Weldable heat resistant alloy

Moderate cost air meltable, highly weldable nickel-chromium-iron alloys of high hot strength and excellent hot gas corrosion resistance especially in the service temperature range of 1600.degree. F. to 2200.degree. F. Which may be formulated to provide for good weldability before and after high temperature aging consisting essentially of: ______________________________________ Nickel 33-40% by weight Chromium 24-30% Iron 14-32% Tungsten 8-17% Carbon Up to 0.12% Manganese Up to 1% Silicon Up to 1.2% Chromium plus tungsten 34-44% ______________________________________

BACKGROUND OF THE INVENTION 
There is a great demand for large cast shapes and large cast weld 
structures in petrochemical, heat treating, and other fields. The desired 
metallic properties are long life, high hot strength, good resistance to 
thermal fatigue and thermal shock and good weldability at moderate cost. 
The nickel-base superalloys developed primarily for aircraft jet engine 
components since World War II achieve hot strength through solid solution 
hardening by inclusion in the alloys of large quantities of molybdenum 
and, to a lesser extent, niobium (columbium), plus precipitation at or 
along the grain boundaries of very fine particles of gamma prime, a 
compound formed between nickel and fairly large amounts of aluminum plus 
titanium. Since aluminum and titanium are readily oxidized in air melting 
practices, these nickel-base superalloys are melted and poured in a vacuum 
or an inert atmosphere. This requirement plus the high cost of such alloys 
make them impractical for the present application. 
Cobalt-base superalloys were also developed for aircraft jet engine 
components. These alloys derive their hot strength primarily by solid 
solution hardening by elements of the group Mo, W, Nb, and Ta, plus 
precipitation of refractory carbides along the metallic grain boundaries. 
A few cobalt-base alloys and cobalt-nickel base alloys have also employed 
gamma prime hardening. Cobalt remains a relatively scarce and expensive 
metal, and therefore cobalt-base superalloys are far too expensive for the 
large structures discussed here. 
Both cast and wrought iron-nickel-chromium-base alloys have been the 
economical choice in these applications. Wrought alloys of less than about 
0.15% C content, sometimes containing about 1.5% or less combined content 
of aluminum plus titanium, have been employed as headers, manifolds, 
cones, and transfer lines. These alloys have excellent weldability and 
good resistance to thermal cracking in service. They have relatively very 
low hot strengths. 
However, the higher hot strength, higher carbon casting alloys permit 
reduction in wall thickness, reduction of metal weight and better rate of 
heat transfer. Nickel-iron-chromium-base alloys have been modified to 
include up to about 1% Nb, up to about 5% W and up to about 15% Co for 
increased hot strength. In general those cast alloys of about 0.45 to 
0.55% C contents have higher hot strengths but very poor to almost no 
weldability, while cast alloys of about 0.35 to 0.40% C contents have 
substantially lower hot strengths but some degree of weldability. 
These carbon strengthened alloys tend to age and embrittle in service. 
Thus, they may easily crack during thermal cycling. Alloys that, on a 
creep-rupture test basis, should last about ten years have sometimes been 
found to fail in cracking after perhaps one year or so in service. Also, 
in some applications it is desirable to weld repair some components after 
periods of service. 
Alloys which include substantial amounts of iron in their formulation are, 
in general, considerably less expensive than nickel-base, iron-free alloys 
for at least two reasons. They may employ much lower cost ferroalloys to 
make up their contents of chromium, and sometimes other components, as 
contrasted to the higher cost pure chromium and other metals required in 
nickel-base alloys, In addition, the mere replacement of even a portion of 
the moderately expensive nickel by very low cost iron represents 
substantial component cost savings. In alloys of the present invention a 
third extremely important advantage of including substantial quantities of 
iron in their formulation is that they develop higher hot strengths than 
the far costlier iron free nickel base alloys of the equivalent hardness 
and weldability and of the same tungsten contents. 
Thus, there has remained a great demand in oil refineries, petrochemical 
plants, heat treating equipment, and other applications or moderate cost, 
for air meltable nickel-iron-chromium-base alloys that do not require 
large amounts of carbon in their formulation for hot strength and that 
have exceptional weldability as cast. It is further desirable in some 
instances that such alloys retain good room temperature ductility and 
weldability after periods of service at high temperature. 
English, et al, U.S. Pat. No. 2,540,107, disclosed a modification of alloy 
type HP containing 40 to 60% Ni, 22 to 34% Cr and 4 to 6.5% W and known 
commercially as alloy 22H. English, U.S. Pat. No. 3,607,250, disclosed an 
improved version known as super 22H, containing 40 to 55% Ni, 27 to 33% 
Cr, 4 to 5% W and 2 to 4.5% Co. Avery, U.S. Pat. No. 3,127,265, disclosed 
an alloy know as Supertherm, which contains 26 to 42% Ni, 22 to 32% Cr, 3 
to 16% W, 9 to 26% Co, 0.3 to 0.95% C, 0.5 to 2% Si, and the balance, if 
any, iron. In practice, this alloy nominally contains about 35% Ni, 26% 
Cr, 15% Co, 5% W, 0.5% C, 0.7 % Mn, 1.6% Si and 21% Fe. British Pat. No. 
1,046,603 disclosed a nickel base alloy known as MO-RE 2, containing 26 to 
38% Cr, 10 to 25% W, less than 1% C, less than 0.2% each of Mn and Si, and 
the balance Ni. All four of these alloys are characterized by low cold 
ductility and little if any weldability. 
Nickel base superalloys, other than MO-RE 2, and cobalt base superalloys 
have employed up to 15% W, up to 14.5% Mo, up to 5.6% Nb and up to 9% Ta 
in order to attain high substitutional solid solution matrix hardening and 
strengthening, and in many cases, to additionally form hard refractory 
carbides. There is some information in the literature concerning the solid 
solubility limits of these four elements at various elevated temperatures 
and as affected by different levels of chromium in nickel base alloys, but 
there is almost nothing reported concerning how much these solubility 
limits are affected by various levels of carbon and iron additions to 
complex alloys which otherwise contain various levels of nickel, chromium, 
manganese, silicon, aluminum, titanium and possibly cobalt. 
SUMMARY OF THE INVENTION 
It is therefore an object of this invention to provide moderate cost air 
meltable, highly weldable nickel-chromium-iron alloys of high hot strength 
and excellent hot gas corrosion resistance especially in the service 
temperature range of 1600.degree. F. to 2200.degree. F. 
According to this invention alloys are provided which consist essentially 
of: 
______________________________________ 
Nickel 33-40% by weight 
Chromium 24-30% 
Iron 14-32% 
Tungsten 8-17% 
Carbon Up to 0.12% 
Manganese Up to 1% 
Silicon Up to 1.2% 
Chromium plus 34-44% 
Tungsten 
______________________________________ 
Optionally, the alloys of the invention may further contain: 
______________________________________ 
Titanium Up to 0.8% 
Aluminum Up to 0.8% 
Zirconium Up to 0.15% 
Boron Up to 0.01% 
Cobalt Up to 0.8% 
______________________________________

DETAILED DESCRIPTION OF THE INVENTION 
In accordance with the present invention alloys are provided which have 
high hot strength and excellent hot gas corrosion resistance to 
2200.degree. F. combined with excellent weldability. They are air meltable 
and castable and of moderate cost. 
In addition to having outstanding weldability and high temperature hot 
strengths there are other advantages to the instant alloys due to the fact 
that they contain substantial amounts of iron and little or no carbon. 
Prior art nickel base and cobalt base superalloys are far too costly to be 
employable in large cast weld structures for which the alloys of the 
present invention are directed. Alloys of the present invention may employ 
the much less costly ferrochromium for their chromium contents instead of 
the much costlier electrolytic or other pure chromium sources. 
Additionally the much lower melting ferrotungsten may be employed in place 
of the very high melting point pure tungsten that is often only partially 
dissolved in production of iron-free or very low iron content alloys. 
The relatively high carbon contents of prior art nickel-iron-chromium-base 
alloys result, not only in rather high room temperature hardness 
regardless of thermal history, but also in a pronounced tendency to 
develop hot tears or thermal shock cracks during the casting and cleaning 
process itself and in service which involves thermal cycling. 
To the contrary, the very low carbon alloys of the present invention are 
characterized by surprisingly low hardness and high tensile elongations at 
all temperatures during the casting, cleaning, grinding, and welding 
processes, but their hardnesses and hot strengths can then be increased by 
a mere short term heat treatment or during the first day of exposure to 
service temperatures to a degree unheard of in carbide strengthened 
alloys. 
The Alloy Casting Institute Division of the Steel Founder's Society of 
America type HP alloy was selected as a nominal starting base composition 
in the development of the alloys of the present invention, because it is 
metallurgically stable even when its carbon content is reduced or removed 
and because it has the highest hot strength of any standard alloys of the 
series at temperatures above about 1600.degree. F. 
Alloy type HP is also the basic alloy that led to the best of the other 
prior art nickel-iron-chromium-base heat resistant alloys that have been 
employed in larger cast-weld structures. Alloy type HP contains nominally 
about 35% Ni, 25% Cr, 0.35 to 0.75% C, up to 2% each of Mn and Si and the 
balance essentially iron. 
The iron levels of the alloys of the invention were deliberately kept high, 
not only to reduce nickel content and to permit formulation with low cost 
ferroalloys, but also in order to reduce the solid solubility of chromium 
and tungsten in the matrix. It was found that tungsten additions of up to 
approximately the solid solubility limit in Ni-Fe-Cr-base alloys increased 
their high temperature rupture lives, typically by a factor of about ten, 
and that further increases of tungsten far beyond the solubility limit for 
a given Ni-Fe-Cr-W alloy composition typically further almost tripled the 
rupture life, that is, an increase of thirty folds. These rupture lives 
were even further increased by additions of certain minor elements. Thus 
the presence of substantial levels of iron in the alloys of the invention 
substantially reduced the tungsten levels required in the iron-free alloys 
to attain these various increases in high temperature service live. Iron 
levels between about 14% and 32% were found to be most desirable. Nickel 
contents of the alloys of the invention must be high enough to maintain an 
austenitic or face-center-cubic matrix crystal structure. Since iron, 
chromium and tungsten all tend to form the opposite undesirable 
body-center-cubic crystal lattice structure, a nickel range of about 33 to 
40% was found to be required in alloys of the invention. While slightly 
lower chromium levels will suffice to achieve adequate oxidation or other 
hot gas corrosion resistance up to about 2200.degree. F. in alloys of 
higher nickel contents, chromium contents between about 24% and 29% are 
required for the nickel levels of the alloys of the invention. Greater 
than about 30% Cr content results in metallurgicallyunstable alloys at 
these nickel and tungsten levels. Some nickel ore deposits also contain 
small quantities of the sister element cobalt. Also, cobalt is present in 
some high nickel alloy scraps. It has been found that at least about 0.8% 
Co may be included in alloys of the present invention without detriment. 
Molybdenum in nickel base alloys is known to form carbides when carbon is 
present and to form an intermetallic compound of Ni.sub.7 Mo.sub.6, when 
little or no carbon is present. Molybdenum is also notorious for forming 
the hard, brittle destructive electron compound known as sigma phase when 
present in many high temperature alloys formulated to contain various 
quantities of nickel, cobalt, chromium, iron and other elements. Thus, 
large amounts of molybdenum remove nickel from the matrix as well as cause 
other undesirable metallurgical effects. It has been found that up to 
about 0.8% Mo may be tolerated in alloys of the invention without serious 
detriment. In a similar manner large amounts of niobium will form carbides 
with any carbon present in nickel base alloys as well as the intermetallic 
Ni.sub.3 Nb compound in low carbon alloys. Niobium is also known to 
promote the brittle sigma phase in many alloys. It is also a very 
expensive element. Therefore, niobium was not deliberately incorporated in 
alloys of the present invention. However, it has been determined that as 
much as about 0.3% Nb may be present in alloys of the invention without 
serous detriment. 
Tantalum is actually a scarce and very expensive element and thus not 
considered as practical addition to the alloys of the present invention. 
Thus tungsten remained the element of choice for both solid solution 
hardening and precipitation hardening of alloys of the invention. To 
achieve these ends, tungsten levels between about 8% and about 17% were 
found to best attain the desired hot strengths, rupture lives, and other 
alloy properties. 
Carbon has been employed in heat resistant alloys to form large amounts of 
carbides precipitated at the grain boundaries. Its effect as an 
interstitial solid solution strengthener is ordinarily either ignored or 
unknown. Carbon is present in alloys of the invention at levels below 
about 0.12%, and preferably in the range of about 0.04% to 0.08%. At these 
low carbon levels only very minor amounts of carbides are observed under 
the microscope in highly polished samples of alloys of the invention, yet 
the increase in rupture lives of these alloys is substantially higher than 
those attained when carbon levels are below about 0.02%. Small amounts of 
carbon probably not only increase interstitial solid solution 
strengthening, but also appear to somewhat lower the substitutional solid 
solution level of tungsten for any specific iron, nickel, and chromium 
levels. Increased quantities of precipitated tungsten particles appear to 
result from the presence of very small amounts of carbon. Higher than 
about 0.08 to 0.12% carbon not only produces large amounts of undesirable 
grain boundary carbides but also combines with tungsten to lower its 
effectiveness as a hot strengthener as described above. 
Titanium, manganese and silicon also appear to lower the matrix solubility 
for tungsten at a given chromium level without forming undesirable 
tungsten compounds. However, a silicon content greater than about 1.2% in 
higher tungsten alloys of the invention and greater than about 0.9% in 
lower tungsten alloys of the invention seem to reduce high temperature 
rupture life. The same is found to be the case with greater than about 1% 
Mn at any tungsten level. Thus, silicon should be held below about 1.2%, 
and preferably below about 0.9%, in alloys of the invention, while 
manganese should also be held below about 0.8%. Small amounts of titanium 
appear to increase the amount of precipitated tungsten, but greater than 
about 0.8% Ti caused much of this precipitate to form within the grains 
rather than at the grain boundaries. Therefore titanium should be held 
below about 0.8%. Up to about 0.005% B and/or about 0/06% Zr were also 
found to enhance rupture life, but greater than about 0.01% B or about 
0.15% Zr lower cold ductility and weldability of the alloys of the 
invention. 
While precise levels of large amounts of both aluminum and titanium are 
virtually impossible to control in air melting practices due to their 
pronounced tendencies to oxidize at the molten alloy temperatures, smaller 
amounts of these two elements may be employed effectively. For example, if 
about 1% Ti is added to an alloy of the invention in the molten state and 
only a half or a quarter of this addition remains unoxidized and in 
metallic solution, the beneficial effects of the addition appear to be 
about the same for both recovery levels. Thus, it is desirable to have 
from about 0.25% to 0.70% Ti in the present alloys. In a similar fashion, 
aluminum may be employed as a deoxidizer in alloys of the invention in 
which both silicon and manganese are held to very low levels. In these 
instances the residual amount of recovered metallic aluminum is relatively 
unimportant, since the metallic bath has been reduced to a very low oxygen 
level by that part of the aluminum addition that was oxidized and 
collected in the slag. However, in very low silicon and manganese alloys 
of the invention, it is preferable to deoxidized with aluminum prior to 
the titanium addition so that some residual metallic titanium is held in 
the final alloy solution. 
It has been well known for many decades that alloys based upon some 
combination of nickel, iron, cobalt, and chromium require, for long 
service life at high stress and at high temperature, a stable austenitic 
or gamma face-centered-cubic crystal matrix structure and that alloys 
which form the brittle sigma phase will be short lived and prone to crack 
in service. Nickel, cobalt, carbon and nitrogen favor the formation of an 
austenitic structure. The metallurgical literature of several decades has 
taught that large combined quantities of elements from the group chromium, 
molybdenum, tungsten, niobium, tantalum, vanadium, titanium, zirconium, 
silicon, and aluminum in solid solution favor the formation of ferritic or 
alpha body-centered-cubic matrix crystal structure. All of the elements of 
the second group except silicon and aluminum also form carbides when 
carbon is present in these alloys. Some, if not all, of those same 
elements tend to form nitrides when nitrogen is present. As noted above, 
titanium, aluminum, molybdenum, and niobium may also combine with nickel 
to form metallic compounds that precipitate from the solid solution and 
therefore no longer enter into matrix reactions. The literature of the 
last several decades has also taught that these ferrite formers promote 
formation of sigma phase at elevated temperatures especially between about 
1500.degree. and 1700.degree. F. Therefore, in the low carbon alloys of 
the present invention large enough quantities of tungsten to form tungsten 
precipitates would surely be expected to form sigma phase when present 
with such high levels of chromium and iron, if the literature teaching of 
these decades were correct. Surprisingly, sigma phase was not detected in 
alloys of the invention that were exposed to temperatures that favor the 
formation of sigma phase. 
As noted above, the presence of tungsten in solid solution within the 
matrices of alloys of the invention increases hot strength and high 
temperature rupture life to some extent. Much greater increases in hot 
strength and rupture life are achieved at higher alloy tungsten levels 
when fine particles of tungsten precipitates are formed at the alloy grain 
boundaries. Further increases in tungsten content eventually result in 
coarse or excessive amounts of tungsten precipitates that then result in 
reductions of hot strength and rupture life. The quantities of tungsten 
that may be retained in solid solution in alloys of the invention are to 
some extent determined by levels of nickel as well as levels of minor 
elements discussed above, but to a major extent by the quantity of 
chromium present. I have discovered that the maximum hot strengths in 
alloys of the invention are achieved when the combined quantities of 
chromium plus tungsten are between about 38.5% and 43% by weight. When 
chromium plus tungsten weight per cent contents exceed this range, rupture 
lives at any temperature and stress level decline very rapidly. At any 
given stress and temperature a combined content of about 41 % Cr plus W 
appears to result in maximum rupture life of alloys of the invention. 
Since it has been found that 24% to 30% Cr is required in alloys of the 
invention for adequate hot gas corrosion to about 2200.degree. F., it 
would be desirable to formulate such alloys to contain about 17% W at 24% 
Cr level down to about 11% W at 30% Cr level. It is desirable to select 
somewhat lower tungsten levels if elements of the group carbon, manganese, 
silicon, boron, titanium, zirconium, and aluminum are present in 
combinations and levels toward the high end of their ranges. 
For maximum hot strength in applications involving minimal thermal cycling 
or rate of temperature change and requiring welding only prior to service 
exposure, the following ranges of component elements have been found to be 
preferable. 
______________________________________ 
Nickel 35-40% by weight 
Chromium 24-29% 
Iron 14-25% 
Tungsten 8.5-17% 
Carbon Up to 0.10% 
Manganese 0.3-0.8% 
Silicon 0.3-0.6% 
Titanium Up to 0.7% 
Aluminum Up to 0.3% 
Zirconium 0.03-0.06% 
Boron Up to 0.005% 
Cobalt Up to 0.8% 
Chromium plus 38.5-43% 
Tungsten 
______________________________________ 
For excellent weldability both prior to and after high temperature service 
exposure that may involve frequent an/or rapid thermal cycling but of 
lower hot strength, the following ranges of component elements have been 
found to be preferred: 
______________________________________ 
Nickel 33-40% by weight 
Chromium 24-29% 
Iron 18-31% 
Tungsten 8-13% 
Carbon 0.04-0.08% 
Manganese 0.2-0.6% 
Silicon 0.1-0.6% 
Titanium 0.3-0.7% 
Aluminum Up to 0.3% 
Zirconium Up to 0.06% 
Boron Up to 0.005% 
Cobalt Up to 0.80% 
Chromium plus 34-38% 
Tungsten 
______________________________________ 
The following examples further illustrate the invention. 
Example 1 
Three one hundred pound heats of three alloys were prepared in accordance 
with the invention by air melting in a high frequency induction furnace 
and casting into appropriate shapes for tensile test bars. The analyses of 
these heats are indicted in Table I as 10.3 W,13.1 W and 15.1 W. In a 
similar manner a 500 pound heat was prepared and cast into similar shapes 
for test bars as well as into welding test plates measuring approximately 
6" by 12" by 1" thickness, with on edge suitably bevelled for multipass 
welding. The analysis of this heat is set forth in Table I as 12 W. In a 
similar manner a 500 pound heat was prepared a portion of which was cast 
into tensile bar and weld plate castings with the analysis indicated in 
Table 1 as 14.6 W. A quantity of pure nickel was then added to the molten 
remainder of the 500 pound heat left in the furnace altering it chemically 
to raise the nickel content but to dilute the contents of other elements. 
This altered metal was also cast into tensile and weld plate castings. The 
analysis of this altered composition is set forth in Table I as 13.6 W. 
Hardnesses and tensile tests were conducted on as cast bars from each heat. 
Other bars from each heat were held at 1650.degree. F. for 24 hours, 
slowly cooled in the oven and then tested for hardness and tensile 
properties The results of these tests are set forth in Table I. 
From decades of experience with experimental and production heats of 
hundreds of different heat resistant alloy types, I have learned that 
those above about 7 or 8% room temperature elongations have good 
weldability. Weld plates of alloys 12 W, 13.6 W and 14.6 W as cast and of 
12W in the aged condition were welded together by SUPER22H welding 
electrodes. These electrodes nominally contain about 48% Ni, 28% Cr, 5% W, 
small amounts of C, Mn, and Si and the balance essentially Fe. 
All of the weldments were examined by x-ray and 10.times. magnification, 
and no cracks or defects were observed in any of them. 
TABLE I 
______________________________________ 
ALLOY DESIGNATION 
10.3 W 
12 W 13.1 W 13.6 W 
14.6 W 
15.1 W 
______________________________________ 
CHEMICAL 
COMPOSI- 
TION BY 
WEIGHT 
Ni 36.65 39.08 37.27 37.13 34.18 39.13 
Cr 28.12 25.51 27.92 28.18 29.41 27.54 
Fe 22.60 22.01 20.05 18.95 19.39 16.76 
W 10.33 11.92 13.11 13.62 14.63 15.12 
Co 0.11 0.08 0.22 0.03 0.03 0.16 
Mo 0.19 0.12 0.12 0.14 0.18 0.09 
Mn 0.76 0.52 0.41 0.74 0.75 0.33 
Si 0.82 0.61 0.48 0.89 1.09 0.41 
C 0.055 0.023 0.065 0.050 0.044 0.060 
Ti 0.26 0.05 0.27 0.22 0.25 .36 
B 0.005 0.003 0.005 Nil Nil .005 
Nb 0.05 0.03 0.04 0.05 0.05 Nil 
Zr 0.05 0.04 0.04 Nil Nil .03 
Cr + W 38.45 37.43 41.03 41.80 44.04 42.66 
As Cast 
Properties: 
Tensile 73,500 72,000 73,000 
74,000 
69,300 
72,000 
Strength, PSI 
Yield 32,000 30,000 38,000 
39,600 
45,300 
40,100 
Strength, PSI 
% Elongation 
45 48 35 27 16 24 
Brinell 108 116 136 143 156 150 
Hardness 
Properties 
Aged 24 Hrs 
at 1650.degree. F.: 
Tensile 91,000 100,000 90,000 
91,100 
82,700 
90,000 
Strength, PSI 
Yield 55,000 52,000 60,500 
63,000 
63,000 
62,000 
Strength, PSI 
% Elongation 
13 8 4 3 2 3 
Brinell 238 217 268 286 302 295 
Hardness 
______________________________________ 
Example 2 
Aged samples from alloys 12 W, 13.6 W, and 14.6 W were polished, etched, 
and examined at various magnifications under the microscope. 
A few polygonal crystals of titanium carbonitride were observed within the 
grain boundaries here and there in 13.6 W and 14.6 W samples but none were 
seen in the 12 W sample. Also, very small amounts of carbides were 
observed at grain boundaries of 13.6 W and 14.6 W but not those of 12 W. 
This suggests that all or almost all of the carbon content of 12 W remains 
in solid solution even after aging at 1650.degree. F. 
Another apparent phase was observed surrounding austenitic grains of 12 W 
and in larger quantities around grains of the 13.6 W sample and in far 
larger quantities in the 14.6 W samples. This phase is apparently related 
to the quantity of tungsten present in these low carbon alloys. The 
samples were etched to detect sigma phase, but none was apparent. All 
three aged samples registered less than 1.01 magnetic permeability on a 
high sensitivity low MU permeability indicator. Since ferrite is strongly 
ferromagnetic the above observed phase cannot be ferrite. As noted above, 
this is contrary to the literature teachings of the past several decades. 
While alloy 12 W retained good weldability after aging, its hardness levels 
before and after aging are consistent with relatively low amounts of 
tungsten precipitates. The lower nickel content and higher carbon, 
titanium, silicon, and boron levels combined with the higher combined 
chromium plus tungsten level of 10.3 W would be expected to form greater 
quantities of tungsten precipitates in 10.3 W than those of 12 W. As will 
be seen below, 10.3 W has much higher rupture life values than 12 W 
despite the lower tungsten content of the former. 
In a contrary manner, the lower chromium plus tungsten content of 15.1 W 
will be shown below to result in much higher rupture life than 14.6 W 
despite the higher tungsten content of the former. 
EXAMPLE III 
Standard 5/16 inch diameter creep-rupture bars were machined from cast 
materials of each inventive alloy of Table 1 and tested at various 
temperatures and stress levels on cantilever creep-rupture testing 
machines. The results of these tests are set forth in Table II below. 
TABLE II 
______________________________________ 
RUPTURE LIVES OF ALLOYS OF THE INVENTION 
AT VARIOUS TEMPERATURES AND STRESSES 
10.3 W 
12 W 13.1 W 13.6 W 
14.6 W 
15.1 W 
______________________________________ 
1600.degree. F. 
10,000 PSI 
-- -- 1279 -- -- -- 
9,000 PSI 
692 -- -- -- -- 1209 
8,000 PSI 
-- -- -- 1023 -- -- 
7,000 PSI 
-- 283 -- -- 442 -- 
1700.degree. F. 
7,000 PSI 
-- -- 1199 -- -- -- 
6,000 PSI 
742 -- -- 783 -- 1650 
5,000 PSI 
-- 242 -- -- 484 -- 
1800.degree. F. 
5,000 PSI 
-- -- 714 -- -- 501 
4,000 PSI 
718 -- -- 2184 331 -- 
3,000 PSI 
-- 922 -- -- -- -- 
1900.degree. F. 
4,000 PSI 
-- -- 273 -- -- -- 
3,000 PSI 
384 -- -- 1506 -- 1071 
2,500 PSI 
-- 273 -- -- 690 -- 
2000.degree. F. 
3,000 PSI 
-- -- -- -- -- -- 
2,500 PSI 
-- -- -- 482 -- 347 
2,000 PSI 
275 103 769 -- 217 -- 
______________________________________ 
These data were then employed in calculations based upon the Larson-Miller 
parameter to estimate the probable 10,000-hour rupture stresses that would 
be expected for each of these alloys at several elevated temperatures. The 
estimated rupture stresses at 1600.degree. F., 1800.degree. F., and 
2000.degree. F. are set forth in Table III below. Estimated rupture 
stresses of numerous prior art heat resistant alloys at these same 
temperatures are also set forth in Table III. These stresses are taken 
from the usual values represented in the metallurgical literature or, in a 
few instances from experimental data as presented in the literature. 
TABLE III 
______________________________________ 
ESTIMATED 10,000- 
HOUR RUPTURE 
STRESS AT 
VARIOUS 
CHEMICAL COMPOSITION TEMPERATURES, 
MAJOR CONSTITUENTS P.S.I. 
BY WEIGHT PERCENT* 1600.degree. 
1800.degree. 
2000.degree. 
ALLOY C Ni Fe Cr W Co F. F. F. 
______________________________________ 
10.3 W .06 37 23 28.1 10.3 -- 6,000 2,500 650 
12 W .02 39 22 25.5 11.9 -- 4,300 2,000 550 
13.1 W .07 37 20 27.9 13.1 -- 7,600 3,200 1,100 
13.6 W .05 37 19 28.2 13.6 -- 6,000 3,200 1,100 
14.6 W .05 34 19 29.4 14.6 -- 4,800 2,300 700 
15.1 W .06 39 17 27.5 15.1 -- 6,700 3,200 1,100 
113 MA .03 58 -- 23 18 -- 7,500 2,700 700 
KSN .02 57 -- 16 26 -- 6,100 2,200 550 
MORE 2 .20 50 -- 33 16 -- 6,600 3,000 1,100 
HA188 .10 22 2 22 14 39 6,000 2,200 650 
L-605 .10 10 -- 20 15 53 6,800 2,200 500 
HK40 .40 20 53 25 -- -- 3,800 1,700 720 
HP45 .45 35 38 25 -- -- 3,900 2,250 900 
22H .45 48 17 28 5 -- 5,500 2,400 850 
SUPER .45 48 14 28 5 3 6,600 2,900 1,200 
22H 
HP50WZ .50 35 32 26 5 -- 6,000 2,500 850 
SUPER- .45 35 17 26 5 15 7,400 3,800 1,300 
THERM 
330 .10 35 44 19 -- -- 2,400 1,000 330 
800 .07 32 47 21 -- -- 2,500 1,000 250 
600 .08 76 8 16 -- -- 2,500 1,100 370 
800H .08 32 46 21 -- -- 3,100 1,200 400 
601 .05 61 14 23 -- -- 3,000 1,200 350 
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*Almost all of these alloys contain small amounts of Mn and Si employed a 
deoxidizers; many contain small amounts of one or more other elements. 
The first group of alloys in Table III contains the experimental alloys of 
the invention. Alloys 10.3 W and 12 W had room temperature elongations of 
45% and 48% as cast and 13% and 8% as aged respectively. They would be 
readily weldable after high temperature service exposure. Thus their hot 
strengths should be compared to the last five alloys listed in Table III. 
The alloys of this last group are those commonly employed when thermal 
fatigue, thermal shock or weldability after periods of service are 
considered desirable. The inventive alloys are obviously of much higher 
hot strengths than the prior art alloys for which weldability is required 
after service exposure. It may also be seen that 10.3 W has higher hot 
strength than 12 W despite its lower tungsten content. This is believed to 
be because its titanium, carbon, silicon and chromium plus tungsten 
contents are all higher, resulting in greater tungsten precipitate 
quantities even with lower W total content. The other four inventive 
alloys may all be readily welded before aging at service temperatures but 
not after. These alloys may be compared to the second and third groups of 
prior art alloys set forth in Table III. 
The second group of alloys in Table III contains prior art alloys of higher 
and lower tungsten and chromium contents than the alloys of the invention 
which were tested and are all either nickel base or cobalt base alloys and 
far costlier than the alloys of the invention. The MORE 2 alloy is said to 
be useful up to 2400.degree. F., but does not provide hot strengths as 
high as those attainable in the inventive alloys. The other prior art 
alloys of the second group in Table III are all of lower chromium contents 
and suitable for service to 2000.degree. F. Again, it may be seen that the 
inventive alloys can be formulated to provide higher hot strengths despite 
their high iron contents and chromium levels that render them suitable for 
service to 2200.degree. F. Alloy 113 MA of this group has 1600.degree. F. 
hot strength comparable to inventive alloy 13.1 W, but the latter gains 
steadily over the former as temperatures increase. 
It may also be noted from Table III that alloys 13.1 W, 13.6 W and 15.1 W 
present the best hot strengths of the inventive alloys and have chromium 
plus tungsten contents of 41.03%, 41.50% and 42.66% respectively. Alloy 
14.6 W has lower hot strengths at 44.09% combined chromium plus tungsten 
content. Alloys 10.3 W and 12 W have combined chromium plus tungsten 
contents falling below the optimum hot strength range but gain the 
advantage of weldability after high temperature aging in service. Alloy 
14.6 W and 15.1 W. have combined chromium plus tungsten contents above the 
optimum hot strength range but have higher aged hardness levels. Thus, 
where hot abrasion is involved, it may be desirable in some applications 
to sacrifice optimum hot strengths for increased hot wear resistance. 
Five of the six alloys of the third group in Table III are typical of prior 
art alloys commonly used in these applications. They all contain 
relatively high carbon levels and have very poor to no weldability as 
cast. The Supertherm alloy has hot strength above the alloys of the 
invention but has not been suitable for the application discussed here due 
to its high cost and lack of weldability. 
From the above it may be seen that alloys of the invention can be selected 
and formulated to perform better than the three comparative groups of 
alloys of the prior art. They may be formulated to give better hot 
strengths than the far costlier nickel base or cobalt base alloys. They 
may be formulated to provide much higher as-cast ductility and weldability 
than prior art cast high hot strength alloys. And they may be formulated 
to provide much higher hot strengths along with excellent weldability both 
before and after service exposure than prior art alloys designed with this 
end in mind. 
As various changes could be made in the above described alloy without 
departing from the spirit and scope of the invention, it is intended that 
all matter contained in the above description shall be interpreted as 
illustrative and not in a limiting sense.