C but is unstable at 760°C, where
an MpC-type carbide is precipitated, resulting in inferior
properties. For the higher titanium concentrations the
MC carbide is stable on aging at 760°C and results in excel-
lent properties after a solution anneal at 1177°C. However,
high-titanium alloys are significantly less ductile if they
are solution amnealed at 1260°C and aged at either 650 or
760°C. The heat with the lowest carbon content (0.04% C)
was most resistant to property ‘changes on aglng up to
10, 000 hr at both 650 and 760 c.

 

INTRODUCTION

We are concerned'witfifmbdifying the composition of Hastelloy N for

better perfbrmance as a material for a mplten-salt thermal breeder reac-

. tor w1th an expected llfetlme of 30 years or more. The ex1st1ng
| Hastelloy N developed primarlly at Oak Ridge over the past 10 to

12 years functioned well in an 8-Mw(t) reactor built five years ago and

ff0perated over the past three years.2 However, this reactor operated at

 

' 1Consultant,fro:mthe Uhifiersity of Tennessee, Knoxville;VTennessee.

24, E. McCoy, Jr.; An Evaluation of the Molten-Salt Reactor Experi-
ment Hastelloy N Survelllance Specimens — Fourth Group, ORNIL~TM-3063
(March 1971) .

 
 

2

650°C, and-when it was construéted very littlé was known about the
problem of high-temperature irradia.tion da.ma.ge‘ to this alldy. Since
that time we have observed® damsage by thermal neutrons on meny mate-
rials, including nickel-base alloys such as Hastelloy N, in which a
larger deterioration in properties is observed than in most iron-base
alloys. -Thus , an alloy with greater resistance to radiation damage is
needed. )

. We found that small additions of reactive elements, notably tita-
‘nium, can significantly enhance the fiostirradiation'creep-rupture life
and ductility.4 Thus, we have proceeded with the development of &
modified Hastelloy N conteining titanium. This report represents the
phase of the program concerned .with the effect of thermal and mechani-
cal treatment on the creep and tensile behavior of these experiménta.l
a.lloys. ' a

EXPERIMENTAL PROCEDURE

The material investiga.ted: in this program was produced by & com-
mercial vendor. These four 100-1b heats wgi'e vecuum induction melted
with initisl ingot brea.kdown at 1177°C and final fabrication to 0.5-in.
plate at 870°C. The titanium content was the primary variable, being
0.15, 0.27, 0.45, and 1.20% as shown in Table 1. Although & nominal
carbon level of 0.07% was specified, the carbon content of Heat 466-548
wes lower than the others. The various thermal and mechanical treat-
ments epplied are given in Table 2. The variables include solution
annealing conditions, cold work, aging time, and aging temperature.

The dimensions of the mechanical property test specimen'are shown ' |
ianig. 1. This specimen was used so that the results could be t.:ompared.
directl_y to postirradiation mechanical pi-operty tests for which this

 

3H. E. McCoy, "Variation of the Mechanical Properties of Irradiated
Hastelloy N with Strain Rate,” J. Nucl. Mater. 31, 67-85 (1969).

- %H. E. McCoSr,‘ Jr., Influence of Titanium, Zirconium, and Hafnium
Additions on the Resistance of Modified Hastelloy N to Irradiation
Damage at High Temperature — Phase I, ORNL-TM-3064 (Januery 1971).
 

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Table 1. Chemical Analyses of Alloys in Aging Program

 

466=535%

Chemical Analyses, wt %

 

 

466=541 L466-548 467-548

Mo 12.8 13.2 12.4 12.0
Cr 7.2 6.8 7.7 7.1
Ti 1 0.15 . 0.27 0.45 1.20
C 0.073 0.07 0.04 0.08

. Fe 0.03 . 0,03 0.03 0.04
Si 0.07 0.05 0.05 0.03
Mn 0.12 0.10 0.14% 0.12
Mg 0.034 0.01 0.023
W 0.007 0.07 0.007 0.004
Zr 0.0002 0.0007 < 0,0001 0.002
v < 0.01 - < 0.01 < 0.01 0.001
Co 0.02 - 0.04 0.03 0.15
Cu 0.002 0.005 0.0005 0.01
Nb- < 0.0001 - - 0.0001 - 0.0003 0.0005
Al < 0.03 < 0.03 < 0.03 0.05

s 0.003 - < 0,002 0.003 < 0.002
P , o : - - 0.0004
B 0.0002 . 0.0002 0.00007 0.0007
0 0.0012 0.0008 0.001 0.0003
N 0.007 0.0005 0.0009 0.0002
H 0.0002 0.0002 0.0002 0.0003
Ni Balance Balance Balance Balance

aHeat numbers of 100-1b heats obtained from the Stellite'Division
ion. - _

of the Cabot Corporat

Table 2. VariablesAConsidered in This Study

 

- Variable

Titanium content, %

Preage treatment -

 Aging temperature, °C
- Aging time, hr o
 Testing after aging

Tensile
. Creep

Levels or Treatments

0.15, 0.27, 0.45, and 1.2

;_;_Soiution ameal 1 hr at 1177°C
‘Solution anneal 1 hr at 1260°C
Anneal 1 hr at 1177°C, then pres
room temperature '

650 and 760

1500, 3000, and 10,000

- 650°C, 0.002/min strain rate

- ':_'__650'0_(;’_-.'40,'000_ psi stress -

train 10% at

 

 
 

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Fig. 1. Mechanical Property Test Specimen.

specimen was designed. Specimens were solution annealed in.argon end
aged in stainless steel capsules that had been evacuated and backfilled
with argon. Conventional tensile measurements were carried out fising
an Instron tensile machined equipped with a specimen furnace. Two
Chromel-P vs Alumel thermocouples were used with a proportioning con-
troller to keep the test specimen Va.t 650 = 3°C.

The mitrostructures developed during aging were studied by optical
and electron transmission miecroscopy, extraction replication, and
scanning 'microscopy. In addition, precipitated phé,ses present in cer-
tain alloys were identified by x-rey diffraction using the Debye- |
Scherrer technique on precipitates extracted electrolytically with a
methanol-10% HC1l solfi;tion.r Details of the microstructure exsmination
have been répbi'ted,5 ‘and the alloying effects of titanium have been
summarized.® We shall concern ourselves here with the detailed mechani-
cal property response to heat treatment and a limited discussion of
’ miérostructures revealed in optica;l meta.llbgraphy.

 

3C. E. Sessions , Influence of Titanium-on the High-Temperature
Deformation and Fracture of Some Nickel Based Alloys, ORNL-4561
Tduly 1970).

6C. E. Sessions, E. E. Stansbury, R. E. Gehlbach, and
H. E. McCoy, Jr., "Influence of Titanium on the Strengthening of a
Ni-Mo-Cr Alloy," pp. 626—630 in Second International Conference on the
Strength of Metels and Alloys, Conf. Proc., Vol. 1I, The American Society
for Metals, Metals Park, Ohio, 1970.
 

 

 

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Because of the large number of variables 1ncluded in this study
and the attempt to measure ‘the ccmplicated and very subtle differences

introduced by titanium.addltions, these experiments were statistically

designed as a full factorial replication7.including the variables in
Table 2. The purpose of this experimental design was to prcvide better
evaluation and separation’ofrthe'influence of the specific variables on

the tensile property response. Results of these statistical analyses

" have been discussed previously.®

RESULTS AND DISCUSSION

Typical results of the influence of preage thermal-mechanical
treatment, aging time, and temperature‘on the tensile and creep proper-
ties at 650°C will be presented first. The phase identification'will
then be presented and discussed in light of the mechanical property
changes. Strength and ductility values for all specimens tested are
given in’the Appendix;r | ' '

- Effect of Preaging:onrthe Mechanical Properties

Flgure 2 shows the yield strength (0 2% offset) and total elonga-
tion values for the three preage treatments investigated. Solution
annealing 1 hr at 1260° C lowered the strength and increased the ductil-
ity at each titanium level as compared to a l-hr solution anneal at
1177°C. Prestraining 10% at room temperature after allphr 1177°C solu-
tion anneal doubled the hlgh—temperature yield strength and reduced the

T{'ductility by one-third The yield strength increased'with titanium

content for these unaged specimens, but the tensile ductility was not

| appreciably affected by the titanium content for these pretest

treatments.

 

- TE. M. Bartee, Engineering Experimental Des1gn Fundamentals,
Prentice-Hall, Englewood Cliffs, N.J., 1968.

| 8C S. Lever and C. E. SeSS1ons Fuels and Materials Development

 Quart. Progr. Rept. Sept. 30 1969, omso, Pp. 215-279.

 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

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TITANIUM CONCENTRATION (%)

Fig. 2. Effect of Titanium Content and Preaging Treatments on the
Strength and Ductility of Modified Hastelloy N at 6570.°C

Effect of Aging on Ductility

Changes in tensile elongation with aging time are shown in Figs. 3
through 5 for both aging temperatures. Aging at either 650 or 760°C
after a solution anneal at 1177°C enhanced the ductility at 650° °C for
the 1.2% Ti heat and decreased it for the 0.15% Ti heat. After a l-hr
solution anneal at 1260°C, the 1.2% Ti heat lost ductility on aging at
either aging temperature to the same extent as did the 1owest-titanium
‘level, 0.15%. Thus, the larger grain size or greater‘amcunt'of 501ute
in solid solution after the 1260°C treatment maede the high-titanlum
heat susceptlble to a detrimental aging reaction to which it was immune
~ after the 1177°C anneal. ,

The effect of prestraining on the aging behavior is shown in
Fig. 5. Samples were solution annealed 1 hr at 1177°C, prestrained 10%
&t room temperature, aged at either 650 or 760°C, and then tested at
650°C. The verietion in tensile ductility with aging time and titanium
was similar to that shown in Fig. 3 for samples not prestrained; how-

ever, the overall ductility was lower for the prestrained material.
 

 

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 Aging Time on the Tensile Ductility at 650°C Aging Time on the Tensile Ductility at 650°C

‘after a l-hr Solution Anneal at 1177°C. after a l-hr Solution Anneal at 1260°C.

 

 
 

 

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AGING TIME (1000 hr)

Fig. 5. Effect of Titanium Content and Aging Time on the Tensile
Ductility at 650°C for Samples Solution Annealed 1 hr at 1177°C and Pre-
strained 10% at Room Temperature.

For aging at 650°C the intermediate (0.45%) and high (1.2%) titanium
heats showed an incresse in ductility with aging time. The low (0.15%)
titanium heat exhibited & rapid loss in hot duetility with aging tifie at
650°C. Aging at 760°C affected the ductility of the low- and high-
tité.nium heats the same as aging at 650°C , whereas the intermediate

(o. 45%) ‘titanium heat showed a ductility deterioration with time at
760°C in contrast to the ductility enhancement found after a.g:.ng at
650°C.

Eff'ect of Aging on the Strength

‘The changes in the yield strength on aging are shown for only the
- 1=hr solution anmneal at 1260°C in Fig. 6. For aging at both 650 and
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AGING TIME (hr)

Fig. 6. Yield Strength Changes as a Function of Aglng Time and
Titanium Content for a l-hr Solution Anneal at 1260°C.

760°C the largest strength increase occurred for the 1.2% Ti heat.
The intermediate (0.45%) titanium heat exhibited the smallest change in
yield strength for aging at either temperature. The lower'yield

‘strength values in this 0 45% Ti heat were most likely due to the lower
carbon concentratlon for this heat (0.04% as compared to O. 07%) .

The changes in yleld strength for samples solutlon annealed at

'1177 ¢ before aging (not plotted) were comparsble in most cases to
*those shown in Flg.;6.- For the l 2% Ti heat solution annealed at 1177°C

and aged at 650°C, the yield strength at 650°C peaked at 46,000 psi
after 1500 hr aglng and exhiblted classical overaging by decreasing to

:.hi37 500 psi after lO 000 hr The increase in strength on aging the
1. 2% Ti heat at 760°C after the 1177 C solution anneal was appreciably

lower than for the 1260°C anneal ' The yield strength peaked:at
41,000 psi after 1500 hr and decreased to 37,000 psi after 10,000 hr.

 
 

10

The strength of the prestrained alloys did not increase on aging.
The 1.2% Ti heat actually showed a yield strength decrease with aging
time at both 650 and 760°C, and the inf.ermedia.te- and low-titanium
heats did not show any large cha;nge,in yield strength after aging.

Stress-Strain Curves

The stréss-stra,in curves for the two ,heé.ts thet showed the lowest
and highest strengths are shown in Figs. 7 and 8. The engineering
. stress and strain are plotted up to the ultimate stress for the 0.45
and 1.2% Ti heats in Figs. 7 and 8, respectively. TFor the 0.45% Ti heat

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STRAIN (%) :

Fig. 7. Stress-Strain Curves for the 0..45% Ti Heat at 650°C for
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Fig. 8. ©Stress- Strain Curves for the 1. 2% Ti Heat at 650°C for
Various Heat Treatments.r'_

~ultimate tensile stresses from 50, 000 to 90 000 psi and uniform elonga-
" tions from approxxmately 15 to 40%‘were achieved by the thermal-

mechanical treatments 1nvestigated The 1.2% Ti heat (Fig. 8) exhibited

'a range of ultimate tensile strengths from.75 000 to 105 000 psi and -

uniform elongatlons from 8 to 40%.

' ComperiSonrof'MEasures of Ductility
Dependlng on the appllcatlon, our crlterla of usable ductlllty
often vary. Thus, Flg. 9 compares three measures of the tensile ductil-

ity at 650°C for a l=hr anneal at 1177°C and aging at 650°C _Thls
partlcular treatment produced a high total elongation for the 0f45 and

 
 

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TOTAL ELONGATION UNIFORM ELONGATION '
50 |-
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0O 2 4 6 8 100 2 4 6 8 100 2 4 6 8 10
: AGING TIME (4000 hr)

Fig. 9. Comparison of Measures of Tensile Ductility for a l-hr
Solution Anneal at 1177°C and Aging at 650°C.

1.2% Ti levels, as previously discussed, but the 1.2% Ti heat actually
deteriorated in uniform elongetion from 35 to approximately 26% with
aging time. Therefore, to account for the high total elongation
observed, the nonuniform elongation must have increased to compensate
for the observed decrease in uniform stré.ih. That is to say, for this
heat, which exhibited the maximum ductility after this particular treat-
ment, we found & change in the stress-strain relationship induced by
precipitation. The change corresponded to an increase in yield strength,
an increase in reduction in ares., and an increase in total eloriga.tion;
however, it corresponded to a deérease in ultimate tensile strength and

uniform elongation.

Room-Temperature Tensile Properties

Thé effect of aging treatment on the room-temperature tensile
properties' is given in Table 3. The yield strengths and ductilities
Table 3. RoompTemperaiure Tensile Properties at 0.002/min Strain Rate for Titanium-
e ‘Modified Hastelloy N After Various Heat Treatments

 

Heafi Treatment ‘
Reduction

 

 

 

- Solution Aging Aging Strength, psl Elongatioh, % in Area  Specimen
Alloy Anneal = Time Temperature Yield Ultimate Uniform Total , %) P
| (°cy  (ar)  (°0) Tensile | | o
I el x 10°  x 102 | o | R
L 466-535 1177 .. O 422 . 119.5  6L.0  64.2 66.4 5292
%177 15000 650 615 - 135.9 40.5 - 40.8 29.5 5997
| Coonav7 o 1500 760 53,3  139.1  43.5 46.0 4.8 5998
466-541 - 1177+ O 38.2 112.7 69.6  173.3 57,1 6093
177 1500 650 54.3 125.6 52.5 537 41.9 6099
1177 1500 760 . 424 110.5 50,1 52.3 40.6 6100
- 466-548 1177 o 137.4 13.3  68.8 72.2 62.6 6270
N | 1177 1500 650 47.6 124.1 57.4  59.0 41.8 6277
| 177 1500 760 48.1 129.6 53.6 55.4 50.9 6278
467-548 . 1260 0 | 35.7 104,9 67.0 ©  68.2 - 18.2 6344
| 177 0 | 47.6 - 134.9 62.5 - 63.8 46.2 11416
1177 1500 650 61.3 145.6 50.5  52.0 41.2 - 6321
1177 1500 760 50.1 121.3 47.8 49.0 43.6 6322
1226 41.6 417 36.7 6989

1260 . 1500 760 59.6

 
 

14

are similer for the various heets as solution annealed at 1177°C.

Aging 1500 hr at 650°C increases yield strength more than at 760°C.
However, the'ductility decréase.on eging was approximately the same for
the two sging temperatures. These trends in strefigth and ductility
were found for heats 466-535, 466-541, and 466-548. For heat 467-548
(1.2% Ti) we measured the effect of using a higher solution anneal tem-
perature (1260°C) on the agihg response. Aging 1500 hr at 760°C
increased the room-temperature yield strength from 36,000 to 60,000 psi
and decreased the creep elongation from 68 to 42% strain. This change
in the room-temperature tensile behavior for the 1260°C solution anneal
was larger than that found after annesling at 1177°C. However, the
microstructural differences found for this high-titénium (1.2% Ti,
467-548) heat are significant and have been discussed previously.® The
point is that, regardless of the differences in microstructfire, the

room-temperature properties are similer.

Creep Properties After 3000 hr of Aging

The effect of titanium content on the creep properties of solution-
annealed and aged alloys was measured for 36 specimens. Each test was
conducted at 650°C and under 40,000 psi stress so that the rupture lives
and creep rates indicate the effects of aging. Creep ductility in these
tests reflected the influence of strain rate, which varied with the
strength of the alloy under this constent-load creep test. The ductili-
ties, rupture lives, and creep rates are given in Figs. 10 and 11 for

some of the treatments used.

Effect of Preage Treatments

For each alloy, except the 1.2% Ti heat, the lowest rupture life
and highest elongation corresponded to samples solution annesled at-the
higher temperature (i.e., 1 hr at 1260°C). In most instances the lowest
secondary creep rate, maximm rupture life, and minimum ductility corre-
sponded to the samples prestrained 10% at room temperature. |

 

°c. E. Sessions, Influence of Titanium on the High-Temperature
Deformation and Fracture of Some Nickel Based Alloys, ORNL-4561

(July 1970).
»

n

ORNL=DWG 69-4760R

n

     

0° 0!
5 T 5
- 2
£ W
& 2 & 2
- .
i u
S o2 O 2
E® _ S©
2 a ASSOLUTION ANNEALED| 2 -
' CHTTC 43
S e H77°C + 3000 hr AT = 5
650°C
H77°C + 3000 hr AT
2 1260°C + 3000 hr AT
760°C
o' o 03 : _
0 03 06 - 09 - 1.2 0 03 .06 09 1.2
TITANIUM (wt %) - : : : TITANIUM (wt %)

Fig. 10. Variation.in Rupture Life and Creep Rate with Titanium

Content and Heat Treatment at 650°C and 40,000 psi Stress.

Effect of Aging Temperature

In most cases (10 of 12 tests) the rupture life after aging 3000 hr
at 760°C was reduced by from 20 to 90% as compared with the unaged sam-

ples. Correspondingly, the'creep rates were significantly increased by a

 factor of 3 to_lQ,‘depending_on the alloy and heat treatment. The creep

ductility of the samples aged 3000 hr at 760°C ranged from 6't9 37%,

with the higher values corresponding to the higher titanium level (1.2%).
Less pz_‘opel.'ty.déteri-c.ji'.atio_.fi was dbserved for aging 3000 hr ,a.{: 650°C.

In fact, several heats, ndtabLy the 0.45% Ti heat, exhibited significantly

‘enhanced rupture.lives andnfracture strains even though the creep rate
. was increased apprecisbly by the aging treatment. In general, the pre-
"-'strginéd samples and-thqse'With the larger grain size (i.e., solution

annealed at 1260°C) showed the largest reduction in rupture life after

aging at 760°C.

. Effect of Titanium

 Generally, the creep-rupture lives and ductilities increased with
increasing titanium content for both the solution-annealed and the aged

samples. The creep rates_décréaséd with increasing titanium content.

 
 

 

16

ORNL-OWG 69-4113

7

55

 

 

-~

‘45 : -~

) I %

 

 

 

 

 

 

A
LA L T
A

 

 

 

 

[ 1LY

5 -
{ s SOLUTION ANNEAL W77°C
® 1177°C + 3000 I 650°C
0 = a H177°C + 3000 br 760°C —

o 1260°C + 3000 hr 760°C

 

 

 

 

 

 

 

 

 

 

 

 

 

0
o i35 03¢ 043 060 075 090 105 {2

TITANIUR (%]

'Flg 11. Effect of Titanium Content and Heat Treatment on the '
40 OOO—psi Creep Elongation at 650°C.

The titanium content influenced the magnitude of the aging effect.
For most treatments, the alloy with the highest titanium content
exhibited the smallest property deterioration after aging. However,
‘solution annealing for 1 hr at 1260°C led to creep property changes on
aging that were quite significant for all levels of titanium.

Comparison of Tensile and Creep Results

Table 4 lists the heat treatments for each heat of material that
corresponded to the maximum and minimum property value found in creep
and tensile testing of the samples aged 3000 hr. Since three treatments
‘were included in this study-[i.e,, solution anneals of (1) 1 hr at 1177°C
(2) 1260°C, and (3) 1 hr at 1177°C plus 10% prestrain at room temperature]
the treatment that is not listed under either minimum or meximum in
Table 4 would result in properties intermediate between the minimum and

maximum values.
 

 

O

.

N

17

| ‘Table 4. Thermal-Mechanical Treatments That Produced the Maximum and

Nhnimum.Creep and Tensile Properties at 650°C for Various Heats of
Tlt&nlumrMbdlfled Hastelloy N

 

- Condition

. Heat Treatment® to Give Extreme Property

 

 

 

| Propéfty o for Each Titanium Content in Percent
R 0.15 0.27 0.45 - 1.2
4 Tensile Resultsb
Unaged  Maximum yield strength  cw oW oW oW
: Minimum yield strength 1260 1260 1260 1260
Maximum ultimate strength cw cw ow cw
Minimum ultimate strength 1260 1260 1260 1260
Maximum total elongation 1260 1260 1260 1260
- Minimum total elongation cw oW cw cw
Aged Maximum yield strength  cw/650 cw/650 cw/650  cw/650
| Minimum yield strength 11260/760 1177/760 1260/760 1177/760
Maximum ultimate strength 1177/760 cw/650  cw/650  cw/650
Minimum ultimate strength 1260/650 1260/650 ~1260/760 1260/760
Maximum total elongation 1177/760 1177/650 1177/650 1177/650,
p 760
Minimnm.total elongatlon cw/650 1260/650 cw/760  1260/760
o | Cre EE Results® ,
Unaged Maximum rupture life cwW 1Y77 - cw 1260
' Minimum rupture life 1260 1260 1260 1177
Maximum secondary creep 1177 1260 1177 1177
rate R |
- Minimum secondary creep  cw cWw . Cw - 1260
- rate e | ' ' | .
. Maximum creep elongation 1260 1260 1260 1177
Minimum creep elongatlon oW ew . ew 1260
Aged . Meximum rupture life . 1260/650 cw/650  cw/650 cw/650
— . Minimum rupture J_.:l_.:t‘e__ 1260/760 1260/760 cw/760  1177/760
Maximum secondary creep = 1260/760 1260/760 1260/650 1177/760
rate Sl o S , :
Minimum secondary creep  1177/760 cw/650  cw/650  cw/650
~ rate e - _ o -
‘Maximm creep elongation cw/650  1177/650 1177/650 1177/650
Minimum creep elongation ew/760 cw/760 . ew/760  1260/760

 

®The preage heat treatments are 1177 1 hr

1260°C; cw, 1 hr at 1177°C, then cold'worked.

followed by the aging temperature_ln °C.

at 1177°C; 1260, 1 hr at
For aged specimens this is

Aging times were 3000 hr.

bTensil.e properties measured at 650°C and 0.002/min strain rate.

Ccreep properties measured at 650°C and 40,000-psi stress.

 
 

Results of Uneged Samples | : | - -

 

In each case the maximum yield strength and maximum ultimate ten-
sile strength at 650°C were found for the 10% prestrained samples. The
minimum yield and ultimate strengths were found for the 1260°C treatment.
The minimum and maximum tensile ductilities corresponded to the strongest
and weakest alloys in tensile tests, with one exception, the 0.45% Ti
heat.

In creep the maximum resistance to rupture (meximum rupture life)
corresponded to the prestrained sample in only two of the fbur heats,
but prestrained samples crept at & lower rate in three of the four cases.
Notably, the 1.2% Ti heat exhibited the minimum and maximum creep resis-
tance (as measured by the creep rate) for the 1177 and 1260°C anneals,
respectively. This indicates that creép of this higher titanium heat
was affected more by solution annealing treatment than by 10% cold
.working The minimum and maximum creep ductzllties of the other alloys
corresponded to the prestrained and to the 1260°C annealed material,

respectively.

Results for Aged Samples

The prestrained material aged 3000 hr at 650°C exhibited the
maximum yield strength for each heat. The minimfim yieid stréngth corre-
spdnded to either 1177 or 1260°C samples aged 3000 hr at 760°C; depending
on the alloy. The maximum ultimate strength was for prestrained samples
aged 3000 hr at 650°C with one exception, the 0.15% Ti heat. The lowest
ultimate strength (UTS) was found for the 1260°C treatment; at lower
tltanium level an aging treatment.of 3000 hr at 650°C caused the lowest
UTS, and at the higher titanium levels aging 3000 hr at 760°C caused the
lowest UTS. The rupture life was & maximum for prestrained samples
aged at 650°C, snd the treatment that produced the minimum rupture 1ife
differed for each alloy. An intereéting point is that the 0.15% Ti
heat gave both the maximum afid.the minimum rupture life for'the'larger
grain size 1260°C treatment, depending on whether it was aged at 650 or
760°C.
n

"

19

Phases Idéentified in Aged Alloys

~ Table 5 lists the phases identified by Debye-Scherrer x-ray dif-
fraction analysis of electrolytically extracted precipitate particles
after 3000 hr of aging at the indicated temperatures. In the 0.15% Ti

‘heat a combination of MC and M.C type phases was present, but at 760°C

only the MpC carbide was found. At the intermediate titanium level

 (0.45%) only MC was present at 650°C, while at 760°C a weak trace of MC

and strong lines for M>C were found. At the highest_titanium level
investigated (1.20%) the MC structure was present after aging at either
temperature. Thus, we conclude that the higher titanium content has

stabilized the MC-type carbide at higher aging temperatures.

Table 5. FPhases Identified in Ni-Mo-Cr-Ti Alloys
by Debye-Scherrer Analysis

 

 

. , o Titenium = ‘ppages
Heat Heat Treatment , , Content Present
f - (%)
| 466-535 1 hr 1177°C + 3000 hr at 650°C 0.15  MC® + MyC
— 1 hr 1177°C + 3000 hr at 760°C 0.15 = M0C
466-548 As annealed 1177°C o | b
' As amnealed 1260°C b
1 hr 1177°C + 3000 hr at 650°C . 0.45 MC_
1 hr 1177°C + 3000 hr at 760°C 0.45 MC® + MyC
467-548 As annealed 1177°C 1.20 MC
“ _As annealed 1260°C - - 1.20 ' MC
1 hr 1177°C + 3000 hr at 650°C - 1.20 - MC
1 hr 1177°C + 3000 hr at 760°C 1.20 MC
1 hr 1260°C + 3000 hr at 650°C © 1,20 MC
1.20 Mo

1 hr 1260°C + 3000 hr at 760°C

 

gMihOr amDUnts'prESént;~'-':n
bToo little précipitateftc_extract electrolytically.

The carbides found in these alloyS'were complex ccmpounds involving
chromium, molybdenum, and titanium, ‘and in several cases we found two
different lattice par&meters that generally fit the hexagonal M,C- type

structure with the same c¢/a ratio of 1.63. In the case of the MC

 
 

 

- 20

structure the -lattice constant did not correspond to that of TiC |
(4.33 R) but rather was significantly lower, which indicates that the
smaller molybdenum or chromium atoms were subStitutihg for titanium in

the face-centered cubic MC structure.

| Metallography

Initial Microstructures

 

The effect of solution annealing at 1177 or 1260°C on the grain

- size for three of the four alloys is shown in Fig. 12, Annealing 1 hr
at 1177°C did not produce & single-phase alloy for most of these heats.
In the alloys containing 0.15 and 1.2% Ti after an 1177°C anneal, large
stringers of second-phase particles were evident. The 0.27 and 0.45% Ti
heats were virtually single phase after the 1177°C heat treatment.

After 1 hr at 1260°C the grain size of each heat was 2 to 4 times that
found after the 1177°C anneal. Also, fewer preéifiitate particles and
stringers were present after 1 hr at 1260°C.

Microstructures Developed During Aging

Typical structures obtained.by interference contrast microséopy
are shown in Fig. 13. This figure compares the precipitate distributions
for samples solution annealed at 1177°C and aged 10,000 hr at either
650 or 760°C. These photographs were taken near the fracture sites of

tensile samples tested at 650°C after sging at the indicated temperatures.

After aging at 650°C the precipitate was generally finer, and the con-
centration of precipitate near grain boundaries appeared to be fairiy
heavy, particulafly for the low-titanium alloj. After aging 10,000 hr
at 760°C the precipitate was rather coarse, and the concentration of
precipitate near the grain bouhdary'was lower than fhat found at 650°C.
The precipitate within the grain boundary was also coarée, &s shown by
the micrographs of the 0.27 and 0.45% Ti heats in Fig. 13. Little pre-
cipitation occurred in.the'l.z% Ti heat, since the stringers evident in
Fig; 13 were actually present to some extent even before aging; this was
shown for the as-solution-annealed condition in Fig. 12.

C
 

. | | - o o Y-93309
Heat 466535 (0. 15% Ti) Heat 466548 (0.45% T)) Heat 467548 (1,2% T1)

    
       

    

 

 

 

1w at
_II77°C

 

 

 

 

 

 

 

 

 

 

 

 

Vl hr at
‘ 1_260°C‘

 

 

 

 

 

 

 

 

Yy-93227

 

 

 

 

Fig. 12. Microstructures of Titanium-Modified Alloys in the Solution-Annealed Condition.
100X. Reduced 38%. )

 

1e

 
 

22

 

 

 

 

 

 

 

 

 

 

 

 

 

0.15%Ti 0. 7% T . 0.45% Ti ' - L2%T

Fig. 13. Microstructures Developed in Ni-12% Mo—7% Cr—0.07% C
Alloys Modified with Titanium, Solution Annealed 1 hr at 1177°C, and
Aged 10,000 hr. 1000X. Reduced 65%.

Microstructures developed during aging 10,000 hr after a l-hr
solution anneal at 1260°C are shown in Fig. 14. These structures are
typical also of those developed in 1500 and 3000 hr of aging and indi-
cate a large variation in optical microstructure with titanium content.
Alloys with lower titanium contents exhibited relatively coarse precipi-
tate distributions. The 1.2% Ti alloy showed bands of precipitate
stringers after aging at 650‘C. Apparently, the stringers present
initially along the hot-working direction of the rod stock were not com-
pletely eliminated even though the material was solution annealed at
1260°C. During subsequent aging these residusl second-phese particles
coarsened. Thus, the 1bngitudinal cross section of these aged alloys
exhibited a banded structure, which is similar to that found in the cold-
worked rod before solution annealing. In the 0.45 and 1.2%'Ti heats
during aging at 650°C, however, some additional precipitation occurred
on both subgrain and grain boundaries. ' |

The microstructures developed at 760°C after e 1260°C solution
anneal, shown in Fig. 14, were actually quite similar to those found
after the 1177°C anneal and 760°C age (Fig. 13). For the 1.2% Ti heat,
 

”

23

 

 

 

 

 

 

 

 

 

 

 

 

 

 

G

 

 

 

 

0.15% Ti T I027%Ti 0.45%T1 1.2% Ti

Fig. 14. Microstructures Developed in Ni-12% Mo-7% Cr—0.07% C
Alloys Modified with Titanium, Solution.Annealed 1 hr at 1260°C, and

-Aged 10,000 hr. 1000X. Reduced 65%

r

however, the structure developed after a 1260°C anneal was much differ-

ent_from that developed in samples annealed at 1177°C. ‘The precipitation

- appeared to be a Widmanstfitten'distribution; however, the concentration

of preeipitate'plateletsishcwnlin Fig. 14 varied within a given grain.

Generally, a higher concentration of "platelets" was found near grain

boundaries and near clusters of primary precipitate particles. This

"platelet";type'precfpitate, which forms primarily after the 1260°C

solution anneal and 760°C age, hes been shown10 by transmission electron

'microscopy to result from.preclpltatlon on stacking faults.

'_Structures Developed in Prestralned Material

 

‘The mlcrostructures developed during aging 10 000 hr at 650 and

- 760°C after annealing 1 hr at 1177 C and prestraining 10% at room tem-

perature are shown in Fig. 15 for the 0.15 and 1.2% Ti heats. Prior

 prestraining greatly reduced the coarseness of the preclpltate in both

 

100, E. Sessions, Influence of Titanium on the High-Temperature
Deformation and Fracture of Some Nickel Based Alloys, 0RNL~4561
(July 1970).

 

 
 

 

    

Y-98357

    

®)

 

e e RLTINY-08354

 

 

 

 

 

 

 

 

 

 

 

Fig. 15. Microstructures Developed in Ni-12% Mo-7% Cr-0.07% C
Alloys Modified with Titanium, Solution Annealed 1 hr at 1177°C, Pre-
strained 10% at Room Temperature, and Then Aged 10,000 hr. 1000X.

(a) 0.15% Ti, aged at 650°C, (b) 1.2% Ti, aged at 650°C, (c¢) 0.15% Ti,
aged at 760°C, and (d) 1.2% Ti, aged at 760°C. Reduced 18%.

alloys at 650°C. In samplés aged at 760°C, presfiraining 10% produced
very little effect on the structure of the 0.15% Ti heat, but the 1.2% Ti
heat showed much more precipitate in the prestrained samples (Fig. 15)
than in samples aged without prestraining (Fig. 13). In the prestrained
1.2% Ti alloy, both a fine random matrix precipitate and a crystallo-
graphic platelet-type precipitate are shown in Fig. 15. Evidently pre-
straining has efihanced the precipitation of this pl&telet-type precipitate
|

»

L

C

25

(the platelets are actually a precipitete on stacking faults,?! as
mentioned prev1ously) This apparent enhancement in formation of
stacking fault precipitate produced by straining this 1.2% Ti alloy

" pefore aging is consmstent’with the effect of prior straining on forma-
~tion of stacking fault prec1p1tates in niobiumpdoped stainless steels

as reported by Silcock and ‘I‘unstall11 and Naybour.l? However, some
differences in the effect of strain on the tendency to precipltate on
stacking faults in steels and in this alloy are discussed in another

paper. ??
Fracture of Samples Tested in Creep

The fracture appearance of creep samples tested after solution

‘annealing at 1177°C and aging 3000 hr at 760°C is compared in Fig. 16.

The rupture lives (64, 135, and 1335 hr) and the total_elongation values
(9. 5 16.7, and 36.8%) of samples with these fractures increased progres-
smvely'with 1ncreasing titanium content. The - fracture appearance con-
firms the measured ductility trends, since the fracture mode exhibits a
trans1tion'w1th increas1ng titanium content from 1ntergranular fracture
for the O. 15% Ti heat [Fig. lfi(a)] to transgranular for the 1.2% Ti

heat [Fig. 16(c)]. The intermediate titanium content (O. 45% Ti) also

: _failed in a predOminantly transgranular fracture mode; however, the
“tips of the fractured grains [Fig. 16(Db)] show regions that appear to

have recrystallized during the creep test. Nevertheless, the creep
elongation of this O. 45% Ti heat after the 760° C heat treatment was only
16.7%, which indicates that the apparent recrystallization that occurred

 during creep testing at 650°C did not signlficantly enhance the ductility.
That is, other heat treatments of this 0.45% Ti heat yielded post-age
creep’ elongation values appreciably greater than 16. 7% even though they

, showed no evidence of recrystallization

 

11J M. Silcock and w J. ‘.T.'unstall Phil. @ 10, 360-—389 (1964)

12p 7, Naybour, Acta Met. 13, 1197-1207 (1965)
13¢. E. Sessions and R.,E.-Gehlbach, “Effects of Heat Treatment and

- Straining on Formation of Stacking Fault Precipitates in Hastelloy N,"

(in preparation)

 
 

 

 

 

 

Y-98588

Y-98591

 

Fig. 16.

Creep Fractures at 650°C and 40,000 psi for Heats

of
Titenium-Modified Hastelloy N After a 1 hr at 1177°C Solution Anneal and

3000 hr Age &t 760°C.

Reduced 18%.

100x. (&) 0.15% Ti, (b) 0.45% Ti, and (c) 1.2% Ti.

 
)

o7

In Fig. 16 it is- also apparent that the grain size is smallest for

“the 1.2% Ti heat. Although the fracture was primarily transgranular in

that heat, many intergranular cracks have opened up along the gage
lenéth'in_contrast to the alloys with lower titanium concentrations.

Apparently, the grain bonndary structure in the higher titanium heat

inhibits the propagation of”intergranular cracks even though the nuclea-

tion of cracks in this alloy is apperently quite easy.

Another comparison of fractures for creep tests after various aging
treatments is given in Fig._l? for three heats with different titanium
contents. The total creep elongations for the Sampies shown in
Fig. 17(a), (b), and (c) were 8.3, 42.1, and 24.5%, respectively. Thus,

the alloy with the intermediate titanium content [0.45% Ti, Fig. 17(b)]

was the most ductile of-the-three alloys, in contrast to the results in

Fig. 16. The grain size and the amount of precipitation evident at this

" magnification is greater in Fig. 17(a) and (c) than in Fig. 17(b). As

in Fig. 16, the incidence of intergranular cracks is higher in the most
ductile alloy, Fig. 17(b), than in the otheér materials. The fracture
mode is a mixture of transgranular and 1ntergranular fracture in
Fig. 17(a) and (c) and is ‘transgranular in Fig. 17(b) |

An influence of agingrtrme;at 760°C on the post—age fracture
appearance for the,0.45%*Ti heat is found by comparing Fig. 17(v)
(1500 nr at 760°C) with Fig. 16(b) (3000 hr at 760°C). These two sam-
ples from the same heat were both solution annealed at 1177°C; aged,

;_andithen.creep tested at 40,000 psi and 650°C. The creep properties
(5% hr rupture life and 42.1% elongation) for the sample aged 1500 hr

at 760°C [Fig. 17(b)] were-mubn:better-than-the properties (135 hr

~ rupture life and 16.7% elongation) obtained after aging 3000 hr at 760°C.

Similari&,-the fracture'appéaranoe'is more transgranular and the amount

of grain deformation is much greater for the samples aged only 1500 hr
[Frig. 17(v)]. .Although a s1gn1f1cant difference was found in creep
, properties after aging 1500 and 3000 hr at 760°C, the optical microscopy

did not reveal any 1arge differences in structure. This effect of aging

. time between 1500 and 3000 hr at 760°C is rather surprising since normally

aging follows a logarithmic rate, which would give small changes for &
factor of 2 difference 1n aging time, but it should change 81gn1ficantly

 

 
 

 

 

 

 

 

 

e W
A
ST =

.y -

 

| Fig. 17. Creep Fractures for Tests at 650°C and 40,000 psi in

; Titanium-Modified Hastelloy N for Various Heat Treatments. 100X. |
(a) 0.15% Ti solution annealed 1 hr at 1260°C and aged 3000 hr at 760°C,

, (b) 0.45% Ti solution annealed 1 hr at 1177°C and aged 1500 hr at 760°C,

L and (c) 1.2% Ti solution annealed 1 hr at 1260°C and aged 3000 hr at 650°C.

: Reduced 18%. '

 

 
-

29

for a tenfold difference in aging time. NevertHeless, this observed

large difference in ereefi'behavior between 1500 and 3000 hr of aging

 was confirmed by the tensile ductility for this 0.45% Ti heat, as

shown in Fig. 3, p. 7. The 650°C tensile ductility decreased from about
38 to 21% between 1500 and 3000 hr of aging at 760°C. Thus in the

0.45% Ti heat, ev1dently either the type of precipitate present or its

amount and distribution differed significantly between 1500 and 3000 hr

- of exposure at a temperature of 760°C (phases were not identified in the

samples aged 1500 hr).
SUMMARY AND CONCLUSIONS

We have investigated the effects of titanium content, solution
annealing temperature, aging temperature, and aging time on the mechani-
cal properties of modified Hastelloy N at 650°C. In the solution-

- annealed condition the 0.002/min yield strength increased and the

ductility was approximately constant with increasing titanium content

from 0.15 to 1.2%. Aging“andtprestraining both generally increased the

~yield strength, but the changes in ductility during aging depended on

(1) alloy content, (2) amnealing temperature, and (3) aging temperature.

~ Phase identiflcation of electrolytically extraced precipitates indi-
eated two types of carbides present, An M>C-type was favored in heats
with low titanium contents_or elloyS'eged at the higher'temperature,
760°C; An MC-type carbide was favored at higher titanium contents or

'fbr'alloys'with'a given titanium coneentration aged at'the'lewer temper-

ature, 650°C. Thus, the titanlum.content determined the carbide type,

-: whether MaC or MC. The carbide type, its stability, and its morphology

'.'eubsequently influenced the deformation and fracture in creep and tensile

teets. Correlations of microstructures and high-temperature mechanical

| fiproperties as influenced by the carbon and titanium contents have pre-

1 :v1ously been investigated 14

 

14¢, E. Sessions and E. E;rStansbury, "Correlation of Structures

and High—Temperature Properties in & Ti-Modified Ni-Mo-Cr Alloy,"

submitted to Metallurgical Transactlons.

 
 

30

Increasing fihe titanium content resulted in improved post-age
properties. Apparently, the M:C carbide precipitate was mbre detrimen-
tal to ductility than was the MC carbide precipitate. However, certain
~distributions of MC carbides were also detrimental to the ductility.
This was indicated by the 1.2% Ti alloy, which lost ductility and
increased its strength on aging after.the 1260°C solution anneal. Since
~ only MC carbides were idehtified in this alloy, the strengthening and
 embrittlement are attributed to the precipitation on stecking faults.
Correlations of microstructures and mechanical properties for stacking
fault precipitates have also been made for stainless steels,?5,16
Inconel,l”? and Hastelloy N.18,19 |
| Several additional points should be made from these results.
The alloy with the lowest carbon content (i.e., heat 466-548) gener-
“ally had the lowest strength and highest ductility. This low carbon
alloy,rhcfievef, did shcfi a pronounced difference in aging response at
650 and 760°C, which we attribute to the fact that different carbides
were precipitated at the two aging temperatures (i.e., MC at 650°C and
MoC at 760°C). D

Overaging could be expected in most of these alloys before 3000 hr
at 760°C, as indicated from the peak in yield strength at 1500 hr in
Fig. 6, p 9.

The comparison of creep and tensile properties indicated that the

higher solution anneal, which increases the grain size and amount of

 

153, M. Silcock and W. J. Tunstell, Phil. Mag. 10, 360-389 (1964).
16R. D. Naybour, Acta Met. 13, 1197-1207 (1965).
17p. 8. Kotval, Trans. Met. Soc. AIME 242, 1651-1656 (1968).

18c. E. Sessions, Influence of Titanium on the High-Temperature
Deformation and Fracture of Some Nickel Based Alloys, ORNI~4561

(July 1970).

19¢. E.. Sessions, E. E. Stansbury, R. E. Gehlbach, and
H. E. McCoy, Jr., "Influence of Titanium on the Strengthenlng of &
Ni-Mo-Cr Alloy," pp. 626-630 in Second International Conference on the
Strength of Metals and Alloys, Conf. Proc., Vol. II, The Amerlcan 8001ety
for Metals, Metals Park, Ohio, 1970.
 

‘

31

_solutelin solution, actually produces the lowest yield strength. The

weakest alloys in creep (as measured by the highest secondary creep rate)
were found for the 1177°C anneal, which produces a smaller grain size

relative to the structure prdduced at 1260°C. This dependence of low

- yield stfength on 1arge grain size and of low creep resistance on small

grain size is typical of the effects of grain size on the strength of

alloys at elevated test temperatures. At such temperatures, inter-

granular fracture rather than transgranular fracture is favored.

Of the treatments 1nvestigated cold working 10% at room tempera-

- ture after a l-hr solution anneal at 1177°C had the greatest effect on
raising the strength and lowering the ductility, both before and after

aging. An exception to this was found for the 1.2% Ti heat, which
exhibited the lowest post-age tensile and creep ductility after a 1260°C
solution anneal and 760°C ege.' Semples that showed this large ductility
loss had a high concentration of the platelet-type stacking fault
precipitate. ' '

The heat treatments 1nvestigated produced variations in the high-
temperature ductility between 5 and 50%, yet the significant p01nt is

that most of these treatments change the room-temperature properties
only slightly. The irony, however, is that these subtle microstructural

modifications caused bj titanium additions have affected the cfeep
behavior of this material even more than we found fram this systematic
look at tensile behav1or. However, testlng in creep does not give a
constant rate of defqrmation,'and_the strain rate sensitivity of the
ductility would have made it more difficult to separate effects'of‘aging
time and'temperature than it°has been for this tensile evaluation How-
ever, the creep properties follow similar trendS‘W1th precipitate type

" and distribution and are, in fact most important in reactor applica-
~ tions of this alloy. | |

 
 

"
 

' APPENDIX
 

-
 

%

 

 

.L“"“

35

Table 6. Heat Treatment Designations .

 

 

Solutlon-Annegl - - Age Treatment

 

Designation Temperatur  Prestrain® Time Temperature

| (cc) | (hr) (°c)
109 1177 _ yes 0 '
121 1177 no 0
123 1260 ' no 0 - _
130 1177 ' no 1,500 650
131 - 1 N . no 1,500 760
132 1177 no 3,000 650
133 1177 , no N 3,000 760
134 | 77 no 10,000 650
135 ' 1177 - no - 10,000 760
136 . 1260 no 1,500 650
137 1260 no 1,500 760
138 ' 1260 : no 3,000 650
139 1260 - no - 3,000 760
140 1260 ' no 10,000 650
141 1260 - . no 10,000 760
142 nunw yes 1,500 650
143 1177 . yes 1,500 760

- Y4 o . yes ' 3,000 650

- 145 | nr7 - ~ yes 3,000 760
146 | a7 yes 10,000 650

147 oouam yes 10,000 760

 

®For 1 hr. | | |
. bStrained 10% in tension a.t room temperature after the l-hr anneal
at 1177°C. o : ' o

 
 

 

36

Table 7. Tensile Data for Heat 466-535 (0.15% Ti)
for Various Heat Treatments

 

 

 

Test
. Strength, psi ' Reduction
Heat Temper- - Elongation, % .
?Pe°im¢n_ Treatment ature Yield Ultl?ate Uniform Total in Ares,
, (cc) Tensile (%)
| x 10® x 10° o | |
5989 121 650 29.3 71.4  24.0 24.6 27.28
5991 121 650 28.6 76.1 30.4 30.9 21.18
6019 123 650 26.1 66.7 29.0 - 30.0 22.61
6017 123 650 24.5 64.5 31.8 33.0 32.09
5993 130 650 38.7 . 77.5 24.9 27.7 26.61
6040 130 650 40.7 81.9 27.0 32.5 - 23.40
5994 131 650 35.2 73.9 15,9 16.4 17.59
6041 131 . 650 - 35.2 76.2 16.6 17.4 13.54
5999 132 650 39.3 73.9 18.5 18.8 22.93
6000 133 650 3.7 66.2 13.4 13.8 116,71
6001 134 650 36.2 62.6 14.8 15.1 11.49
6002 135 650 32.3 73.4 16.6 16.9 13.29
6021 136 650 33.1 66.1 21.0 21.6 24.20
6022 137 . 650 33.0 66."7 13.8 16.5 17.43
6027 138 650 35.2 57.5 12.1 - 12.6 17.73
6028 139 650 30.5 63.3 4.3 16.0 - 17.86
6029 140 650 38.4 62.2  10.2 10.7 10.54
6030 141 650 - 31.9 65.3 13.7 1.4 17.36
5997 130 25 61.5 135.9 40.5 40.8 29.54
5998 131 25 53.3 139.1  43.5 46.0 47.81
6003 109 650 47.8 78.4 18.7 19.3 21.38
6005 109 650 49.3 79.4  19.1 19.6 14.72
6007 142 650 57.9 90.2 15.2 15.6 11.67
6008 - 143 650 49.5 80.9 11.8 12.0 9.67
6013 144 650 54.1 - 80.7 10.8 11.5 13.71
6014 145 650 bdr .5 4.7 11.0 11.4 17.31
6015 146 650 48.9 66.2 7.4 7.7 3.98
39.5 68.1 11.8 12.2 7.14

6016 147 650

 

W
 

 

 

niromtr b mmmapee i s

2"

37

| ‘Table 8. Tensile Data for Heat 466-541 (0.27% Ti)
for Various Heat Treatments '

 

Test

 

 

. i - Strength, psi . Reduction

| . Heat Temper- = Elongation, % .
‘Specimen Treatment ature <~ Yield Ultlgate Uniform Total in Area

- (°c) Tensile (%)
X 10% x 103

6090 121 650  26.5 = 68.8 29.2 30.6 28.41
6092 121 650 27.5 76.2 33.7 34.6 28.30
6121 . 123 650 21.5 544 31.8 33.4 31.22
6119 123 650 ~ -21.4  53.9 '29.7 - 30.5 41.14
- 6128 123 650  21.8  53.6 29.8 31.3 27 .48
- 609 130 650 - 35.1 77.3. 35.1 43.7 38.68
6095 131 650 33.3 - 68.8 17.8 18.3 23.17
6101 132 650 34.6 4.3 32.9 39.6 35.31
: ' ' 6102 133 650 = 31.4 63.4 15.4 16.0 20.73
- 6103 134 650 33.8  69.8  28.6 34.6 27.36
o 6104 135 650 28.3 61.0 19.0 119.5 16.01
- 6123 136 650  26.6  46.8 16.4 17.4 26.43
. - 6124 137 650 31.3 59.2 15,5 16.7 15.92
- 6129 138 650 - 29.9 - 46.1 10.4 11.4 25 .41
6120 139 650 29.7 57.6 13.6 14.9 16,03
6131 140 650 © --36.0  48.5 7.4 8.9 10.39
6132 141 650  31.0 58.1 13.1 14.7 13.57
6099 130 25 . 54.3 125.6 52.5 53.7 41.88
6100 131 25 42,4 110.5 50.1 - 52.3 40.64
6107 109 650 45.4 77.5 22.% - 23.2 17.77
6105 109 650 444 76.4 22.9 - .23.7 27 .82
6109 142 650 48,1 81.3  21.6 244 20.33
6110 143 650 45,1 T71.5 11.1 - 11.6 - 13.07
6115 144 650 - 49.3 - 81.6 - 18.0 19.7 19.39
6116 145 650 42.9 | 68.4 2.8 10.4 11.66
6117 Y6 - 650 - 52,2 8l.6  13.7. 1.5  12.60
6118 147 650 444 T73.1 11.8 12.2  11.42

 

 
 

38

Table 9. Tensile Data for Heat 466-548 (0.45% Ti)
for Various Heat Treatments

 

 

 

) Test - Reduction
| Heat = Temper-- Strgngth, pel Elongation, % in Area
Specimen Treatment s&ature Yield Ultimate Unif Total
, (°C) Tensile 110rm oLa (%)
- x10® x10? .
6268 121 650 25.2 66.6 28.2 29.2  35.43
6270 121 650 24.2  66.9 30.4 31.9 = 26.33
6299 123 650 20.3 = 54.2 29.8 32.2 27.74
- 6297 123 650 - 20.6 55.3 30.2 32.9 34.73
7400 130 650 29.7 76.6 40.4 43.3 - 40.10
7402 131 650 30.1 79.2 35.3 37.7 28.94
6279 132 - 650 29.9 4.3 39.6 42.9 - 41.38
6280 133 650 30.1 67.3 20.1 20.8 27.04
6281 134 650 29.4 71.9 35.4 43.3 31.45
6282 135 650 28.9 70.3 23.6 24.3 20.41
6301 136 650 ' 24.3 71.6 47.2 49.8 46,04
6302 137 650 - 27.2 86.5 32.3 33.5 27 .64
6307 138 650 26.1 68.6 bl 462 52.56
6308 139 650 25.8 61.8 24.9 26.2 23.42
6309 140 650 29.6 72.4 39.3 41.7  40.83
6310 141 650 - 27.0 63.9 24,9 26.2 26.33
6277 130 25 47.6 124.1 57 .4 -59.0 41.79
6278 131 25 48.1 129.6 53.6 55.4 50.95.
6285 109 650 41.7 70.1 20.7 21.7 18.04
6283 109 650 43.9  71.9 19.3  20.4 24.16
6287 142 650 45.1  84.3 27.2 30.0 22.49
6288 143 650 41.3 69.7 14.8 15.5 13.90
6293 144 650 44,3 82.0 26.8 - 28.8 - 3147
6284 145 650 40.8 68.8 14.5 14.9 16.75
6295 146 650 47.7  88.2 247 28.2 20.49
6296 147 650 42.1 71.9 15.0 15.6 12.60

 
 

o

e

 

39

Table 10. TenSile,Data;for Heat 467-548 (1.2% Ti)
' ”for Various Heat Treatments

 

Test

 

 

‘ - Strength, psi , L Reduction
Specimen Trgzzzent zzfifzr' yielg Ultimate fifiigaséii%fié_% in Area
_ s Tensile  Uniform Total (%)
| | x 10  x 10° |

6313 121 650 . 39.4 96.3 29.8 35.3 31.58
6315 121 2 650 . 33.8 - 96.7 39.3 - 42.9  3l1l.22
6341 123 650  29.7 - T4.6 - 39.4 40.8 39.95
6343 123 - 650 28.6 = 82.4 43.2 43.9 28.53

- 6317 130 650 46.5 92.1 = 28.3 50.2 45,00
6355 . 130 650 42,9 93.6 25.7 45,4 38.51
6357 131 650 = 41.1 1 94.5 26.5 b4y 7 37 .44
6318 131 650  41.8 92.8 25.9  45.0  39.99
6323 132 650 43.1 = 87.9 26.1 43.3 42.18
6324 133 650 40.8 87.1 24.5 42.1 41,05
6325 134 650 . 38.2 84.3 26.7 44,1 35.74
6326 . 135 . 650 " 36.6 83.1 25.5 4o, 1 36.05
6345 136 650 42,7  83.3 - 244 25.7 24.07
6346 137 650 ~ 46.3 71.5 . 2.6 -11.0 10.60
6351 138 650 = 44.8 79.0 17.3 20.1 27 .41
6352 139 650  44.6 0 91.7 13.3 15.4 18.51
6353 140 650  49.6 8l.4 14.0 15.2 18.29
6354 141 650 . 49.8 76.4 10.8 -11.7 9.52
6321 130 25 ~61.3 145.6 50.5 52.0 41.21
6322 131 25  50.1 121.3 47.8 49.0 43.57
6327 109 - - 650 - 73:1 0 102.2 20.0- - 27.7 30.35

- 6329 109 650 . 64.5 105.5 26.3 32.2 25.14
6331 142 650 - 64.5  104.5 21.6 -33.2 30.56
6332 143 650 - 62.0 101.9 - 20.7  32.2  27.30
6337 . 14 650  6l.4  100.3 18.3  26.4 - 25.45
- 6338 U5 650  56.1 = 9.1 - 20.6 26.5 25.79
6339 146 650  62.9  103.3 20.0  30.4  24.07
6340 147 . 650 6l.6 - 101.9 18.7  31.5  28.11
6989 137 25 59.6  122.8 41.6 417 36.73

47.6 134.9 62.5 - 63.8 46.25

om0 121 25

 
40

Table 11. AC'reep-Rupture Properties at 40,000 psi and 650°C for

Commercial Heats Used to Determine the Thermal Stability

 

Secondary

 

) ‘ ture Total Reduction
- Heat ’Trggz‘:aen & Test Rg?_fé g:::p Elongation in Area
| (nr) (%/nr) (%) (%)
466-535 121 7130 166.5 0.0133 9.50 10.60
_ 123 7132 50.9 0.0119 13.60 24.39
131 7277 136.5 0.0489 11.30 17.36
132 7368 137.8 0.0630 13.60° 14.27
133 7369 63.6 0.0683 9.50 13.19
138 7336 217.5 0.0346 13.70 15.22
139 7340 26.2 0.0868 8.30 9.51
-109 7131 . 193.8 0.0048 3.60 - 6.41
144 7464 135.4  0.0179 5.70 11.24°
_ 145 7465 103.2 0.0420 6.70 10.79
466-541 = 121 7133 341.6  0.0060 15.20 17.75
123 7135  189.0 © 0.0080 21.50 26.61
131 7276 217.0 0.0538 18.90 19.48
132 . 7370 507.6 0.0315 27.60 26.93
133 7371 67.8 . 0.0692 10.50 118.42
138 7337 50.0 0.0716 9.70 12.96
139 7341 41.2 0.0844% 9.90 17.62
109 7134 308.9 0.0043 5.90 - 13.90
144 7470 515.4 0.0159 14.00 20.73
145 7471 43.1 0.0730 6.36 11.96
: 121 7136 358.0 . 0.0117 16.70 22.31
466-548 123 7138 125.1 0.0110 19.30. 28.19
130 7024 569.6 0.0203 27.20 30.09
131 7275 59%.0 0.0251 42.10 34.19
132 7372 574.6 0.0391 41.30 41.30
133 7373 135.5 0.0535 16.70 21.94
138 7338 345.5  0.0968 29.10 37.27
139 7342 175.2  0.0660 24,10 37 .54
109 7137 620.6 0.0065 11.70 Y7.77
144 7468 855.8 0.0166 25.60 22.62
145 7469 130.7 0.0255 6.50 15,97
: 130 7024  669.6 0.0203 27.20 30.09
467-548 121 7434 1207.6 0.0144 36.60 33.44
- 123 7141 2472.9 0.0038 21.60 17.62
131 7303 1006.2 0.0195 32.40 4774
132 - 7374 1582.7 .0.0159 54.40 47 . T4
133 7375 1335.2 0.0162 36.80 35.17
138 7339 2095.5 0.0072 24..50 26.93
139 7343  1454.2 0.0099 20.40 28.64
109 7140 2149.2 0.0049 30.50 28.9% ,
144 7466  2638.1 0.0053 33.30 33.60
145 7467  1853.1 0.0089 32.10 30.33

 
 

 

s e sttt b sl

p

26.
27‘
28

30.
31.
32.
33.

35.
36.
37.
38.
39,
40.
41.
42.
43.
b,
45.
46.
47,
48.
49.
50.
51.
52.
53,

54.
- 56.

- 57.

58,

59.

60,

61.
63.

65,

66.

29.

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115. D. L. Manning , 150. J. L. Scott
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Park, CA 91305 :

E. E. Stansbury, Department of Chemical and Metallurglcal
Engineering, University of Tennessee, Knoxville, TN 37916

D. K. Stevens, RD, AEC, Washington, DC 20545
-R. F. Sweek, RDT' AEC, Washington, DC 20545

A. Taboada, RDT, AEC, Washington, DC 20545

G. A. Whltlow,_Westlnghouse, ARD, Waltz Mill Slte, P 0. Box 158
‘Madison, PA 15663 -

M. J. Whitman, AEC, Washlngton, DC 20545

R. F, Wilson, Atamlcs Internatlonal ‘P. 0. Box 309 Canoga Park,

CA 91304 .

- Director, Division of Reactor Llcen81ng, RDT AEC, Washlngton,

DC 20545 :
Director, D1v1510n of Reactor Standards, RDT AEC, Washlngton
- DC 20545
Laboratory and Unlver31ty Division, AEC, Oak Rldge Operatlons
Division of Technical Informatlon Exten81on : ; :