Note: Descriptions are shown in the official language in which they were submitted.
ROD 15789
NICKEI,-BASE SUPERALLOY SYSTEMS
Background of the Invention
Nickel-base alloys both cast and forged are
extensively used in the design of turbine components
requiring weld ability and high temperature capabilities,
particularly those alloys providing a good combination of
strength and ductility.
High-strength nickel base superalloy, which
usually contain aluminum and titanium as the major
hardening elements are strengthened by the precipitation
of gamma prime (I ') phase with ordered ice structure.
When aluminum and titanium are partially or completely
replaced by niobium or tantalum, a different precipitation
phase can be produced having toe ordered bat structure
designated as gamma double prime (I "). These "-
strengthened alloy systems provide remarkably good
tensile properties to intermediate temperatures.
Inconel 718 (IN 718), also referred to herein
as the "base alloy", contains 25~ by volume, more or
less, of the I" phase as well as a small amount of
ordered ice I' precipitates. Investigations utilizing
transmission electron microscopy have established that
coherent " precipitates are in disc-shape morphology
with a (100) habit plane and have a cubic-cubic
orientation relationship with the ice matrix. More
detailed characteristics of the phase chemistries of
' and " are given in "Phase Chemistries in Pro-
cipitation-Strengthening Superalloy" by EEL. Hall, YAM.
Koch, and KIM. Clang [to appear in Pro. Electron Microscopy
~2~3~
RD-15789
Society of America, August 1983]. The chemical combination
of IN 718 alloy is set forth in TABLE I.
TABLE I
element wow at %
No bet. bet.
Or 18.6 20.7
Fe 18.5 19.2
My 3.1 1.9
5.0 3.1
To 0.9 1.1
Al 0.4 ox
C 0.04 0.19
Despite the relatively low volume fraction of
strengthening phase (~25%) therein, IN 718 alloy, when
forged and heat treated, has a room temperature yield
strength of 165 ski, which is higher than that of Udimet 700
(~140 ski), which contains US volume % I' precipitate. This
unique strength characteristic is responsible for the
extensive use of IN 718 alloy in many turbine engine apply-
cations.
In addition to its strength and ductility capably-
it'll, another notable property of IN 718 alloy is its
excellent weld ability, a characteristic which is apparently
related to the sluggish precipitation kinetics of the
coherent I" strengthening phase. This characteristic is of
particular importance, because some welding processes are
mandatory in the manufacture and repair of certain turbine
engine component. Most precipitation-hardening super-
alloy, when welded, develop cracks in the heat aff~ted
zone and in thy weld metal during welding or during
I ,
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post-weld heat treatment. Cracking accompanying the welding
operation or subsequent heat treatment causes excessive and
costly reworking of welded components and prevents optimum
design latitude for components requiring joining during
fabrication. IN 718 alloy is known to be the only nuances-
c~ptible alloy that also provides adequate strength. It is
for that reason that IN 718 has been selected as the base
alloy against which improvement is to be measured herein.
Unfortunately, the tensile strength of IN 718
alloy is relatively sensitive to temperature compared to
conventional I' strengthened alloys. Further, the stress
rupture life of IN 718 deteriorates rapidly at temperatures
in excess of 1200 F. There is a continuing demand for new
high-strength weldable, cartable, forceable superalloy
having improved temperature capability for operation above
1200F, because of the continuing increase in the turbine
engine operating temperature.
The problem of providing weld ability in a nick
Colby cast alloy is addressed in U.S. 4,336,312 - Clark
et at. In accordance with the Clark et at. invention,
conventional nickel-base cartable superalloy are modified
by reducing the aluminum content and increasing the carbon
content thereof. In addition, as-cast modified nickel-base
alloy components are subjected to a prowled thermal condo-
toning cycle, which is believed by the patentees to resulting a precipitate that retains adequate ductility within the
gnat no .
U.S. 3,046,1C8 - ~iselstein is directed to a
malleable, age-hardenable, nicXel-chromium base alloy in
which the emphasis is on the presence of "controlled and
coordinated amounts of alloying elements" (column 1, lines
45 and 46). The composition of IN 718 lies within the
teaching of they'll patent. The exclusion of iron, the
~2~3~ US
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inclusion Of tantalum and the inclusion of cobalt are merely
options.
Certain terminology and relationships will be
utilized herein to describe this invention, particularly
with respect to the precipitation hardening elements such as
aluminum, titanium, tantalum and niobium. The approximate
conversions of weight percent to atomic percent for nick
Colby superalloy are et forth a follows:
Aluminum (wit%) x 2.1 = Aluminum (at)
Titanium (wit%) x 1.2 = Titanium (at%)
Niobium (wit%) x 0.66 = Niobium (at%)
Tantalum (wit%) x 0. 33 = Tantalum (at%)
The hollowing are definitions useful in understanding this
invention:
Iota% TOTAL" is the term representing the total
content of aluminum, titanium, niobium and Tanya-
lump expressed in atomic percent.
"Rgdp" is the value of the sum of the niobium and
tantalum contents (in at%) divided by at% TOTAL.
When this value is 0.62 or greater I" is the only
precipitation strengthening phase present.
The following U.S. patents disclose various
nickel-base alloy compositions: U.S. 2,570,193; U.S.
2,621,122; U.S. 3,~61,426; U.S. 3,151,981; U.S. 3,~6,412;
U.S. 3,322,534; U.S. 3,343,g50; U.S. 3,575,734; U.S.
~,207,098 and U.S. 4,336,312. The aforementioned U.S.
patents are representative of the many alloying situations
reported to date in which many of the same elements are
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I
RD-15789
combined to achieve distinctly different functional rota-
tionships between the elements such that phases providing
the alloy system with different physical and mechanical
characteristics are formed. Nevertheless, despite the large
S amount of data available concerning the nickel-base alloys,
it is till not possible for the metallurgist to predict
accurately the physical and mechanical properties of a new
combination of known elements even though such combination
may fall within broad, generalized teachings in the art.
Description of the Invention
Major alloying modifications of the base alloy
have resulted in new alloys for the production of weldable
castings and, further, of weldable, cartable, forgeahle
alloy heat treatable to produce an improvement of greater
than 100F in high temperature capabilities over the base
alloy. A number of criteria to provide weld ability have
been determined for this new alloy system: at% TOTAL is to
be between about 5.0 and about 8.0; the value of Rgdp is to
be equal to or greater than about 0.62 and equal to or less
than 0.95; the sum content of aluminum and titanium (i.e.,
Al + Tip is to be equal to or less than about 3.0 at% and
equal to or greater than about 0.5 at% and the sum content
of niobium and tantalum (i.e., Nub + Tax is to be equal to or
greater than about 3 0 at% and equal to or less than about
7.5 at%, thereby assuring that the alloy will be free of
gamma prim phase. In order to add to the weld ability
property certain desired high temperature capabilities (high
temperature strength and stress rupture strength), it is
preferred to eliminate iron as a constituent except insofar
as it may be present as an impunity. Limited amounts of
iron ., lies than about 5.0 wit%) may be tolerated
RD-15789
realizing that some minor reduction in high temperature
properties may be incurred. To optimize the increase in
high temperature strength and stress rupture life afforded
by this invention, Or, Co and To are added in amounts
5 ranging from about 18 wit% to about 22 wit% Or, from about 8.0
wit% to about 14.0 wit% Co and a minimum of about 2.0 White Tax
In its overall compositional definition, the
nickel-ba~e alloy of this invention contains (in White about
12% to about 24% chromium, about 5% to about 20% cobalt,
about 1% to about 8% from the group consisting of molybde-
numb tungsten, rhenium and mixtures thereof, about 2.0% to
about 23% tantalum, up to about 10.5% niobium, up to about
2.7% aluminum, up to about 3.7% titanium, about 0.003% to
about 0.05% boron, up to about 0.10% carbon, up to 0.1%
zirconium, up to about 5.0% iron, up to about 0.5% silicon,
up to about 0.5% manganese and the balance essentially
nickel. In respect to nickel the term "balance essentially"
it used to include in addition to nickel in the balance of
the alloy, small amounts of impurities and incidental
elements, which in character and/or amount do not adversely
affect the advantageous aspects of the alloy. Molybdenum
may be replaced in part or entirely by an equal weight
amount of tungsten anger rhenium. Iron is an undesirable
element in alloys of this invention and its content level
must not exceed about 5.0 wit%.
In a preferred overall compositional definition,
the nickel-base alloy of this invention contains (in White)
about 16% to about 24% chromium, about 8% to about 16%
cobalt, about 1% to about 8% from the group consisting of
molybdenum, tungsten and mixtures thereof, about 2.~5% to
about 22.5% tantalum, up to about 10.1% niobium, up to about
1.45% aluminum, up to about 2.54% titanium, about 0.005% to
about 9.02% boron, up to about 0.04% carbon and the balance
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RD-15789
essentially nickel. The minimum content of Al To is about
Owe% and the minimum content of Nub To is about 4.70%. The
maximum content of Al To is about 2.54% and the maximum
content of Nub + To is about 22.5%. Impurities, which may be
present in the alloys of this invention, include iron,
silicon, manganese, sulfur, copper and phosphorus. The
maximum permissible concentrations of these elements as
impunities are as follows:
Iron ... l.00 wit%
Silicon ... 0.35 wit%
Manganese ... 0.35 wit%
Sulfur ... 0.015 wit%
Copper ... 0.30 wit%
Phosphorus ... 0.015 wit%
Brief Description of to Drawing
The features of this invention believed to be
novel and unobvious over the prior art are set forth with
particularity in the appended claims. The invention itself,
however, as to the organization, method of operation, and
objects and advantages thereof, may best be understood by
reference to the following description taken in conjunction
with the accompanying drawing wherein:
Fig. l it a graphic representation of measured
comparative tensile and yield strengths (1) of the base
alloy and (2) of the base alloy modified by removing iron
and introducing 1 at% tantalum;
Fig. 2 is a graphic representation of investiga-
lions crusade out to study the effect of alloying modifica-
Sheehan of the base alloy on the creep rupture properties
thereon;
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RD-15789
Fig. 3 is a graphic representation of the rota-
tionship between rupture fife and yield strength of a cast
optimal alloy composition subjected to a number of thermal
processes;
Fig. 4 is a graph schematically displaying the
relationships between (Al + Tax and (Nub + To), expressed in
at%, required for the production of weldable alloys accord-
in to this invention;
Fig. 5 is an enlargement of the portion of Fig. 4
bounded by ABCDA;
Fig. 6 is a graphic representation of yield
strength (0.2% YE) data obtained in tests at 1300F for
compositions HOWE through HOWE located in region ABCDA of
Fig. 4;
Fig. 7 it a graphic representation of tensile
strength (US) data obtained in tests at 1300F for the same
compo~itiohs for which data are given in Fig. 6;
Fig. 8 is a graphic representation of yield
strength (0.2% YE) data obtained for compositions HOWE
through HOWE to demonstrate the changes in this parameter
wit changes in cobalt content, the tests being conducted at
1300F on sample previously annealed and aged;
Fig. 9 is a graphic representation of tensile
strength (US) data obtained for the same compositions for
which data are given in Fig. 8, the tests being conducted at
1300F on samples previously annealed and aged;
Fig. 10 is a graphic representation of rupture
life data obtained for the same compositions for which data
art given in Fig. 8, the tests being conducted at 1300F and
90 ski on sample previously annealed and aged;
Fig. 11 is a graphic representation of yield
strength ~0.2% 'IS data obtained in tests similar to those
RD-15789
conducted in fig. 8, the tests being conducted at 1300F on
samples previously exposed to 1300F for 1000 his;
Fig. 12 is a graphic representation of tensile
strength (US) data obtained for the same compositions for
which data axe given in Fig. if, the tests being conducted
at 1300F on samples previously exposed to 1300F for Lowe
his;
Fig. 13 is a graphic representation of rupture
life data obtained for the same compositions for which data
are given in Fig. 11, the tests being conducted at 1300F
and 90 ski on samples previously exposed to 1300F for 1000
ha;
Fig. 14 is a graphic representation of yield
strength (Owe% YE) data obtained for compositions HOWE
through HOWE to demonstrate the changes in this parameter
with change in chromium content, the tests being conducted
at 1300F on samples previously annealed and aged;
it. 15 is a graphic representation of tensile
strength (US) data obtained for the same compositions for
which data are given in Fig. 14, the tests being conducted
at 1300F on samples previously annealed and aged, and
Fig. 16 is a graphic representation of rupture
life data obtained for the same compositions for which data
are given in Fig. 14, the tests being conducted at 1300F
and 90 ski on samples previously annealed and aged.
Manner and Process of Making and Using the Invention
,, . . . _
In the development of the base alloy, iron ~18-20
wit%) was added to maximize room temperature yield strength.
The main effect of introducing iron into the base alloy is
to control the volubility of hardening elements at aging
temperature. By not introducing iron the degree of
1 I rJS
RD-15789
supersaturation is reduced. This results in a reduction in
the amount of precipitation phase, which can form, and
thereby in a decrease of yield strength. It was found in
the making of the invention disclosed herein that the
decrease in supersaturation by leaving out the iron can be
restored by adding more of the precipitate-forming element.
Thus, it has been found that tantalum, as well as niobium
(columbium), can form the I" phase in nickel-base super
alloy About 1 at% of tantalum is sufficient to compensate
for the decrease in yield strength caused by the removal of
iron from the base alloy.
FROG I NAGS CAMP Elm
Measurements of the tensile properties of a
forging of such an alloy (i.e., -Fe + l at% Tax over the
temperature range from room temperature (i.e., 68-70F) to
1400F are plotted in Fig. 1, which also includes the
requisite data fox the base alloy in the forged condition.
The tensile strength and yield strength test results of the
(-Fe Tax forging is represented by curves _ and c, respect
lively. Curves b and d represent the tensile strength and
yield strength, respectively, of the base alloy. Commercial
forging practices were used.
As may be observed in Fig. l, in the iron-free,
tantalum-modified alloy system:
l. With the same room temperature yield
strength, a higher ultimate tensile strength
is developed whereby this alloy system can
sustain more plastic deformation (i.e., curve
a vs. curve b).
2. With the same room temperature yield
strength, a better strength level is attained
--10-- .
I
ROD 15739
at intermediate temperatures, i.e., the
alloy system becomes less sensitive to
temperature (i.e., curve c vs. curve d).
Extensive investigations were carried out to
study the effects of individual alloying elements on the
creep rupture properties of the base alloy forgings.
Results of some of these investigations are shown in
Fly. 2 wherein comparisons are made to the base alloy.
Values shown along the vertical axis in Fig. 2
are values of rupture stress and the values given along the
horizontal axis are values of the Larson-Miller rupture
parameter (P). This latter term is defined by the
relationship:
P = (T -I 460) x (22 log t) /1000
where
T is temperature (OF)
t is rupture time (his.).
The rupture properties of the base alloy forging is
represented by curve m. By fixing t=100 hours, rupture
curves _, _, p and q were plotted to provide a measure of
whether or not an alloy being compared to the base alloy
does, in fact, reflect improvement in performance at
higher temperatures. As shown, the curves are plotted at
50F intervals. Test data from these investigations are
superimposed on Fig. 2 and the extent of temperature
improvement can be readily seen thereon.
The following conclusions have been reached
from these data:
1. The addition of cobalt to the (-Fe + tax
alloy in proper amounts can improve rupture
life remarkably; thus, introducing 12 wit%
cobalt provides more than an order-of-
magnitude increase in stress rupture life
at 1200F, and
RD-15789
2. Increasing the hardening element content
(e.g., Tip Tax can improve the alloy strength
and subsequently increase rupture life.
However, the improvement from adding tile-
Nemo, or tantalum (without cobalt addition)
it limited.
3. The refractory elements (My, W, Rev have very
little effect on the tress rupture proper-
ties.
1 a CASTINGS COMPARED
Because of the difficulties encountered in the
case of forged specimens, but not in the case of cast
specimen , in relating results obtained in tests on one
composition to a different composition, the more comprehend
size studies ox the individual and combined effects of
alloying elements were performed using as-cast alloys after
appropriate heat treatment. Conclusion reached from the
testing of cast alloys are applicable as well to forged
at Lowe .
In the effort to accomplish the goal of modifying
the cast base alloy to produce a new alloy system yielding
lo a weldable cast alloy and (2) a weldable cast alloy
with improved high temperature live., base alloy + l00F~
capabilities, four candidate alloy compositions were select-
Ed A 3-1/2 in. diameter, 30 lb. cylindrical ingot of each
alloy was melted in a vacuum induction melting (VIM) fur-
nice. The chemical compositions of these four alloys are
jet forth in TABLE II.
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~2~3~
~D-15789
_ _ c L
_ O O
an L I
E
to 3 o us
to- I ED r_
Of 'I o O O _ N Ox
00 00 00 00
O O O 0 00 0 0
00 00 00 00
Lo O O O O O O
00 00 00
It
"I 0 o cut
_ Ox
owe it
I_ O 3 O N
No N ---- N N ICKY
--I I, . ,
O Ox
0
0 3 O CUD
O t-- O-- O 0 0 Us
of N N N --O Us J
O O I
'I I-- I I'`' 03
O Ox 0 I 3
I _ _ _
Sol
I N N N N
-- I I I T
I
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I
ROD 15,789
The composition of each alloy is set forth for each alloy
designation both as wit% (upper set of figures) and at%
(lower set of figures). SHEA is a low volume fraction
I' precipitation strengthening alloy; SHEA is a
modification of the base alloy in that (a) iron has been
deleted, (b) cobalt has been added (12 wit% and (c) tantalum
has been added (3 wit%). These changes in the cast base
alloy improve the tensile and creep strengths at elevated
temperature without diminishing the slow aging character-
is tics of the I" strengthening mechanism.
A macro (~0.225" thick) slice was cut from the
center of each of the 3-1/2 in. diameter ingots. A
slice adjacent to the initial slice was cut from the
bottom of the top half of each ingot and a slice adjacent
to the initial slice was cut from the top of the bottom half
of each ingot. The top half of each ingot was homogenized
at 2150F/4 his. and air cooled (ARC.). The bottom half of
each ingot was hot isostatically pressed (Whopped) at
2125-2150F/2 hrs./15 ski. Later, the slices were subjected
to the same homogenization or hot isostatic pressing treat-
mint and held for later studies. Small sections of each
ingot half were heated to determine the I' or I" solves
temperature. One hour heat treatments were performed
starting with 1900F, the temperature being increased by
25F to a maximum of 2050F. Optical metallographic
examinations of these specimens revealed that the solves
temperatures of SHEA and SHEA were below 1900F, while the
solves temperature of SHEA was in the range of 200F-
2050F and the solves temperature of SHEA was above
2050F.
Based on the solves temperature, the case alloys
were subjected to the following heat treatment: alloys
SHEA and SHEA were heat treated in vacuum at fly hr.
and then at 1400F/5 his., followed by furnace cooling to
- 14 -
,, s -
I
ROD 15789
1200F at fry., upon reaching 1200F the alloys were
held at temperature for 1 hour. Alloy CHIHUAHUAS was heat
treated in vacuum at fly hr., air cooled and then
heated at 1600 F/4 he's., followed by air cooling. Alloy
SHEA was heat treated in vacuum at fly hr., air
cooled, heated at 1600 F/4 his., air cooled and then heated
at 1400F/16 his. and air cooled.
Creep and tensile specimen bars were fabricated
from the ingot after the heat treatment. The bars were
fabricated from the ingots so that thy central axis of the
completed bars had been parallel to the cylindrical axis of
of the ingot. The specimen geometry and dimensions were the
same for each bar fabricated. Tensile properties were
evaluated at room temperature and at 1300F; creep proper-
lo tie were evaluated at 1300F/90 ski.
The results of tensile and creep rupture tests are summarized in TABLES III and IV. The alloy SHEA showed the
bet ten tie properties at room temperature and at 1300F
among the four experimental alloys evaluated.
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I So
RD-15789
r O
I 0 Us N a O O N D Jo Ox O Us O O
I r--3 0-- I N O +
I N N --N N N N I I-- I-- Us
Z I N I Cry O O O . TV
N N ON I Jo 00 CUD I N No O I . C L L L
__ _ _ _ arc
. Inn-- us 00 ox OX OX I.. 0
O It _ _ _ ¦ O O Ox O Ox MU O O
O N + +
Jo 3 I O 0 O I O JO N O Us N N E O
ON O O 3 N J 3 _ I _ _ O O N Ox _ L L
1- -- -- -- __ __ __ NO c
I Ox It I N O Us O O O O 0
--Jo O¦ Irk N N _ _ _ I" I N N O CLUE 0 I 2
. _ _ _ I _ _ _ _ _ _ _ _ _ _ _ _ _ ED N N
Q r _ _ . I> . O O . . O O I/: N O v o v -
1- I ID l_ ._
Z N N N N N N N N ¦ N N N N N N N Q O =-- Q r
Z . . . T I N N
O O ' g O O to O Jo >.
U: O --O -- I T T T l O _ 2 T T I to US
J N r N _ I_ Z
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'?~ I. y I.
I
RD-15789
. Y
3 ~"~
-- Jo o I DO L L
3 I 0. I D t_ Jo N I
I 1 3 i ¦ N J . . _ -- E O .
us _ ¦ , l us I ED I-- C 6 O
I: I t-- ¦ or N L -- L L "t
N Jo ION N _ I oh
ox _ 80 a o O
3 C C C NO o
j_ 0 o o 3
O- O O - V I
J to ::~ O O C L 0 0
--O N
of I 3 aye 0 0 + Us Of
Us I N 'J 01 N N O O I L I C
'-- 0 E 0 o ) L Lo.
0 - >0>5 Ox ~0>0
on _ Q v_ ED
W T T E I T I *
J N I : 1.. 1
l 12 N Z
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~233t~'7,S
RD-1578g
The SHEA alloy at 1300F exhibits values for ultimate
tensile strength (US), 0.2% yield strength (YE), elongation
(LONG) and reduction of area (ROY.) comparable to the
values displayed by specimens of IN 718 prepared as both
Cast to Size CUTS and Cut from Casting (C.F.C.) specie
miens and tested at 1200F. Manifestly the data displayed
herein employs C.F.C. specimens of SHEA The data for cast
I It 718 is CUTS data, which is known Tao higher test
values than C.F.C. data. Thus, even on this disadvantageous
basis of comparison the SHEA alloy displays a EYE ad van
tare over cast IN 718.
The SHEA alloy exhibited lower tensile properties
than SHEA, though it had a high tensile ductility indicate
in good weld ability. The C~-23 and SHEA alloys, which
were compositional modifications of Rune '41 and Rune '63
respectively, displayed tensile and creep properties equiva-
lent to the cast Renew alloys. Notably, the lower carbon
levels of these alloys do not appear to degrade the tensile
and creep properties.
The creep rupture test data in TABLE IV display
results at the test conditions of 1300F/90 ski during which
the time varied from 22.8 hours to 232.5 hours.
Having established the superiority of cast SHEA
alloy relative to the other three cast alloys tested, a
US property comparison was made with IN 718 by testing these
two cast superalloy in parallel. As established by high
temperature tensile strength and stress rupture life tests
shown in TABLE V and VI the SHEA alloy shows a clear-cut
advantage over IN 718. It should be noted that significant-
lye greater loads were applied to the SHEA specimen than to
the IN 718 specimen in the stress rupture tests.
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Compositional, ingot processing and thermal processing data
follow. Tests were conducted on ~0.225" thick specimens.
Alloy Composition-
OH 22 (#33) - Ni~18Cr-12Co~3.OMo-5.ONb-3.OTa-l.OTi-0.5Al-
O.OlB-O.OlSC
IN 718 (#34) - Ni-19Cr-19Fe~3.0Mo-5.1Nb-0.9Ti-0.05A1-
0.006B-0.003C
Inn Processing.
Vacuum Induction Melting
Casting: Cylindrical Cut mold 3-5/8" diameter x 8-1/2"
length
HIP: 1150C/15 ksi/4 his
Heat Treatment
SHEA (#33) - 1075C, lhr/water quench + 750C, 8
hr~/furnace cool 650C, 10 hrs/water quench
IN 718 (#34) 950C, 1 hr/water quench -I 720C~ 8
ho furnace cool 620C, 10 hrs/water quench
TABLE V
(Tensile)
. _ . .
JEST 0.2% YE US LONG ROY
ALLOY TEMPT I ski) (%) I
SHEA 1000 126 133 7.3 60
~#33~ 1200 135 139 13 43
IN 718 1000 111 121 16 19
(#34) 1200 115 118 12 62
, _ . . .. _ _
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TABLE VI
(Stress Rupture)
TEST Rupture Life L.-M.* LONG ROY.
ALLOY CONDITION __ _ _Lhr2~ 25) - (%~ -~%~
SHOWOFF ksi118 47.65 6.0 7.4
(~33)1200F/100 ski 811** 46.33 0.22 ---
IN 7181300F/75 ski 20 46~29 5.1 7.8
(#3~1200 F/90 ksi214 45.37 6.7 9.8
....
Larson Miller rupture parameter
**Runt
In addition to the superior performance of the
SHEA alloy V3. IN 718 displayed for the parallel testing
repented in TABLES V and VI, a comparison of TABLES III and
V provides additional insight into the improved capabilities
provided by alloys of this invention. Thus, the SHEA alloy
at 1300 (TABLE III) exhibits values for ultimate tensile
strength (US), 0.2% yield strength (YE), elongation (LONG)
and reduction of area (ROY.) comparable to the values
displayed by specimens of IN 718 at 1200F. Manifestly, the
cat SHEA alloy (as heat treated for tests of TABLE III)
exhibits a 100F~ advantage over cast IN 718 (as heat
treated for tests of TABLE V) for these parameters.
Phase stability studies were made on the unseaters-
sod samples after their exposure at various temperatures and
times. After the heat treatment exposure, tensile specimens
were machined and tested at 1300F to ascertain the effect
of time and temperature on the stability of SHEA alloy.
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The tensile properties of SHEA alloy after long term
exposure are set forth in TABLE VII below.
TABLE VII
LONG TERM US 0.2% YE LONG ROY.
5 EXPOSURE (I, (ski) (%) (%)
1300F/1000 his 125 124 6.4 40
132 131 8. 3 35
. . . _ _ . .
101400F/216 his 134 126 7 . 8 7. 7
1400F/5û0 ho 118 108 3.3 5.2
Two rupture tests at EYE ski were conducted
on exposed SHEA samples from the ingot prepared for tests
reported in TABLES III and IV and the results (shown in
TABLE Via) of these tests indicate that the rupture lives
are longer than those of the unexposed samples of SHEA
(TABLE IVY These observations establish that alloys of
20 this invention exhibit excellent thermal stability at
temperatures up to 1300 F.
-21-
28~ 5
ROD 15789
ABLE Viva
LONG TERMRUPTIJRE LONG ROY.
EXPOSURE HOURS (%) (~)
1300F/1000 his 19~ 5.6 12
159 2.7 4.3
_ _ . _ . . . _ _ _ _ . _ _ .
Comparison of TABLES VII and Viva with TABLES III and
IV suggest that the alloys of this invention can
be heat treated to still further improve both
their high temperature strength and their rupture
properties. These properties are both of great
value in alloys used in the manufacture
of turbine engine parts.
Heat treatment and aging studies were
performed on SHEA alloy to identify and standardize
thermal processing parameters for enhancing the
strength and stress rupture life of alloys
encompassed by this invention. The results of
the effects of two thermal processes (Schedules
A and B) on the tensile and rupture properties of
SHEA are shown in TABLE VIII. These results
together with SHEA data from TABLES III, IV and
Viva are displayed in Fig. 3. The heat treatment
(solution anneal plus aging) of Schedule B is
considered a feasible and very effective thermal
processing sequence for the alloys of this invention.
Results for the testing of IN 718 are located as a
point on Fig. 3. Desist the significantly more
severe rupture test conditions for the SHEA alloy,
the Schedule B heat treatment for this alloy produces
(as compared to IN 718) an alloy of greater strength
and significantly longer rupture life.
- 22 -
` ;
RD-15739
TABLE VIII
Heat Treatment:
Schedule A - 1075C, 1 hr/W.Q. + 750C, 8 hrs/F.C. -I 650C,
10 hrs/W.Q.
Schedule B - 1075~C, 1 hr/W.Q. + 775C, 4 hrs/F.C. 700C,
10 hrs/W.Q.
Tensile (1300F~:
0.2% YE ITS LONG ROY.
(ski) (I
lo Schedule A 111 116 16 44
108 ~11 12 46
Schedule B 121 122 7 13
124 129 26 64
Radiator Ll300 F/90 ski):
LOOPHOLES. LONG ROY.
- his Parameter (%) (%)
P25
Schedule Aye 46.96 4.7 11
33.7 46.69 5.3 16
Schedule B89.6 47.43 5.8 lo
247.4 48.21 5.1 12
Weld ability texts were conducted on plates sliced
(about .225 inch thick) from each ingot prepared for tests
reported in TABLES I I I and IV in both the homogenized and
hot i~ostatic pressed condition. Two grooves, each about
3/4-inch wide were machined into one surface of each plate
and two additional grooves were machined, spaced apart, into
the opposite surface of the plate with top and bottom
grooves being in alignment with each other. The stock
remaining in the juxtaposed depressed region was about 0.06
--23--
33~7S
RD-15789
inches thick. A series of electron beam (EN) welds and
tungsten inert gas TWIG) welds were made lengthwise of the
0.06 inch thick stock. Visual inspections were made for
welding cracks before and after each welding pass and heat
treatment employed subsequent to the welding. TABLE IX
summarizes the results of these weld ability tests.
-24-
I
ROY
-- -- N
I I_
. .
'I z A:
(nQI Al us - Z I
~31 awl O O 'i
I Z 2: "I JO
I Z Z Jo,
3 e m Al j L I- o
a N we Z Z Z : F . . .
. _ Ill L
w I '' Z o I) ONE
e I E O + O O
I _ 2 O O O o _ 3 11~
O I
an owe
-- I or Z Z it I N .-- '- SO L Jo
Jo Q U .
to :
01 o _ o o z o o I__
' ox g ox ox
o
I-- I I-- a _ 1-- J
of 1~1 N N N *
T 2
so
I
. RD-15789
The SHEA alloy was the most weldable alloy. Only one crack
was observed in the TWIG welding after the third-weld-plus-
heat treating cycle. The SHEA alloy is the next best alloy
followed in turn by SHEA and SHEA.
Another set ox specimens for weld ability tests
were prepared as plates as described hereinabove and homage-
Ned at 2150F for 4 hours. A series of EN welds and TWIG
weld were made in passes perpendicular to the grooves with
all weld penetrating the plates. The EN passes each
extended across both grooves; the TWIG passes each extended
across one of the grooves. Visual inspections were made for
welding cracks after each welding pass. TABLE X summarizes
the results of these weld ability tests setting forth the
number of cracks, if any, counted for each pass. These
alloy identified in TABLE X as to at% TOTAL and Rgdp are
located on Fig. 5, which is the enlargement of a portion of
Fig. 4. The balance of the contellts of these alloy compost-
lions are substantially the same as for the SHEA alloy
except that change in Sal + To + Nub + Tax are accommodated
by varying the No content.
___
- -26- ,
33~
RD-15789
TABLE X
_ .,_~.
CRACKS CRACKS
EN TWIG
ALLOY at TOTAL R WELDS WELDS
do- -
WOW 5.5 0.63 0 2
HOWE 5.5 0.91 0
.
HOWE 7.5 0.64 0 6
_ _ _ _
HOWE 7.5 0.73 0 2
Wylie 7.5 0.93 0 2
. _ _ _ . _ . .
SHEA
. _ . . . .
SHEA 0
(#33)
. . . _ _ . . . . . . _
IN 718 9 3
I
_ _ _ .. . .. . _ _
Interestingly, the Al To levels in nickel base
alloys Jay be the most importallt variable affecting the
weld ability. The lower the level of Al + Tip the better the
welclability of nickel-base alloys becomes. Lowering the Al
+ To level below 2 wit% appears to be beneficial to achieve
good weld ability. Differences in weld ability appear to
exist between hot isc~static pressed specimens and homage-
sized specimens depending upon the alloy investigated. The
Ben fits of the alloyinc3 system of this invention are
-27-
it
RD-lS789
optimized in tile specific combination of elemellts in which
quantities of cobalt and tantalum axe substituted for the
iron content of the base alloy and I" phase material having
a preselected relationship of at% (Al + Tip to at% (Nub + Tax
5 it elected as the sole precipitation strengthening mocha-
noisome.
The particular relationships between at% (Al + Tip
and at% (Nub + Tax, which contribute to the excellent weld-
ability characteristics of the alloy system of this invent
lo lion are defined in Figs. 4 and 5 and discussion related
thereto. It must be appreciated that each of the defining
lines displayed in Figs. 4 and 5 actually represents a thin
longi~udinally-~xtending band to account for the inevitable
error encountered in the chemical analyses made to acquire
lo the data establishing these lines. Lines W and Y, which
pass through the origin of the graph, delineate three
different precipitationC~strengthening mechanisms (i.e., all
' mixed with I", and all I"). Thy mixed
mechanism prevails when the value of Rgdp is between about
0.35 and about 0.62, and IN 7l8 fulls into this region of
Fig. 4. In addition to having only I' phase as the precipi-
station strengthening material, another criterion displayed
in Figs. 4 and 5 is to be met for alloys of this invention
for which optimum weld ability is desired. Thus, the value
25 of at% TOTAL for such alloys is to be equal to or greater
than Abbott 5 .0 (line T) and be equal to or less than about
8.0 (line Z).
Applying these criteria, it can be seen from Figs.
4 and 5 that the (Al + Tip to (Nub + Tax relationships most
broadly encompassed within this invention fall approximately
within the area ABCDA. Preferred compositions fall
approximately within the area of the quadrilateral A, B, E,
F, A. Ropre~entative weldable alloys in addition to SHEA
-28-
33~ I
ROD- 15789
are set forth in TABLE XI. These alloys were cast and
subjected to microscopic examination whereby it was deter
mined that I" phase was the only precipitation strengthening
phase present therein. This information was utilized in
locating line Y.
In addition to data points for PEW PI, PUG and
SHEA, the data points for IN 718, Wasp alloy and IN 706 are
plotted on Fig. 4.
\ '
_ ,,
_ _ __
- --29-- ,
I
RD-15789
Jo N or 3
I: O O O
it t it
'I 'i .
.0 m Ox _ N
Z _ it ,.
I- O
_,~
t
O O O O O O
L I O Ox t t t
''1 I`
I -- Us
JO O O Ox O
'I\ I 'O it O I> O
I O N O to! O N
it O
--I ' O Cut --
O .. I O, ox
it t
t
Al t I t
L I O I 0 1
of I N 0 O I
it L)
, . ye
C C t
a - -
I' O - e
-30-
I Yo-yo
RD-15789
The numerical expression for the relationships set
forth in Figs. 4 and 5 for ABCDA are as follows:
at% wit%
Al 0 to about 3.05 0 to about 1.45
To 0 to about 3.05 0 to about 2.54
Alto 0.5 to about 3.05 0.24 to about 2.54
Nub 0 to about 6.75 0 to about 10.1
To 0.75 to about 7~50 2.25 to about 22.5
Nb+Ta 3.1 to about 7.50 4.70 to bout 22.5
Similarly the numerical expressions for the more
preferred relationships of A, B, E, F, A are as follows:
we
Alto 1.0 to about 3.05 0.48 to about 2.54
Nub 0 to about 5.65 0 to about 8.56
To 0.75 to about 6.4 2.25 to about 19.4
Nb+Ta 3.1 to about 6.4 4.70 to about 19.4
The most preferred values are the following in
which the Al to To ratio is about 1:1 and the Nub to To ratio
it about 1:0.3:
at White
Al 0.95 to 1.50 0.45 to 0.71
To 0.95 to 1.50 0.79 to 1.25
Nub 2.38 to 4.69 3.61 to 7.11
To 0.75 to 1.41 ~.25 to 4.~7
Yield strength, tensile strength and rupture life
tests were conducted using alloys HOWE through HOWE
. -31- .
Tao
RD-15789
located within the compass of area ABCDA (Fig. 5) and also
identified in TABLE both as to at TOTAL and Rgdp~
I; Changes in (Al + To + Nub + Tax are accommodated by varying
the No content. Changes in (Nub Tax content as a function
of at TOTAL are plotted as Rip in the graphs of Figs. 6
and 7. Two tests were performed at 1300F for each sample
composition and the Results of the yield and tensile tests
conducted are set forth in TABLE XII and displayed in Figs.
6 and 7, respectively. The temperature and extent of heat
treatment for each alloy it shown below TABLE XII.
TABLE XII
at% 0.2% YE US LONG
ALLOY TOTAL Rip (ski) ski I
HOWE 5.5 .6379.4 92.4 14.2
82.3 98.7 17.0
HOWE 5.5 .9112r.6 127.8 2.9
121.5 123.4 7.5
HOWE 7.5 .64122.8 123.8 5.3
133.9 137.6 5.2
HOWE 7.5 .73151.3 152.7 5.9
151.5 153.2 5.2
Hoyle 7.5 .93127.8 153.8 10.7
135.6 161.3 7.8
HEAT TREATMENT:
HOWE HOWE HOWE Hoyle WOW
Solution 975C Luke 1125C 1075C 1025C
(1 sir)
Aging 775C/4 ho + 700/10 ho
The results of tile rupture tests are shown in TABLE XIII.
Test conditions were 1300F and 90 ski. The test data is
--32--
I I
RD-15789
recast in TABLE XIV in order to better reflect the regions
of area ABCDA in which the (Nub + Tax and at% Total will
provide improved rupture life.
TABLE XIII
._ . .. . _
ALLOY RUPTURE LIFE ELONGR.A.
_ (ho) _ Lo (%~
HOWE 115 6.2 19
(J) 6~6 4.0 9.2
HOWE 34.5 3.6 7.5
No 1.0 3.8 8.0
HOWE 99.1 I 5.3
(K) 78.7 2.4 2.4
Hoyle 69.8 2.0 3.8
(Q) ~5.8 2.0 8.6
I Wylie 53.4 5.6 10
UP) 45.3 6.0 13
TABLE XIV
RUPTURE LIFE RUPTURE LIFE
(or) FOR (ho) FOR
Rip await TOTAL _ .5 at TOTAL 7.5
0.63 115/6.6 99.1/78.7
0.73 69.8/65.8
.~. 0.93 3~.5/1.0 /45.3
I
Tests were conducted to determine the optimum
range of Co. Tile balance of the contents of these alloys
WOW through HOWE) are substantially the same as for the
SHEA alloy except that changes in Co content are
-33-
? 3tJ1~y~
RD-l5789
accommodated by varying the No content. The results of yield
strength and tells to strength tests are reported in TABLE XV
and displayed in Figs. 8 and 9, respectively. Samples were
annealed and aged as indicated below TABLE XV and the tests
were conducted at 1300F. Results of the rupture tests are
shown in TABLE XVI and are displayed in Fig. lo
TABLE XV
ALLOYCOLALT 0.2% YE US ' LONG
~Wt~Q~_(ksi)_ us I
HOWE 0.00 127.0 133.5 20.9
Hoyle 4.00 126.2 131.9 10.6
HOWE 8.00 127.2 131.8 10.0
HOWE 12.00 125.6 130.3 8.1
WOW 16.00 130.0 135.6 9.9
HOWE 20.00 109.8 120.9 14.4
HEAT TREATMENT: 1075C/1 ho + 750C/8 ho + 650C/10 ho
~34~
YO-YO
R~-1578
TABLE XVI
_ _
ALLOCABILITY RUPTURE LONG
__ (wit) LIFE I
(ire
5 WOW 0.0020.27 5.8
Wylie 4 0047.14 4.7
WOW ~.0085.15 3.8
HOWE 12.00138.18 5.3
HOWE 16.00~2.23 5.6
37.46 4.2
Halsey 20.0022.78 7.1
71.37 5.1
HEAT TREATMENT: 1075C/1 11l t 750C/8 ho + 650C/10 ho
Samples of the same composition were tested at
1300F, the samples having been exposed for 1000 his at
1300F. Results of yield and tensile tests are shown in
TABLE XVII and displayed in Figs. 11 and 12, respectively.
Stress rupture tests on samples of the same composition
subjected to the same heat treatment are reported in TABLE
XVIII and shown in Fig. 13. Tests were conducted at 1300F
and 90 ski.
-35-
I
ROD- 1 S 7 89
TABLE XVI I
ALLOY COBALT . 2% YSUTS LONG
_Iwt~)(ksi I
HO- 10 0 . 00120 . 612~ . 814 . 5
121 . 3 aye . 120 . 6
HO- 11 4 . 00128 1134 . 19 . O
131.0 131.912.2
HOWE 8 . 00138 . 9141. 88 .1
134 . 0 138 . 5~3 . 8
HOWE 12 . 00 133 . 5137 . 6 3 . 7
HOWE 16 . 00 135 . 4139 . 7 6. 0
131 . 4 135 . 17 . O
TABLE XVI I I
ALLOY COBALTRUPTUREELONG
_ to AL I FE
_. Sir)
HOWE 0.0051.14 3.8
27.95 4.4
HO 11 4. 0060. 58 4.0
79 . 53 2 .
HO- 12 8 . 00 107 . 124 . 4
67.94 2.9
OW- 13 12 . 00 112 . 723 . 3
147.64 4.2
2$ HO- 14 16 . 00 60 . 41 3 . 3
93 . 18 2 . 7
Additional tests were conducted using alloys HOWE
through HO 45. The balance of the contents of alloys HOWE
-36--
RD-15789
through HOWE are substantially the same as for SHEA,
except that the lob and To contents of these alloys were 6. 5
wit% and O wit%, respectively, while the compositioll for HOWE
it the same as thy composition of SHEA. Yield strength end
5 tensile strength data for all these alloys are set forth in
TABLE XIX and in Figs. 14 and 15, respectively. The data
from rupture lift tests conducted at 1300 F and 90 ski are
reported in TABLE XX and in Fig. 16.
TAB LYE X I X
ALLOY CHROMIUM O . 2% YE US ELOl\JG
__ White ski) ski I
HOWE 12.0 53.1 69.4 39.2
54.8 67.5 32.4
HOWE 15.0 105.7 117.5 19.7
1û5.3 106.S OWE
HOWE 18.0 111.4 117.3 12.8
124.1 129.4 9.3
HOWE 21.0 118.4 123.8 7.2
124.7 127.8 6.1
HOWE 24.0 112.7 113.3 6.4
122.8 125.9 10.8
HOWE 18.0 119.0 124.0 7.5
102.2 116.5 4.5
HEAT TREATMENT: 1075C/1 ho + 750C/8 ho 650C/10 ho
-37- ,
RD-15789
TABLE XX
ALLOY CHROMIUM LIFE LONG
Jo
HO 40 12.00.00 28.0
0.00 19.0
WOW 15.02.20 5.6
1.95 3.1
HOWE 18.088.75 15.0
14.94 3.1
HOWE 21.091.86 3.1
74.05 8.4
HOWE 24.012.48 3.6
14.43 4.2
WOW 18.0165.18 7.8
90.24 3.1
HEAT TREATMENT: 1075C/1 ho + 750C/8 ho + 650C/10 ho
In a more preferred composition, the nickel-base
alloy of this invention is substantially free of iron and
contains (in wit%) about owe to about 22% chromium, about 8%
to about 14% cobalt, about 2.8% to about 3.0% molybdenum,
about 2.5% to about 3.5% tantalum, about 4.5% to about 5.5%
niobium, about 0.3% to about 0.7% aluminum, about 0.8% to
about 1.2% titanium, about 0.005% to about 0.015% boron, up
to 0.03% carbon and the balance essentially nickel. In the
optimized compost Zion twit Rgdp equal to or greater than
0.62 and equal to or less than 0.95 and at% TOTAL in between
about 5.0 and about 8.0) the minimum content yin await of Al
+ To is about 1.9% and the minimum content (in at%) of Nub +
To is about 3.1%. The maximum content (ill at%) of Al To
-38-
I S
RD-l5789
is about 3.0% and the maximum content (in at%) of Nub + To is
about 6.1%. In this optimized composition -the balance of
the contents of the alloy will be substantially the same as
for the SHEA alloy (except that W may be substituted for
some of the Mow with the balance essentially Nix The best
mode of this invention as it is now known is the composition
of SHEA (in White): Ni-18Cr-12Co-3Mo-5Nb-3Ta-lTi-0.5~1
O.OlB-0.015C. The preferred compositional relationships
between aluminum and titanium and between niobium and
lo tantalum, when expressed in at%, is the following: Alto is
about 1:1 and Nb:Ta is about 1:0.3.
The data presented herein define the following
relationships between weld ability and at% TOTAL and Rgdp
the content of Al, Tip Nub and To being set thereby) within
the area AWAKED:
1. weld ability improves as at% TOTAL is
decreased and
2. weld ability improves as Rdgp is increased.
Similarly, the effect of Co and Or content on
yield strength ~0.2% YE), tensile strength (US) and stress
rupture life establishes that given the OH 22 composition of
other components, optimum high temperature strength and
stress rupture life are obtained by using contents of Co in
the range of about 8 to about 14 wit% and/or by using con-
tents of Or in the range of about 16 to about 22 wit%.
Unless otherwise specified, percentages given rein weight percent.
- -39
