CA2099358C - Single crystal nickel-based superalloy - Google Patents

Abstract

This invention relates to a nickel-based superalloy comprising the following elements in percent by weight: from about 5.0 to about 7.0 percent rhenium, from about 1.8 to about 4.0 percent chromium, from about 1.5 to about 9.0 percent cobalt, from about 7.0 to about 10.0 percent tantalum, from about 3.5 to about 7.5 percent tungsten, from about 5.0 to about 7.0 percent aluminum, from about 0.1 to about 1.2 percent titanium, from about 0 to about 0.5 percent columbium, from about 0.25 to about 2.0 percent molybdenum, from about 0 to about 0.15 percent hafnium, and the balance nickel + incidental impurities, the superalloy having a phasial stability number NV3B less than about 2.10.

Description

Expre.j ~k.il RB345060034 ,. ~
20g93~8 SINGLE CRYSTAL NICKEL-BASED SUPERALLOY

BACKGROUND OF THE INVENTION
1. Field of the Invention This invention relates to single crystal nickel-based superalloys and, more particularly, single crystal nickel-based superalloys and articles made therefrom for use in advanced gas turbine engines under high stress, high temperature conditions.

2. Description of the Prior Art Advances over recent years in the metal temperature and stress capability of single crystal articles have been the result of the continuing development of single crystal superalloys, as well as improvements in casting processes and engine application technology. These single crystal superalloy articles includP
rotating and stationary turbine bl~des and vanes found in the hot sections of gas turbine engines. However, gas turbine engine design goals have remained the same during the past decades.
These goals include the desire to increase engine operating te~.perature, rotational speed, thrust-to-weight ratio, fuel efficiency; and enaine component durability and reliability.

The basic technology of alloys for the casting of single crystal components is described in U.S. Patent Nos.

3,494,709; 4,116,723 and 4,209,348. Development work resulted in f ~

first generation nickel-based superalloys, which were materially improved over those described in the aforementioned patents.
However, these first generation nickel-based superalloys contained no rhenium. Examples of such first generation nickel-based superalloys, commercially known as CMSX-2 alloy and CMSX-3 alloy produced by Cannon-Muskegon Corporation, assignee of the present application, are described in U.S. Patent No. 4,582,548.
Further development work resulted in second generatlon nickel-based superalloys having improved creep strength/creep rate.
These second generation nickel-based superalloys have a moderate rhenium content of about 3 weight percent. An example of such a second generation nickel-bzsed superalloy is described in U.S.
Patent No. 4,643,782. Thls patent discloses a superalloy, commercially known as CMSX-4 alloy, having a specific nickel-based composition including a rhenium content in the range of 2.8-3.2 weight percent. The present invention provides the next generation of nickel-based superalloys having higher total refractory element (W+Re+Mo+Ta) content and improved mechanical properties.

Single crystal articles are generally produced having the low-modulus (001) crystallographic orientation parallel to the component dendritic growth pa.~ril or bl-ade s.acking 'aul' axis. Face-centered cubic (FCC) superalloy single crystals grown in the (001) direction provide extremely good jthermal fatigue resistance relative to conventionally cast articles. Since these single crystal articles have no grain boundaries, alloy design 20993'58 ' -without grain boundary strengtheners, such as carbon, boron and zirconium, is possible. As these elements are alloy melting point depressants, their reduction from an alloy design provides a greater potential for high temperature mechanical strength achievement since more complete gamma prime solution and microstructural homogenization can be achieved relative to directionally solidified (DS) columnar grain and conventionally cast materials. Their reduction also makes possible a higher incipient melting temperature.

These process benefits are not necessarily realized unless a multi-faceted alloy design approach is undertaken.
Alloys must be designed to avoid tendency for casting defect formation such as freckles, slivers, spurious grains and recrystallization. Additionally, the alloys must provide an adequate heat treatment window (numeric difference between an alloy's gamma prime solvus and incipient melting point) to allow for nearly complete gamma prime solutioning. At the same time, the alloy compositional baiance should be designed to provide an adequate blend of engineering properties necessary for operation in gas turbine engines. Selected properties generally considered important by gas turbine engine designers include: elevated temperatur2 creep-rupture strength, thermo-mechanical fatigue resistance, impact resistance plus hot corrosion and oxidation resistance.

An alloy designer can attempt to improve one or two of -,1 .: " 2091935~
-these design properties by adjusting the compositional balance of known superalloys. However, it is extremely difficult to improve more than one or two of the design properties without significantly or even severely compromising the remaining properties. The unique superalloy of the present invention provides an excellent blend of the properties necessary for use in producing single crystal articles for operation in gas t-~rbine engine hot sections.

SUMMARY OF THE INVENTION

This invention relates to a nickel-based superalloy comprising the following elements in percent by weight: from about 5.0 to about 7.0 percent rhenium, from about 1.8 to about 4.0 percent chromium, from about 1.5 to about 9.0 percent cobalt, from about 7.0 to about 10.0 percent tantalum, from about 3.5 to about 7.5 percent tungsten, from about 5.0 to about 7.0 percent aluminum, from about 0.1 to about 1.2 percent titanium, from about 0 to about 0.5 percent columbium, from about Q.25 to about ~.0 percent molybdenum, from about O to about 0.15 percent hafnium, and the balance nickel plus incidental impurities, the superalloy having a phasial stability number NV3B less than about 2.10.

Advantageously, this superalloy composition may be further comprised of (percentages are in weight ~ercent) from about O to about 0.04 percent carbon, from about O to about 0.01 209935g percent ~oron, from about O to about 0.01 percent ~ttrium, from about ~ to abou~ 0.01 percent cerium and from about O to about O.01 ~e~cent lanthan~m. A~t~ough incidental i~purities should be kept to the leas~ amount possi~e/ the s~uper~lloy can also be comprised Oc from about O to a~out 0.04 percent manganese, from about O to a~ou~ 9.05 percent silicon, f~om ~bout ~ to about O.01 percen~ zi~conium, from about O ~ about ~.001 percent sulfur, and fxom abou, O .o about 0.10 percent vanadium. In all cases, the base elemen~ is ~ic~el. Furthermore, this superallo~ can advantageously have a phasi~l sta~ility number ~ less than abou~ 1~85, and a ~hromiu~ content of from abo~ 1.8 to abou~ 3~0 percent, a rhenium content of ~rom a~out 5.5 to about 6,5 perc~nt, and a co~alt con~ent of from about ~.0 ~o a~ou~ 5.0 percent. This invention provides a cuperalloy ha~ing an increased resistan~e to creep under high st~ess t high temperature condi~ions, par~icularly up to abou~ 5~F.

Single cr~vstal articles can ~e suita~ly made from the supexalloy of ~his invention. The article can be a component for a turbine engine and, more particularly, the component can be a gas tur~ine blade or gas turbine vane.

The s~peralloy compositions of this inYention ha~e critically ~alanced alloy c~ 7 ~try which resul~s in a unique blen~ o~ desirable properties~ These properties include:
excellent single cx~stal component castabili~y, particularly for moderately sized blade and vane componen~s adequate cast - 209935~

componen~ soluticnability: ex~ellen~ resistance to single crysta~
cast co~ponent recrystallization; ultxa-high creep-rupture stre~g~h to about ~975~F; extremely goo~ low cycle fa~igue s~rength; extx~ely good high cycle fatigue strength: high impact streng~h; very ~oo~ ba~e hot corrosion resistance; very good bare oxi~tion resist~nce, and adequate microstxuctural stability, such as re~istance to the undesix~ble, brittle phases called topologically c}ose-packed (TcP~ pha~es.

Ac~o~dingl~, it is an ob~ect of the prese~ t~ provi~e superalloy compositions and singl~ cryst~l articles mzde the~e~rom having a unique blend of desirable properties. It is a ~urt~er object o~ the presen~ invention to provide supe~lloys ~nd singie crystal articles m~de ~herefrom for use in ~dvanced gas tu~ine engines under hig~ stress, high temperat~re conditions, such as ~p to ab~ ls7s~F. These and other ob~ects and adva~tages of the presen~ invention will ~e apparent to those skilled in the a~t ~p~n reference .o the followi~g description of the preferred embodiments.

BRIEF ~i~;SC~I.PTION OF T}~E I;~WINGS

~ IG. 1 is a cha~t of hot corrosion ~est results performed t~ 117 hours on cne embodiment ~f ~he alloy of this in~ention and on two prior art alloys;

.

20993~8 FIG. 2 is a chart of hot corrosion test results performed to 144 hours on another embodiment of the alloy of this invention and on a prior art alloy.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

The nickel-based superalloy of the present invention comprises the following elements in percent by weight:

Rhenium about 5.0-7.0 Chromium about 1.8-4.0 Cobalt about 1.5-9.0 Tantalum about 7.0-10.0 Tungsten about 3.5-7.5 Aluminum about 5.0-7.0 Titanium about o.l-1.2 Columbium about 0-0.5 Molybdenum about 0.25-2.0 Hafnium about 0~0.15 Nickel + Incidental balance Impurities This superalloy composition also has a phasi~l stabilit~ r.ulm.~er N~B less than about 2.10. Further, this invention has a critically balanced alloy chemistry which results in a unique blend of desirable properties. These properties~include increased creep-rupture strength relative to prior art single , CA 020993~8 1998-06-01 crystal superalloys, single crystal component castability, cast component solutionability, single crystal component resistance to recrystallization, fatigue strength, impact strength, bare hot corrosion resistance, bare oxidation resistance, and microstructural stability, including resistance to TCP phase formation under high stress, high temperature conditions.

Unlike prior nickel-based superalloys known in the art, the superalloys of the present invention have a low chromium, low cobalt and high rhenium content. The chrominum is about 1.8-4.0%
by weight. Advantageously, the chromium content is from 1.8% to 3.0% by weight. This chromium content is significantly lower than that typically found in prior art single crystal nickel-based superalloys. In the present superalloy, chromium provides hot corrosion resistance, although it may also assist with the alloy's oxidation capability. Tantalum and rhenium also assist toward hot corrosion property attainment, and aluminum is present at sufficient levels to provide adequate oxidation resistance, so that relatively low addition of chromium is tolerable in this alloy. Besides lowering the alloy's gamma prime solvus, chromium contributes to the formation of Cr, Re, W-rich TCP phase and must be balanced accordingly in these compositions.

The cobalt content is about 1.5-9.0% by weight.
Advantageously, the cobalt content is from 2.0% to 5.0% by weight. This cobalt content is lower than that typically found in prior art single crystal nickel-based superalloys. In the present superalloy, cobalt assists in providing an appropriate 20.993~-8 heat treatment window since it has the effect of lowering the alloy's gamma prime solvus while generally not affecting its incipient melting point. Rhenium-containing alloys are generally designed with much higher cobalt content than the present invention for the purpose of imparting increased solid solubility and phasial stability. However, the superalloys of the present invention unexpectedly show that much lower cobalt contents are possible and desirable toward providing optimized phasial stability, including control of TCP phase formation.

The rhenium content is about 5.0-7.0% by weight and, advantageously, rhenium is present in an amount of from 5.5% to 6.5% by weight. The amount of rhenium in the superalloy of the present invention is significantly greater than the rhenium content of prior art single crystal nickel-based superalloys.
Furthermore, the superalloys of this invention are generally designed with an increased level of refractory element content, e.g., W+Re+Mo+Ta. The tungsten content is about 3.5-7.5% by weight and, advantageously, the amount of tungsten is from 3.5%
to 6.5% by weight. Tungsten is added since it is an effective solid solution strengthener and it contributes to strengthening the gamma prime. Additionally, tungsten is effective in raising ' h2 alloy's incipient melting temperature. The amount ~f tungsten added to these superalloys is balanced with the amount of rhenium added since they both contribute to.the formation of "freckle" defects during the single crystal inve~tment casting process. They also both strongly effect the propensity for TCP

209~35g phase formation.

Similar to tungsten, rhenium is effective in raising the alloy's incipient melting point. However, rhenium is a more effective strengthener than tungsten, molybdenum and tantalum in terms of elevated temperature creep-rupture and, therefore, rhenium is added appropriately. Additionally, rhenium has a positive influence on this alloy's hot corrosion resistance.
Moreover, rhenium partitions primarily to the gamma matrix, and it is effective in slowing gamma prime particle growth during high temperature, high stress conditions. Besides requiring the balancing of rhenium with tungsten for castability reasons, W+Re must also be set at a level consistent with minimizing TCP phase formation. In general, the TCP phases which occur in such material are rich in chromium, tungsten, and rhenium content, with rhenium being present in the greatest proportion. Thus, careful Re/W ratio control is necessary in this alloy to control the propensity for TCP phase formation.

The molybdenum content is about 0.25-2.0% by weight.
Advantageously, molybdenum is present in an amount of from 0.25%
to 1.5% by weight. Molybdenum is a good solid solution strellg.h~ner, but t is no~ ~s effect~ve as tungsten, rhenium and tantalum. However, since the alloy's density is always a design consideration, and the molybdenum atom is lighter than the other solid solution strengtheners, the addition of moly~denum is a means of assisting control of the overall alloy density in the ~0993~8 compositions of this invention.

The tantalum content is about 7.0-10.0~ by weight and, advantageously, the tantalum content is from 8.0% to 10.0% by weight. Tantalum is a significant contributor to this alloy's strength through means of solid solution strengthening and enhancement of gamma prime particle strength (tantalum also partitions to the gamma prime phase). In this alloy, tantalum is able to be utilized at relatively high concentration since it does not contribute to TCP phase formation. Additionally, tantalum is an attractive single crystal alloy additive in this composition since it assists in preventing "freckle" defect formation during the single crystal casting process. Tantalum is also beneficial in this composition since it tends to raise this alloy's gamma prime solvus, and it is effective toward promoting good alloy oxidation and hot corrosion resistance, along with aluminide coating durability.

The aluminum content is about 5.0-7.0% by weight.
Furthermore, the amount of aluminum present in this composition is advantageously from 5.3% to 6.5% by weight. Aluminum and titanium are the primary elements comprising the gamma prime phase. These elements are added in this composition in a proportion and ratio consistent with achieving adequate alloy castability, solution heat treatability, phasial stability and high mechanical strength. Aluminum is also added to this alloy in proportions sufficient to provide oxidation resistance.

,' ' ~ ,t'~- 2099358 ~

The titanium content is about 0.1-1.2% by weight.
Advantageously, titanium is present in this composition in an amount from 0.2% to 0.8% by weight. Titanium is generally beneficial to the alloy's hot corrosion resistance, but it can have a negative effect to oxidation resistance, alloy castability and alloy response to solution heat treatment. Accordingly, the titanium content must be maintained within the stated range of this composition.

The columbium content is about 0-0.5% by weight and, advantageously, the columbium content is from 0 to 0.3% by weight. Columbium is a gamma prime forming element and it is an effective strengthener in the nickel-based superalloys of this invention. Generally, however, columbium is a detriment to alloy oxidation and hot corrosion properties, so its addition to the composition of this invention is minimized. Moreover, columbium is added to this invention's composition for the purpose of gettering carbon, which can be chemi-sorbed into component surfaces during non-optimized vacuum solution heat treatment procedures. Any carbon pick-up will tend to form columbium carbide instead of titanium or tantalum carbide, thereby preserving the greatest proportion of titanium and/or tantalum for gamma prime and/or solid solution strengthening in this alloy.

The hafnium content is about 0-0.15% bX weight and, advantageously, hafnium is present in an amount from 0.02 to ~-' 20~93S8, 0.05% by weight. Hafnium is added in a small proportion to the present composition in order to assist with coating adherence.
Hafnium generally partitions to the gamma prime phase.

r--~' The balance of this invention's superalloy composition is comprised of nickel and small amounts of incidental impurities. Generally, these incidental impurities are entrained from the industrial process of production, and they should be kept to the least amount possible in the composition so that they do not affect the advantageous aspects of the superalloy. For example, these incidental impurities may include up to about 0.04% by weight manganese, up to about 0.05% by weight silicon, up to about 0.01% by weight zirconium, up to about 0.001% by weight sulfur, and up to about 0.10% by weight vanadium. Amounts of these impurities which exceed the stated amounts could have an adverse effect upon the resulting alloy's properties.

Additionally, the superalloy may optionally contain about 0-0.04% by weight carbon, about 0-0.01% by weight boron, about 0-0.01% by weight yttrium, about 0-0.01% by weight cerium and about 0-0.01% by weight lanthanum.

Not only does the superalloy Qf this invention have a composition within the above specified ranges, but it also has a phasial stability number N~B less than about 2 10.

Advantageously, the phasial stability number N~B is less than 1.85 and, preferably, the phasial stability number N~B is less '~ ~ ' r 20~g3~X ( than 1.75. As can be appreciated by those skilled in the art, NV3B is defined by the PWA N-35 method of nickel-based alloy electron vacancy TCP phase control factor calculation. This calculation is as follows:
EOUATION l ~onversion for weight percent to atomic percent:
Atomic percent of element i = Pi = Ui/Ai X100 ~ i ~Ui /Ai ) where: Wi = weight percent of element i Ai = atomic weight of elemént i Calculation for the amount of each element present in the continuous matrix phase:
Element Atomic amount Rii remaininq Cr RCr 0.97PCr-0.375PB-1.75Pc Ni RNi=pNilo szspa-3~pAllo o3pcr~pTi-o~5pc+o~pv PTa PCb Hf~
Ti, Al, B, Ri=o C, Ta, Cb, Hf V Rv=0.5Pv W R(v) PV-O 167PC Pv PMO + P~J

M~ ~MO) P~MO) ~ ~ 75PB 0.167Pc PM~

Note: weight percentage Re is added to weight percentage W
for the calculation above.

Calculation of Nv3~ using atomic factors from Equations l and 2 above: - ;

Nji = Rl._ then NV3B - ~jNi(NV) 2û~9358 ~_ where: i = each individual element in turn.
Nji = the atomic factor of each element in matrix.

(NV)i = the electron vacancy No. of each respective element.

This calculation~is exemplified in detail in a technical paper entitled "PHACOMP Revisited", by H. J. Murphy, C. T. Sims and A. M. Beltran, published in Volume 1 of International Symposium on Structural Stability in Superalloys (1968) As can be appreciated by those skilled in the art, the phasial stability number for the superalloys of this invention is critical and must be less than the stated maximum to provide a stable microstructure and capability for the desired properties under high temperature, high s'ress conditions. The phasial stability number can be determined empirically, once the practitioner skilled in the art is in possession of the present subject matter.

The superalloy of this invention can be used to suitably make single crystal articles, such as components for turbine engines. Preferably, this superalloy is utilized to make a single crystal casting to be used under high stress, high temperature conditions characterized by an increased resistance to creep under such conditions, particularly high temperature conditions up to about 1975~F. While this superalloy can be used for any purpose requiring high strength castings incorporating a single crystal, its particular use is in the casting of single crystal blades and vanes for gas turbine engines. This alloy ~ 15 ~ 0993~

possesses an unusual resistance to component recrystallization during solution heat treatment, which is considered an important alloy characteristic that is necessary when producing advanced technology, multi-piece, cast bonded single crystal airfoils.
Additionally, this superalloy provides the alloy castability characteristics believed necessary to produce conventional-process-cast, moderately-sized turbine airfoils with intricate cooling passages.

While this superalloy's primary use is in aircraft turbine engines, there are stationary engine applications requiring the specialized high performance characteristics of this alloy. This is particularly the case in turbine engines which require performing characteristics with very restricted clearances, thereby materially limiting the amount of permissible creep. Engines designed to devélop high performance characteristics are normally operated at higher component temperatures and, therefore, the problem of creep is increased.
Generally, creep in excess of 1% is considered unacceptable in these cases. The creep characteristics of known state of the art alloys have limited operating temperatures and, thus, maximum performance capability. The superalloy of this invention has an increased resistance to creep under high stress, high temperature conditions, particularly up to 1975~F.

The single crystal components made from this invention's compositions can be produced by any of the single . CA 020993~8 1998-06-01 crystal casting techniques known in the art. For example, single crystal directional solidification processes can be utilized, such as the seed crystal process and the choke process.

The single crystal castings made from the superalloy of the present invention are advantageously subjected to a high temperature aging heat treatment in order to optimize the creep-rupture properties of these alloys. This invention's single crystal castings can be aged at a temperature of from about 1950~F to about 2125~F for about 1 to about 20 hours. However, as can be appreciated by those skilled in the art, the optimum aging temperature and time for aging depends on the precise composition of the superalloy.

This invention provides superalloy compositions having a unique blend of desirable properties. These properties include:
excellent single crystal component castability, particularly for moderately sized blade and vane components; excellent case component solutionability; excellent resistance to single crystal cast component recrystallization; ultra-high creep-rupture strength to about 1975~F; extremely good low cycle fatigue strength; extremely good high cycle fatigue strength; high impact strength; very good bare hot corrosion resistance; very good bare oxidation resistance; and microstructural stability, such as resistance to formation of the undesirable TCP phases. As noted above, this superalloy has a precise composition with only small permissible variations in any one element if the unique blend of properties is to be maintained.

2~g93~8 !

, In order to more clearly illustrate this invention and provide a comparison with representative superalloys outside the claimed scope of the invention, the examples set forth below are presented. The following examples are included as being illustrations of the invention and its relation to other superalloys and articles, and should not be construed as limiting the scope thereof.

. --EXAMPLES

A large number of superalloy test materials were prepared toinvestigate the compositional variations and ranges for the superalloys of the present invention. Some of the alloy compositions tested and reported below fall outside the claimed scope of the present invention, but are included for comparative purposes to assist in the understanding of the invention.
Representative alloy aim chemistries of those materials tested are reported in Table 1 below.

See "Key" Be l o~
Al loy C B Cr Co ~o ~I Cb Ti Al Ta Re H~ Ili llv3B- 1 2 3 4 aSX-10~ ~ - 3.0 8.5 .70 7.2 .30 .65 6.0 7.6 5.0 .05 8AL 2.08 12.46 6.65 14.5j 20.76 -lOB - - 2.6 8.2 .70 6.ff .30 .68 6.0 7.9 4.~5 .06 BAL 2.02 11.9 6.68 14.88 20.5 -10C - - 2.5 7.7 .70 6.6 .30 .65 5.9 8.2 4.8 .05 BAL 1.90 11.4 6.55 15.05 20.3 -100 - - 4.0 4.8 .60 6.4 .30 .60 5.7 8.2 4.9 .03 8AL 1.95 11.3 6.30 14.80 20.1 -lOE - - 2.2 7.2 .70 6.3 .25 .n 5.85 8.3 4.8 .042 CAL 1.84 11.1 6.57 15.12 20.1 -10F.02 .02 2.4 7.6 .65 6.45 .28 .63 5.9 8.5 5.0 .046 SAL 1.89 11.45 6.53 15.31 20.6 -10G - - 2.4 6.3 .50 6.4 .20 .70 5.8 8.0 5.5 .04 BAL 1.82 11.9 6.5 14.7 20.4 -10Ga - 2.4 4.0 .50 6.2 .15 .55 5.8 8.3 5.6 .04 BAL 1.72 11.8 6.35 14.8 20.6 10K-10Gb - - 2.3 3.3 .40 5.5 .10 .30 5.7 8.4 6.3 .03 BAL 1.60 11.8 6.0 14.5 20.6 -10H - - 2.2 5.9 .50 6.4 .15 .80 5.9 8.0 5.5 .04 BAL 1.82 11.9 6.7 14.85 20.4 -101 - - 2.5 4.7 .50 6.4 .15 .70 5.8 7.9 6.0 .04 BAL 1.81 12.4 6.5 14.65 20.9 -1OIa - - 2.5 3.3 .40 6.1 .10 .60 5.8 7.9 6.0 .04 BAL 1.69 12.1 6.4 14.4 20.4 -10J.015 .01 2.65 4.0 .50 6.0 .20 .65 5.8 9.0 5.5 .04 BAL 1.79 11.5 6.45 15.65 21.0 10L - - 2.0 2.7 .40 5.3 .10 .20 5.65 8.4 6.3 .03 BAL 1.50 11.6 5.85 14.35 20.4 aSX-12A - - 3.0 4.5 .35 5.S - 1.0 5.65 9.0 5.5 .04 BAL 1.84 11.0 6.65 15.65 20.35 -12B - - 3.5 3.0 .35 5.0 - .90 5.60 8.8 6.0 .04 B~L 1.80 11.0 6.5 15.3 20.15 -12C - - 2.8 3.5 .40 5.3 - .75 5.60 8.8 5.8 .04 BAL 1.70 11.1 6.35 15.15 20.3 12D-12Ca - - 2.5 3.2 .45 4.7 - .50 5.60 8.7 6.3 .03 B~L 1.61 11.0 6.10 14.8 20.15 -12E - - 2.0 3.0 .45 4.7 - .40 5.60 8.7 6.3 .03 BAL 1.50 11 0 6.0 14.7 20.15 aSX-10Ri - - 2.65 7.0 .60 6.4 .40 .80 5.8 7.5 5.5 .06 BAL 1.91 11.9 6.6 14.5 20.0 aSX-12Ri - - 3.4 8.0 .50 6.0 - 1.0 5.6 7.6 5.3 .06 BAL 1.92 11.3 6.6 14.6 19.4 I~eY~
2 ~ T i 3 - ~ l ~ T i ~ Ta ~ Cb 4 - ~1 ~ Re I l~o I Ta ~ Calculated usir~ A ~1-35 ~lethod .
,~-'- 20993~

Third generation single crystal alloy development to investigate the compositional variations for the superalloys of the present invention began with the definition and evaluation of a series of experimental compositions. Increased creep-rupture strength was the primary objective of the initial development effort, with elemental balancing to provide a combination of useful engineering characteristics following the definition of a base concept for increased strength.

T~.e initial materials explored the utility of higher levels of refractory element and gamma prime forming elements than are present in similar prior art compositions. As shown in Table 1, the alloy chromium content was reduced to improve alloy stability. Cobalt content, initially thought to be required for increased solid solubility, could be significantly reduced.
Refractory element content (W+Re+Mo+Ta) was varied, while the summation of the primary gamma prime partitioning elements (Al+Ti+Ta+Cb) was also varied. The alloy's Re content was initially explored at conventional levels, but it was found that the Re level had to be increased.

Standard Nv3~ calculations were performed during the initial alloy design stage to assist respective alloy phasial stability predictions, with that number varying 'rom one alloy composition to another.

CA 020993~8 1998-06-01 Some of the alloys were produced using production-type procedures. These alloys were vacuum induction melted in the Cannon-Muskegon Corporation V-1 furnace, yielding approximately 200-300 lbs. of bar product per alloy (see Table 2 below).
Quantities of each compositional iteration, as reported in Table 2, were made into test bars and test blades by vacuum investment casting. Solution heat treatment procedures were developed in the laboratory in 3" and 6" diameter tube furnaces. Gamma prime aging treatments were also performed in the laboratory.

Heat Alloy No. C B Cr Co Mo W Cb TiAl Ta Re Hf Ni CMSX-lOA VF 778 .001 <.0012.9 8.5 .77.2 .3 .70 6.05 7.6 5.0 .05 BASE
-lOB VF 831 .002 c.OOl2.6 8.2 .76.9 .3 .68 6.06 7 9 4-9 05 BASE
-lORi VF 965 .001 c .0012.65 7. O .66.4 .4 .80 5.72 7.6 5.5 .06 BASE
-lORi VF 966 .001 c.OOl2.69 7.0 .66.3 .4 .80 5.66 7.6 5.4 .06 -lORi VF 980 .001 c.OOl2.66 7.0 .66.3 .4 .79 5.78 7.6 5.4 06 BASE
-12Ri VF 963 .001 <.0013.3 8.0 .48 6.0 ~.05 1.01 5.69 7.6 5.3 ,07 BASE
-12Ri VF 964 .001 c.OOl3.4 8.0 .48 6.1 c.05 1.00 5.60 7.6 5.3 .06 BASE
-12Ri VF 979 .001 c.OOl3.4 8.0 .50 6.1 c O5 1.00 5.56 7-6 5.3 ~06 BASE
-lOGa VF 983 .001 c.0012.4 3.95 .41 6.1 .14 .56 5.83 8 4 5 9 -03 BASE
-12C VF 985 .001 c .0012.7 3.5 .45 5.3 c .05 .75 5.66 8.8 6. O .025 -lOGb (-lOK) VF 994 .001<.001 2.2 3.3 .40 5.5 .09 .24 5.74 8.2 6.4 .025 BASE
-12Ca (-12D) VF 993 .001<.001 2.4 3.2 .46 4.8 <.01 .50 5.64 8.6 6.4 .025 BASE

20993S8 ~

All other specimens reported in Table 1 above were produc~d by blending base alloy bar stoc~ with the virgin elemental additions necessary to achieve the desired composition. The blending was done during test bar and blade manufacture. The base alloy bar stoc~ plus virgin additions were placed into the casting furnace melt crucible, mel~ed and the bath homogenized prior to pour~ng into an appropriate shell mold. It is believed that good correiation between alloy aim chemistry and test bar/blade chemistry was routinely achieved (see Table 3 below).

AlloY C 8 Cr CoMo V Cb Ti Al Ta Re Hf Ni llv38 CMSX-10A - - 2.9 8.5.687.4 .29 .69 6.0 7.5 5.1 .07 aAL 2.09 -108 - - 2.7 8.1.696.95.29 .69 6.0 7.8 4.8 .06 BAL 2.01 -10C - - 2.6 7.7.696.4 .30 .62 5~7 8.3 4.7 .07 BAL 1.86 -10G - - 4.0 5Ø626.0 .31 .59 5.44 8.1 4.7 .04 8AL 1.83 -10~ - - 2.2 7.2.706.4 .26 .63 5.89 8.2 4.8 .05 8AL 1.84 -10F .014 .027 2.4 7.7.65 6.4 .28 .63 5.96 7.9 5.0 .04 BAL 1.86 -10G - - 2.5 6.5.535.5 .20 .68 5.6 8.2 4.6 .05 BAL 1.68 -10Ga - - 2.4 4Ø416.2 .14 .55 5.79 8.3 6.0 .025 aAL 1.73 -10Gb ~10K~ - - 2.3 3.5.42 5.9 .10 .43 5.67 8.5 6.0 .OZ4 BAL 1.63 -10H - - 2.3 5.6.516.2 .17 .76 5.58 7.8 5.4 .05 8AL 1.69 -101 ~ ~ 2.6 4.8.526.6 .14 .67 5.65 7.4 5.4 .04 3AL 1.70 -101a - - 2.7 3.5.475.2 .10 .60 5.80 8.0 5.8 .04 8AL 1.67 -10J .017 .01 2.6 4Ø4a6.0 .19 .62 5.74 8.8 5.7 .04 8AL 1.76 -12A - - 3.0 4.6.395.3<.01 .96 5.61 9.4 5.0 .05 8AL 1.aG
-128 - - 3.5 3Ø385.1<.01 .84 5.52 8.8 6.1 .05 BAL 1.79 -12C - - 2.7 3.5.455.4<.01 .75 5.62 8.8 6.0 .04 BAL 1.72 -12Ca ~12D) - - 2.5 3.2.46 5.0 <.01 .61 5.56 8.7 6.0 .03 BAL 1.60 -tZE
-10Ri - - 2.65 7Ø606.4 .40 .80 5.67 7.6 5.5 .065 3AL 1.87 -12Ri - - 3.4 8Ø486.1<.01 .99 5.54 7.6 5.3 .07 BAL 1.92 P\JA N-35 Method 2'099358( For the CMSX-lOD specimen (see Table 1), high quality virgin elemental additions were vacuum melted and the refined material was poured into 2" diameter bars. In turn, a quantity of the resulting bar was used to produce single crystal test bar/blade specimens by investment casting.

It was apparent that considerable variation in the investment casting process integrity may have occurred during specimen manufacture since varying levels of test bar freckle formation, secondary dendrite arm spacing and propérty attainment were apparent. Derivative alloy response to solution treatment (reported in Table 4 below) varied, and was a function of both alloy composition and test specimen quality.

Heat treatments developed for the alloy iterations are reported in Table 4 below. Full gamma prime solutioning was desired for each material, however, this objective was not universally achieved. Primary gamma prime aging was performed to effect a more desirable gamma prime particle size and distribution.
Secondary gamma prime aging was performed to effect precipitation of conventional matrix gaD a prime precipitates along with ultra-fine gamma prime precipitates located within the matrix channels between the primary gamma prime particles for these specimens.

CA 020993~8 1998-06-01 TI~BLE 4 Heat Treatment Detail Peak Solution ~Y PrimarYy~
AlloY Tem~erature Solutioned* Aqinq+ SecondaryY Aqinq+
~F ~C
CMSX-lOA2460 134997.0-98.0 1975~F/4 Hrs 1600~F/20 + 1400~F/24 -lOB2465 135297.0-98.0 1975~F/4 Hrs 1600~F/20 + 1400~F/24 1975~F/19.5 Hrs -lOC2470 135499.0-99.5 2100~F/8 Hrs 1600~F/20 + 1400~F/24 1975~F/10 Hrs -lOD2450 1343 99.9-100 1975~F/10 Hrs 1600~F/22 + 1400~F/24 -lOE2465 1352 100 1975~F/15 Hrs 1600~F/20 + 1400~F/24 1975~F/21 Hrs 1600~F/25.5 + 1400~F/23 -lOF2444 1370 95 1975~F/16 Hrs 1600~F/23 + 1400~F/24 -lOG2475 135799.0-99.5 1975~F/12 Hrs 1600~F~24.5 + 1400~F/17 -lOGa2485-901363-6599.5-100 2075~F/5 Hrs 1600~F/ZO + 1400~F/23 2075~F/6 Hrs 1600~F/24 + 1400~F/30 1612~F/48 + 1414~F/22 -lOGb(lOK) 2485 1363 100 2075~F/6 Hrs 1600~F/24 + 1400~F/30 -lOH2475 135798.5-99.0 1975~F/16 Hrs 1600~F/27.5 + 1400~F/27 1975~F/18 Hrs 1600~F/101 + 1400~F/46 -lOI2475 1357 100 2075~F/5 Hrs 1600~F/22 + 1405~F/24 -lOIa2480 1360 99.5-100 2075~F/5 Hrs 1600~F/24 + 1400~F/24 -lOJ2480 136098.0-99.0 1975~F/15 Hrs 1600~F~24 + 1400~F~30 2075~F/5 Hrs 1600~F/24 + 1400~F/30 -lOL
-12A2475 135798.5-99.0 1975~F/16.5 Hrs 1600~F/24 + 1400~F/32 1975~F/12 Hrs 1600~F/24 + 1400~F/27.5 -12B2480 136099.0-99.5 1975~F/13 Hrs 1600~F/57 + 1400~F/39 -12C2485-901363-6599.5-100 2075~F/5 Hrs 1600~F/20 + 1400~F/5 2075~F/6 Hrs 1600~F/24 + 1400~F/30 -12Ca(12D)2485 1363 100 2075~F/6 Hrs 1600~F/24 + 1400~F/30 -lORi2460 134998.5-99.8 2075~F/6 Hrs 1600~F/24 + 1400~F/30 -12Ri2455 1346 100 2075~F/6 Hrs 1600~F/24 + 1400~F/30 ~ Determined by visual estimation + Specimens air cooled from all aging treatments Fully heat treated test bars were creep-rupture tested. The specimens were machined and low-stress ground to ASTM standard proportional specimen dimension. The specimens were creep-rupture tested at various condition of temperature and stress, according to standard ASTM procedure.

- ! 2099358 A significant ~actor o~ the CMS~-lOA a~loy design was the shift ~o higher Re con~ent. At the same ti~e, W, Cr, Ta and other ga~a prime strengtheners were balanced to provide the desired allo~ charac~eris~ics and pr~perties. The allo~s hîgher Re le~el resulted in significantl~ improved creep-rupt~re strength th~oughout t~e entire test regi~e, as indicated by the results reported in Ta~le S below for ~he CMSX-lOA specimens.

f 20993S8 TI~E IN ROURS
CUPTURE TI~E X X ~I~AL C~ctr ~~ADI~G TO ~~ACN
TEST CO~DITIO~ ~OURS ELO~G. ~A t, hours S J~r~, tion 1.0S 2.0S
1600-F/75.0 ksi534.4 24.2 26.9 534.2 Z.331 10.9 21.0 328.4 22.0 27.8 328.3 21.055 6.3 8.7 527.3 21.1 26.3 526.3 17.552 28.4 72.2 1700-F~50.0 ksi305.0 31.1 34.5 304.2 28.614 62.1 108.9 292.4 19.2 19.9 291.6 19.324 71.5 123.7 87.6 2.6 5.8 85.7 1.474 65.9 -1800-FnO.0 ksi 415.6 16.1 21.4 413.8 15.643 182.7 246.1 848.0 37.1 33.0 846.3 34.326 460.4 524.3 1016.2 33.2 30.5 1014.3 32.984 476.8 655.1 1800-F/36.0 ksi586.5 38.1 38.0 585.6 33.050 -3ff.0 425.0 572.7 36.9 35.3 570.7 29.029 395.0 4Z.0 546.5 26.4 34.2 545.7 ' 25.843 373.0 406.0 420.3 22.4 26.3 418.7 18.105 286.7 317.6 426.0 14.8 17.0 425.1 10.244 326.5 353.2 239.8 24.3 23.8 239.7 23.264 94.1 123.9 255.7 19.9 27.4 253.6 18.510 115.2 152.7 1gOO-F/25.0 ksi32.3 5.5 11.0 31.0 2.075 26.7 30.7 129.7 43.2 38.9 128.7 39.556 30.4 48.~
168.7 34.7 36.4 166.1 30.816 58.2 78.4 ~8.1 S8.1 32.3 226.4 16.926 146.3 160.6 277.7 29.5 31.1 276.4 27.323 9.9 29.9 ~23.4 39.7 38.3 422.7 35.121 218.4 250.9 383.8 35.9 36.1 382.7 34.861 192.9 226.7 373.3 31.3 35.7 371.6 26.138 211.6 238.0 2000-F/18.0 ksi138.0 22.3 33.0 136.3 19.052 33.9 77.0 134.9 40.7 36.5 134.7 38.328 54.7 71.9 122.9 23.2 34.9 122.0 19.050 50.1 69.4 115.6 34.2 36.6 114.4 30.861 40.8 56.8 245.2 35.1 36.2 244.3 29.844 135.7 157.9 Z1.9 36.3 35.4 221.8 33.737 113.0 140.0 181.2 32.1 34.2 180.1 29.249 53.1 61.4 2050-F~15.0 ksi126.4 47.9 49.0 124.1 30.086 45.8 69.8 150.5 45.5 47.8 148.1 39.308 16.8 34.5 140.5 30.6 40.0 138.7 23.596 30.6 76.4 120.8 29.5 39.7 120.0 29.479 16.3 55.6 79.0 11.7 14.4 79.0 11.644 41.7 54.8 112.2 24.3 31.3 112.1 21.401 55.9 69.5 2100~F/12.5 ksi94.1 22.1 27.5 94.1 20.520 42.2 62.6 112.5 39.4 33.1 112.2 29.126 28.0 58.8 96.6 25.9 35.4 ff.5 14.542 52.3 62.5 123.6 43.4 40.4 122.9 31.050 40.9 63.5 50.8 21.7 29.6 49.9 9.330 35.1 37.6 90.5 41.6 43.7 89.7 37.422 13.6 - 38.5 1800-F/36.0 ksi ~ 420.6 23.9 35.1 419.9 23.196 213.8 286.0 396.1 37.1 34.0 394.7 31.623 239.4 264.9 384.9 31.1 34.0 382.9 25.554 220.5 247.9 As-Solutioned Condition ~ .. . .

,-- 2099358 Microstructural review of the failed rupture specimens of this alloy revealed that TCP phase precipitation occurred during the respective creep-rupture tests, particularly those at 1900~F and a~ove. It became apparent that the Nv3~ phasial stability num~er calculation would be an effective tool in predicting alloy stability and, effectively, high temperature creep strength for the invention.

Wherein the CMSX-lOA specimen's NV3B number was 2.08, CMSX-lOB was designed to the 2.02 level. This was accomplished by the further reduction of alloy Cr content and similar reduction to Co and W+Re level. W was reduced more than the Re in this specimen since Re is more effective in the solid solution. Additionally, wherein some loss in W contribution to the gamma prime could be anticipated, it was sufficiently replaced by the modest increase tc Ta content in this composition. These changes resulted in the CMSX-lOB alloy specimen exhibiting even more improved creep strength at 1800~F. Table 6 reported below illustrates that three specimens achieved an average life of 961 hours, with 1.0%
creep occurring at an average of 724 hours. However, it was observed that TCP phase was present at higher temperature.

t .. : . , ....

20~9358 015X-lOG c~ RE
TIIIE Ill Ill~tS
~L~E TUIE X X FIIUL L~~~ ~DIIIG TO ~EACN
TEST C~lDI~Ial llaJ~S ELCIIG, ~A t, han~sX ' fw L ion 1.OX 2.0X

1800~F/36.0 ksi 907.1 19.2 34.0907.0 17.832697.2 752.7 989.3 18.9 33.5 988.517.657 768.11~17.8 988.4 35.9 36.1 987.331.813 705.8767.5 507.0 44.1 45.4 505.741.5U 317.9352.6 598.1 46.9 43.4 596.142.340 386.5415.2 408.3 62.6 52.1 407.254.479 187.3256.5 265.3 39.7 43.7 262.737.102 87.6 119.2 385.3 45.5 46.2 383.539.031 177.4213.4 412.8 43.4 40.5 410.638.771 189.1233.4 389.3 51.5 44.2 386.836.920 ~O.5 249.2 459.5 40.0 46.3 458.039.513 210.2291.1 258.0 38.1 40.6 257.936.743 32.1 90.2 U4.1 27.9 40.0 U3.426.296 288.1326.
376.9 16.4 20.4 376.8' 16.088 96.0 226.6 U1.0 50.5 48.2 478.834.557 264.4297.5 461.5 35.1 40.6 460.130.786 181.1265.3 U3.0 47.1 46.8 482.143.714 286.2320.7 500.1 33.4 37.0 499.730.486 11.9 280.1 U00-F/40 ksi 436.7 4G.2 44.1436.2 39.818294.6 318.9 390.8 50.1 42.8 390.3~.1.817 250.9276.2 336.9 52.7 U.1 335.246.697 226.5240.9 190C-F/25.0 ksi 237.8 55.9 45.7237.4 53.85433.0 113.5 295.7 57.4 49.1 295.646.592 123.7170.9 2000-F/18.0 ksi 192.7 31.5 26.6191.6 27.73356.3 88.6 166.5 41.4 25.3 166.534.102 46.2 72.7 173.3 36.6 ~7.0 171.431.U1 24.0 66.1 2050~F/15.0 ksi 219.6 40.1 40.4218.6 37.87113.2 56.8 122.3 Z8.2 47.9 120.626.614 37.0 63.7 118.4 33.2 60.0 116.929.986 36.7 56.5 179.7 44.1 U.1 179.139.188 8.4 75.3 74.9 44.2 U.6 74.634.800 6.8 14.5 168.3 U.6 49.7 167.043.171 36.9 77.1 104.8 17.0 27.2 102.81.62~ 66.1 155.9 46.3 49.8 155.238.388 64.4 81.9 90.6 15.1 21.4 87.11.046 75.5 120.5 46.3 55.8 118.735.143 10.3 27.7 150.7 39.8 49.7 150.133.903 21.4 60.9 149.5 33.2 46.2 lU.923.166 73.3 88.3 142.9 42.0 47.5 142.541.524 54.9 70.5 2050~F/15.0 ksi 163.0 52.5 49.2161.9 46.14620.5 76.9 151.1 66.4 45.6 150.759.115 52.7 75.5 131.8 57.3 44.4 131.548.310 26.3 57.1 156.0 54.4 41.0 155.945.502 55.5 78.3 133.7 57.2 56.0 132.741.753 67.5 W.7 135.1 59.7 52.3 134.346.317 54.9 71.5 151.1 66.4 45.6 150.759.115 52.7 75.5 131.8 57.3 44.4 131.5U.310 26.3 57.1 2100- FJ15.0 ksi 69.7 54.2 U. l69.4 47.67425.3 36.3 ~s-Solutioned Condition CA 020993~8 1998-06-01 only about 97-98~ gamma prime solutioning was achieved in the CMSX-lOA and -lOB materials (see Table 4) which was insufficient for the purpose of optimizing alloy mechanical properties and microstructural homogeneity. Attainment of a greater level of gamma prime solutioning, therefore, became an equal priority in tandem with improving microstructural stability at temperatures above 1900~F.

To conform the suspected composition of the TCP phase forming in the alloys, scanning electron microscope (SEM) wavelength dispersive x-ray (WDX) microchemistry analyses of CMSX-lOB test bar contained needles was undertaken and compared to the alloys gamma and gamma prime compositions. The results, reported in Table 7 below, confirm that the needles were enriched in Cr, W
and Re.

20993~8 f-CUSX-10B ~icro-ChemistrY Analyses - Cast Test Bar (VF 831) - T~ r~ Se~tion. ~ott x Bar Location.
- Solutioned to ~465~F
- Aged 1975~F/19.5 Hrs./AC
160û F/20 Hrs./AC
1400~F/24 Hrs./AC
GAMMA PHASE GAMMA PRIME PHASE NEEDLE CO~STITUENT
ELEM K Z A F ELEM K Z A F ELEM K Z A F
ALK 0.0101 1.090 0.324 1.000 ALK 0.0145 1.084 0.322 1.000 ALK 0.0116 1.107 0.347 1.000 TIK 0.0069 1.007 0.930 1.051 TIK 0.0084 1.002 0.934 1.052 TIK 0.0077 1.026 0.908 1.039 CRK 0.0428 1.008 0.963 1.108 CRK 0.0250 1.002 0.965 1.117 CRK 0.0390 1.028 0.949 1.083 COK 0.0970 0.994 0.984 1.018 COK 0.0761 0.988 0.987 1.022 COK 0.0755 1.016 0.977 1.025 NIK 0.6891 1.033 0.988 1.010 NIK 0.7270 1.026 0.991 1.005 ~IK 0.6143 1.056 0.9O3 1.024 TAL 0.0485 0.794 1.020 1.000 TAL 0.0697 0.788 1.024 ~.000 TAL 0.0389 0.814 1.018 1.000 U L 0.0329 0.788 0.963 1.000 U L 0.0311 0.783 0.962 1.000 U L 0.0682 0.808 0.968 1.000 REL 0.04Z2 0.785 0.968 1.000 REL 0.0085 0.779 0.968 1.000 REL 0.1083 0.805 0.973 1.000 UT Z UT X ~T X
ELEM CPS ELEM ELEM CPS ELEMELEM CPS ELEM
AL K 12.1800 2.87 AL K 17.9400 4.19AL K 11.9900 3.02 TI K 5.5200 0.71 TI K 6.8400 0.86TI K 5.2500 0. ~
CR K 27.6400 3.98 CR K 16.4500 2.31CR K 21.5ôO0 3.69 CO K 40.6800 9.74 CO K 32.5400 7.64CO K 27.1700 7.42 NI K 253.1300 66.K NI K 272.3800 71.11 NI K 193.7500 57.84 TA L 6.5667 5.99 TA L 9.6329 8.64TA L 4.5259 4.70 L 4.077~ 4-33 ~ L3-9375 4.13 ~ L7.2620 8.71 RE L 4.6000 5 56 RE L 0.9500 1 13RE L 10.1300 13 82 TOTAL ;00.00 TOTAL ' 100.00 TOTAL 100.00 The calculated Nv38 numbers were 1.90 for CMSX-lOC and 1.95 for CMSX-lOD. Re was maintained at around 5% while W was further reduced to improve stability in these specimens. Alloy Ta was increased since it did not participate in TCP formation and the Ta/W ratio was effectively improved, which assisted with alloy castability. Chromium was reduced in the -lOC specimens but increased to 4.0% in the -lOD specimens to provide an opportunity to determine the suitability of the Cr levels from a hot corrosion standpoint. Co was reduced in both materials, significantly in the -lOD specimen, while Al+Ti level was also CA 020993~8 1998-06-01 reduced to assist in achieving more complete gamma prime solutioning. Creep-rupture results for the two specimens are reported below in Tables 8 and 9, respectively. Even though the -lOD alloy specimens were observed to exhibit full gamma prime solutioning (as opposed to 99.-99.5~ for CMSX-iOC) the alloys greater Cr content, which necessitated a lower Al+Ti level, effected lower properties than attained with CMSX-lOC. However, both materials exhibited improved alloy stability and higher temperature properties, so that attempts to balance the alloys low and high temperature creep response were favorable.

CMSX-lOC CREEP-RUPTURE

R~ ~ FINAL CREEP READING TIME IN HOURS
TIME % % % TO REACH
TEST CONDITION HOURS ELONG. RA t, hours deformation 1.0% 2.0%
1800~F/36.0 ksi 556.1 31.4 30.5 555.226.615316.1 376.3 636.6 43.9 37.5 636.4 38.460416.6455.4 609.2 23.3 34.7 607.6 19.074410.6460.6 635.7 44.9 45.6 635.3 34.991407.3443.4 612.8 43.5 38.8 611.9 41.951409.8438.7 1850~F/36.0 ksi 252.2 30.2 37.8 252.022.03361.1 166.3 298.1 41.3 39.0 297.6 37.953170.3194.8 231.1 33.6 39.5 230.2 29.689127.8146.0 1922~F/20.3 ksi 492.4 52.5 52.4 491.648.922176.5 251.7 529.8 38.6 45.5 528.9 33.353269.6306.2 637.5 48.9 43.3 635.2 45.804189.5318.3 2000~F/18.0 ksi 258.8 35.0 41.5 258.732.44474.2 127.5 293.1 49.2 44.1 292.1 42.079145.6170.9 221.9 43.0 48.5 220.9 33.50755.6123.3 266.1 35.1 44.0 264.6 33.759113.6143.6 2050~F/15.0 ksi 196.6 39.7 40.3 194.127.75526.0 134.8 170.4 30.1 46.3 169.2 25.62411.151.4 193.2 38.1 42.9 191.9 32.28846.576.5 247.3 33.1 40.5 246.0 26.494122.0150.8 C~SX-lOD ~ URE
11 ~S
PUPTU~E TI~E X X FI~AL CR~~P ~E~DIHC TO ~EACHTEST 0DlTla~l ~SE~ G~ ~A t, haurg S dLrv~ ~ian 1.0X z.m 1800-F/36.0 ksi U8.0 26.7 29.3 426.3 24.166 189.2 2U.3 1850-F/36.0 ksi 141.0 23.1 26.8 140.1 20.660 57.8 79.7 140.7 14.7 26.1 140.2 13.741 56.2 77.6 166.0 17.5 28.9 165.0 15.640 76.5 100.1 1922-~t20.3 ksi 519.9 23.8 24.9 518.9 2Z.64B 202.0 345.6 667.0 17.6 23.7 665.2 16.819 151.8 391.4 680.3 14.9 28.2 678.9 14.476 340.2 500.3 2000-F/18.0 ksi 370.3 18.8 21.3 369.9 15.560 20.9 106.9 401.5 11.1 18.0 ~00.0 8.903 19.8 125.5 366.6 17.5 25.8 366.6 ~8.049 223.9 306.1 2050-F/15.0 ksi 465.3 12.9 20.5 465.2 12.639 61.0 305.9 338.8 9.8 24.8 337.7 9.468 30.8 204.4 The acceptability of the alloys' low Cr content was confirmed through extremely aggressive short-term burner rig hot corrosion tests performed at 1650-F, 1% sulfur, 10 ppm sea salt condition.
FIGS. 1 and 2 illustrate the results for tests performed to 117 and 144 hours for the CMSX-lOC and CMSX-lOD specimens, respectively. In both cases, the materials performed similar to MAR M 247-type materials, thereby confirming the suitability of the low Cr alloy design concept.

With the above-noted results, another series of alloys, CMSX-lOE, -lOF, -lOG, -lOH, -lOI,and -12A were designed, produced and evaluated. The alloys explored Re level ranging 4.8-6.3%, 2.2-3.0% Cr level, 4.7-7.6% Co level and the remainder balanced to maintain castability, improve solutionability and improve phasial stability. The NV38 number ranged between 1.81-1.89.

One of the series, CMSX-lOF, contained .02% C and .02~ B. These additions were observed to improve casting yield and may have assisted in providing more consistent control of single crystal cast article orientation. However, the melting point depressants, C and B, restricted the specimen's response to solution heat treatment. The CMSX-lOF creep-rupture properties are reported in Table lO below.

C~SX- 1 OF ~c~r ~J~ I U~_ TI~E IX ~OURS
eUPTU~E TI~E S X FI~UL CREEP DF-nl~G TO RE~L~
TEST CDHDITTO~ HOUQS ELO~G. R~ t, hours S dLfo. Iion 1.0S 2.0X
1840~F/36.0 ksi 616.0 18.1 22.4 615.816.898 439.9 477.6 666.6 45.6 48.0 666.4~3.261 464.6 49Z.3 603.1 25.3 24.3 602.524.281 398.4 444.0 1850-Ft36.0 ksi 243.9 19.6 28.2 243.018.045 129.1 160.9 285.9 26.8 32.1 285.525.701 187.8 206.0 258.6 19.2 29.1 258.318.175 168.3 189.5 1922~F/20.3 ksi 499.5 40.0 41.0 498.537.756 208.2 2n.6 649.2 55.6 52.9 648.351.045 197.6 338.8 361.0 15.8 21.9 357.7 2.599 273.2 335.7 2000'F/18.0 ~si235.4 39.6 51.7 235.437.881 100.8 133.2 276.1 43.7 52.8 274.436.762 115.1 155.9 290.0 36.7 47.3 289.133.304 125.3 162.1 2~50'F/I~.0 ~si255.4 Z8.7 36.6 255.027.426 67.4 131.0 255.1 33.4 43.1 254.931.378 46.2 102.2 254.5 25.4 33.3 254.423.737 50.9 118.7 . CA 020993~8 l998-06-0l The CMSX-lOE, G, H, and I, plus CMSX-12A creep-rupture specimen results are reported below in Tables 11, 12, 13, 14, and 15, respectively. The results show a general improvement to alloy creep-rupture strength above 1900~F while maintaining extremely good strength at lower températures.

CMSX-lOE CREEP-RUPTURE

RU~ 1 UK~ FINAL CREEP READING TIME IN HOURS

TIME % % % TO REACH

TEST CONDITION HOURS ELONG. RA t, hours deformation 1.0% 2.0%

1800~F/36.0 ksi 664.5 31.4 36.3 663.530.435 436.5 470.8 604.4 35.1 36.7 603.3 33.371253.7 355.9 582.5 41.5 36.1 581.7 39.79278.9 329.3 553.5 35.9 37.0 552.5 33.172326.4 357.1 1850~F/36.0 ksi 257.9 25.3 32.0 257.022.734 149.4 170.3 199.2 18.4 32.1 198.6 16.261122.4 139.4 260.5 33.6 33.4 259.7 31.315159.9 174.0 1922~F/20.3 ksi 810.6 38.6 33.0 808.433.523 210.2 378.2 800.9 35.3 36.4 799.1 32.405339.7 434.2 859.9 39.0 35.4 859.6 37.036364.6 465.2 2000~F/18.0 ksi 362.8 27.7 29.3 362.424.887 98.4 177.3 411.2 29.4 27.0 409.9 26.426173.6 218.6 369.7 15.3 28.2 368.8 12.941170.3 221.9 379.7 26.4 26.1 379.2 27.656177.9 206.6 2050~F/15.0 ksi 476.9 21.8 23.4 476.318.233 196.6 255.9 418.4 27.5 24.7 417.5 25.854180.0 227.3 397.7 19.0 23.8 396.8 17.522112.6 198.2 CPSX- 1 OG CREFP- RUPTUR~
TI~E IX ~OU~S
RupTuoE TIK ~ X FI~AL C~EEP ~~AD1~6 TO eEACH
TEST OO~DIT10~ ~OURS ELO~G. R~t, hours X JLr~ lion l.OS Z.O~

1700~F/55.0 ksi 671.8 19.6 28.6 670.5 14.m 447.2 508.1 693.6 26.0 24.2 691.7 21.750 441.2 493.4 724.9 23.3 29.7 723.2 19.913 464.8 520~4 582.5 18.6 20.1 581.1 15.200 77.0 356.7 681.2 20.9 24.1 679.2 19.115 56.4 314.8 538.4 21.6 17.5 538.3 17.857 242.1 ' 308.7 523.0 17.7 21.8 522.4 14.157 235.3 308.0 569.7 17.5 19.8 568.5 15.035 287.0 354.9 1800-F/36.0 ~si m.2 29.6 29.3 773.8 28.826 315.0 539.9 719.7 29.5 28.5 717.8 27.266 321.2 486.4 741.6 28.0 25.9 740.3 24.870 284.5 464.2 6~2.8 45.6 34.7 681.1 39.289 409.1 452.4 764.0 23.2 33.7 764.0 '22.884 543.6 586.6 790.4 41.4 35.6 789.4 38.172 511.6 565.3 799.1 27.0 32.3 797.4 25.737 529.8 579.1 1850-Fn6.0 ~si 354.4 19.3 30.2 351.9 16.000 246.7 271.4 344.5 28.5 31.9 344.3 26.174 220.8 241.9 315.4 23.7 30.7 315.1 23.571 183.4 205.6 1922-F~20.3 ksi 753.4 31.7 34.8 753.2 27.914 352.3 462.1 728.0 31.5 33.5 727.1 28.362 281.1 422.1 731.6 34.3 38.8 730.5 30.770 339.3 437.3 1976-F/28.1 ksi 95.4 29.3 29.4 94.9 22.842 41.5 50.9 ff.7 26.7 27.2 94.7 20.130 45.8 54.7 104.6 30.4 33.2 104.4 27.517 41.8 54.4 100.8 25.6 35.1 98.9 21.577 49.2 58.1 95.8 25.9 28.9 93.6 19.7U 41.1 51.4 110.0 29.3 30.3 108.0 22.669 48.5 60.1 108.2 43.8 104.8 45.8 104 3 U.6 2000~F/18.0 ksi 464.4 23.1 21.3 463.6 18.190 257.7 293.5 411.9 18.3 23.0 410.4 16.347 103.5 227.6 370.9 27.0 38.7 369.8 25.326 7.6 47.3 2012~F/14.5 ksi 790.2 31.2 34.9 788.7 24.939 299.9 406.0 671.4 23.6 25.7 670.3 13.397 303.3 396.3 512.1 22.6 28.1 510.4 21.094 192.5 277.7 651.7 27.4 39.7 651.3 16.328 315.7 434.7 75-.6 29.7 25.4 753.1 24.032 193.8 388.7 908.3 17.7 18.3 758.9 30.8 26.5 758.7 24.090 388.7 438.2 740.0 19.8 20.5 739.5 16.962 316.5 426.7 671.5 26.4 23.8 669.3 15.578 359.8 .412.4 2050-F/15.0 ksi 410.8 22.9 27.4 410.0 18.655 ~6.5 272.2 283.5 18.0 31.2 283.5 15.303 156.4 191 2 320.0 16.8 17.4 318.3 12.979 156.4 191.2 389.7 ~.o Z2.1 389.7 18.488 29.9 189.1 381.4 27.0 24.1 381.1 24.758 69.5 197.9 2100-F/12.0 ksi 254.4 12.7 30.4 252.9 8.984 108.4 185.5 ~19.8 20.5 26.0 419.8 18.917 201.1 274.3 2100-F/12.5 ksi 331.4 16.9 21.7 331.1 15.069 25.2 83.2 367.7 19.2 23.2 366.5 17.530 76.2 177 4 387.3 16.8 17.2 386.5 12.742 236.9 282 0 383.1 34.1 32.4 381.6 32.135 10.5 164.3 . ~ ~
'- 2099,358 C~SX-lON CREEP-RUPTURE
TI~E IR RCURS
AUPTURE TI~E X X ~I~UL C~6Y READI~G TO R~ACN
TEST CO~DIT]ON ~ouRs ELONG. RA t. hours S JL~o.~ t;on 1.0X 2.0X

1800~F/36.0 ksi563.4 23.2 27.2 563.2 22.669 318.5 366.2 553.1 24.5 Z3.0 552.7 21.324 373.1 402.8 526.9 20.7 27.3 526.4 19.715 358.2 390.7 594.5 35.1 41.4 594.4 32.090 32B.8 372.8 lJS0-F/36.0 ksi242.9 24.3 20.1 242.2 20.686 107.3 155.6 221.9 17.0 21.0 221.0 14.888 115.9 150.4 223.4 21.3 21.0 221.7 19.196 128.4 144.7 1922-F/20.3 ksi5Z0.6 Z6.1 Z9.3 520.4 Z3.183 234.3 319.1 470.4 26.3 21.2 469.2 19.333 176.1 253.2 574.7 16.8 23.0 573.0 '14.411 282.1 373.0 2000~F/18.0 ksi434.0 21.5 18.7 432.1 20.234 103.5 233.1 437.3 27.1 33.8 437.3 '26.306 182.6 240.8 430.7 24.6 20.4 430.7 23.244 68.8 192.1 430.1 21.1 19.3 428.9 19.050 73.7 213.8 2050-F/15.0 ksi366.1 16.3 12.0 365.5 11.326 239.8 273.3 3~4.0 17.4 16.0 38~.3 12.055 168.2 242.9 420.2 12.2 13.3 418.6 10.017 127.3 273.2 203935~8 ~ TABLE 14 a9t-10 1 ~E P m~TutE
TUIE Ill 11~5 ~lUE TIIIE X X FIIUL aEEP IEADIIIG ~O ~EAUI
TEST er4DlT~ S ELO~IG. ItA t. ha~rsS Ar~n lial 1.m 2.0X
1800-F~36.0 ksi 565.1 35.2 32.0 564.8 29.774 297.0 368.9 581.9 32.4 29.3 580.228.689 371.9 402.5 514.1 24.1 30.2 514.121.207 318.3 358.2 1850/36.0 260.5 25.0 24.8 259.3 23.255 156.7 175.3 247.5 22.4 29.1 245.7 17.no 131.9 169.0 246.1 23.7 29.0 246.120.277 137.6 156.7 1922/20.3 916.3 24.9 30.3 914.8 ~.465 472.9 549.3 934.8 32.2 33.0 934.830.165 353.7 475.2 863.6 27.8 28.5 862.927.057 295.6 442.5 1976/28.1 116.1 19.5 20.1 116.1 19.155 57.4 70.1 65.6 22.9 20.6 64.2 21.368 17.8 26.4 91.6 23.2 25.3 90.4 15.544 37.6 49.7 ~000/18.0 430.1 ~.7 25.7 429.2 18.449 58.9 193.0 U3.8 19.8 25.1 U3.8 17.8~0 102.4 245.4 2050/15.0 397.7 17.9 30.0 397.3 13.264 239.8 292.9 U7.7 21.4 21.9 U7.1 18.854 248.2 318.4 468.3 18.4 25.5 467.915.800 194.1 300.1 2100/12.0 501.3 10.1 15.9 498.7 0.615 401.3 16.8 26.3 399.715.429 6.6 25.5 210.6 11.5 12.7 2~0.3 0.373 C~15X-12A ~ac~
TI~IE Ill lla~S
I~U7TlJtE TIIIE X X FIIIAL aEEP IEIIDI~G TO IIEACII
TEST ~IOIT1~llalFS EL~IIG. RA t. halrsS cLr~. tia~ 1.m 2.m 1800~F~36.0 l~si 491.9 40.2 41.6 491.8 38.605 254.0 293.7 420.4 23.5 31.9 420.319.299 234.9 277.9 383.4 25.3 26.2 3U.9 22.920 198.1 244.3 456.2 24.1 26.1 454.5 ~.582 89.9 265.5 458.0 30.7 32.7 457.126.155 253.2 292.8 386.8 30.1 30.4 3~6.327.031 172.7 216.9 403.7 34.5 28.8 402.731.033 140.2 204.9 398.7 21.6 23.5 398.420.277 181.1 236.1 1850/36.0 2W.5 32.1 40.5 2W.3 31.2U 10~.8 119.6 189.5 21.2 25.2 189.420.461 99.1 116.3 1922/20.3 829.6 46.5 45.3 U8.8 44.U8 315.8 400.7 797.0 33.5 32.5 796.932.856 315.3 400.5 200~/18.0 500.3 31.7 29.6 499.2 24.922 218.4 268.5 227.6 36.5 41.2 227.126.825 90.6 113.9 '30.4 18.5 23.3 430.418.180 181.0 234.1 2~50/15.0 U4.8 17.0 27.5 423.3 15.832 263.5 301.2 366.1 26.2 42.8 365.520.399 146.6 197.8 '.00.8 18.2 25.4 400.716.910 184.6 251.3 2100/12.0 255.4 25.8 45.8 253.6 22.920 64.1 125.8 U3.9 10.1 19.3 4U.7 8.602 378.6 421.9 325.1 7.1 16.6 324.7 4.315 268.8 302.5 CA 020993~8 1998-06-01 Varying the primary gamma prime aging treatment was explored with most of the development activity concentrated on achieving optimized gamma prime size and distribution through longer soak times at 1975~F (see Table 4) since higher temperature aging treatments accelerated TCP phase formation during the aging cycle.

Ten to twenty-one hour soak times at 1975~F were successful since they provided average gamma prime particles of about 0.5 ~m dimension. However, it appeared that shorter primary gamma prime aging time at higher temperature may be more practical, once more stable microstructures were defined.

Microchemical SEM WDX needle particle analyses was performed on a failed CMSX-lOG creep-rupture specimen. The specimen, tested at 1976~F/28.1 ksi condition, exhibited needles in its microstructure. The results of the analysis are reported in Table 16 below and indicate, again, that the needles formed in this class of material are particularly rich in Re, but are also enrichened with Cr and W.

(--' 20993,~8 '~ TABLE 16 CMSX-lOG

1976~F/28.1 ksi 104.6 HRS.
ELEM K Z A F

CRK 0.0426 1.105 0.793 1.049 COK 0.0584 1.094 0.888 1.086 NIK 0.1740 1.140 0.910 1.116 W L 0.2107 0.941 0.972 l.C00 REL 0.4767 0.941 0.979 1.000 NEEDLE CHEMISTRY

WT %~
ELEM CPS ELEM
CRK 113.7000 4.63 COK 112.1100 5.54 NIK 305.1425 15.02 W L 134.8988 23.03 REL 276.4000 51.76 100. 00 A standardized test for resistance to recrystallization was performed on a CMSX-lOG test bar. The test method and the results are reported in Table 17 below. The test results indicate that the CMSX-lOG specimen exhibited similar resistance to cast process/solution treatment/bonding process recrystallization level in comparison to CMSX-4 alloy.

20393~8 ;- TABLE 17 ~eehod: A controlled level of compressive stress is im4arted on the entire surface of an As-cast test bar. The bar is then solution heat treated. Follo~ing solution ~re~t ,t, the bar is sectioned and the tld.,~v~r~e section is observed ~etallograFhically. Depth of recrystallization measurements are taken.
Eva~uation Standards:

Resistance To RX
Anticipated in Blade Alloy Depth of RX Castings CMSX-4 .004" Very Good SX 792~ntire Bar Very Poor CYSX-10G .004" Very Good The CMSX-lOGa -lOIa, -12B,-12C, -lOJ, -lORi and -12Ri compositions were defined and evaluated. NO creep-rupture properties were generated for the CMSX-lOJ specimen, although test bars were produced and a solution heat treatment developed.
Again, the inclusion of C and B in the -lOJ composition appeared to have positive effect to single crystal test specimen yield.
Additionally, the lower levels of C and B than evaluated in CMSX-lOF specimen, particularly lower B, made the material more amenable to solution heat treatment. Ninety-eight to ninety-nine percent gamma prime solutioning was achieved, as opposed to the approximate 95% level typical of the CMSX-lOF composition.

The CM~X-lOC-~ -r.d -lOIa alloys were designed with NV3B numbers of about 1.70. These alloy specimens contain about 2.5~ Cr, 3.3-4.0% Co, 5.6-6.0% Re, greater Ta/W ratio, reduced Cb, and reduced Al+Ti content. Such reduction to Cb+Al+Ti level improved the - ~ 20993s8 solutioning char~cteristics of the materials tsee Table 4), plus assisted achieveme~ of increased alloy stability. Both specimens exhi~ited nearly ~ull gamma p~ime solutioning.

The lowered Ny;~ number continued t~ show effectiveness in providing be~ter creep-rupture capa~ility at temperature greate~
than lgOO~F, while ~aint~i~i ng extreme~y good creep-strength at lower ~empera~ure. ~MSX-lOGa test results from specimens produ~ed with impro~ed casting process co~tr~ls exhibited 700 hours o~ more life with abou~ 475 h~uxs required to creep to 1.0%
fs~ 1800~F/36.0 ~si condition. Fo~ higher temperature exposu~e, the specimen provided the improved average lif2 of abo~t 500 hou~s a~ 2050~F~lS.0 ksi condition and average 1.0~ creep d~formation ~hat occurred at about 250 hours, as indicated by the results reported in Table 18 below.

TABLE 18 ~-C~5X-lOCa CREEP-RUPTURE
TI~E IY bOURS
RUPTURE TI~E S X FI~AL C~EEP UEADI~G ro UEACH
TEST COXDITIC~ ~OURS ELO~G. U~ t. hour~ S JLf~. Lion l.OS 2.0S
1800-F/36.0 ksi 500.7 19.9 25.2 499.7 19.541 316.5 360.1 5K .2 29.1 25.4 583.9 26.395 370.0 401.8 505.1 Z .6 29.8 503.7 18.212 307 4 347 3 730.9 42.0 42.8 730.7 40.216 477 6 516 1 460.6 428.5 1850/36.0 lK .5 41.0 33.9 183.2 37.154 U .3 94.5 291.5 27.3 29.9 290.2 19.323 191.6 207.8 279.5 33.9 32.5 278.1 29.054 155 3 180.5 323.9 30.9 36.6 322.9 29.218 194 1 217.1 326.5l 8.9 12.6 295.2~ 33.3 33.5 174.1 162.3 300.1~ Z.8 22.4 1976/28.1 88.6 34.9 33.9 88.6 25.502 39 7 48 9 100.1 28.2 29.2 98.8 19.706 53 9 61 3 107.9 28.8 31.4 107.0 23.657 51.1 62.1 87.1 27.0 33.8 87.1 24.177 39.2 U.5 82.8 - 23.3 27.7 81.0 '17.301 20.6 38.0 88.2 31.2 35.2 86.4 24.463 33.6 44.4 83.7 34.0 34.3 83.4 29.718 36.3 45.1 114.1 24.3 26.3 113.0 20.544 62.1 73.2 1~ .3 18.3 21.3 120.7 15.740 76.5 86.0 117.7 23.2 25.6 117.5 Z .2K 78.0 85.3 ~I~TERRUPTED TESTS) 40.2 1.036 39.9 43.4 1.187 42.3 99.3l 60.1 38.8 127.9 41.5 34.5 127.5 37.493 51.2 62.6 96.8 22.9 27.9 96.5 20.124 45.9 54.4 118.9 31.3 27.1 118.0 24.603 49.5 61.3 111.1 25.0 22.8 110.2 21.521 46.4 58.0 96.6~ 24.1 22.9 120.51 25.8 29.4 113.01 27.6 20.5 197U18.85 ~I~TERRUPTED TESTS) 261.5 1.015 260.3 207.2 1.017 204.6 592.1 25.8 22.4 590.4 23.596 210.1 305.9 570.7 27.2 26.9 570.7 26.289 293.3 332.6 535.5 19.3 23.9 535.2 17.513 308.2 344.2 240.5 307.6 2050/15.0 536.8 28.5 27.5 535.6 20.662 232.3 321.3 49~.0 23.7 23.9 496.2 17.600 260.3 317.9 514.8 23.4 24.4 513.1 12.500 230.4 340.4 454.1 16.6 35.2 453.7 15.476 263.2 317.1 420.3~ 33.7 33.2 ~I~TERQUPTED TESTS) - - - 239.1 - - - 189.6 ~ - 2W .3 560.1i - 22.9 2012~14.5 536.6+ 7.3 8.1 424.6 2100/12.0 354.1 14.8 36.5 353.8 12.646 91.2 219.1 343.4~ - 27.2 - -91.4 147.2 ~91.01 - 16.7 1700/50.0 Xbchined From Bl~dk Specimen CA 020993~8 l998-06-0l 1~ creep strength is a significant property. Limiting creep strains to 1.0~ to 2.0~ is extremely important to gas turbine component design, since a component's usefulness is generally measured by its resistance to creep to an approximate 1-2~ level, not its ultimate rupture strength. Many prior art alloys may exhibit attractive rupture strength at the >1900~F level, however, they lack the level of useful strength, i.e., creep strength to 2.0~, that this invention provides in tandem with its far superior strength in test conditions below 1900~F.

The CMSX-lOIa specimens also provided significantly increased creep strength at the higher temperature extremes, but it did not appear to develop strength as good as the CMSX-lOGa specimens in lower temperature tests, as indicated by the results in Table 19 below.

CMSX-10 Ia CREEP RUPTURE

RUPTURE FINAL CREEP READING TIME IN HOURS
TIME % % % TO REACH
TEST CONDITION HOURS ELONG. RA t, hours deformation 1.0% 2.0%
1800~F/36.0 ksi 532.0 34.832.7 530.733.000259.1 312.5 474.6 23.8 29.2473.1 22.886201.0269.2 374.3 20.0 21.0372.8 19.238171.1214.7 1850/36.0 256.0 28.7 28.5256 0 27.867135.4157.1 251.4 34.4 30.3250.7 33.055121.6144.6 217.8 30.5 22.4217.2 27.000 94.2117.9 1976/28.1 85.7 27.5 28.983.8 21.754 36.9 46.2 81.9 33.6 31.881.0 24.384 32.1 42.1 68.9 26.1 25.867.6 20.960 23.1 32.4 2012/14.5 930.2 10.0 14.4928.4 9.649104.6455.7 844.4 17.7 23.2842.8 16.132339.7502.3 864.2 15.3 11.9862.8 14.558179.9453.4 2050/15.0 510.2 17.8 19.7508.4 15.703187.2312.7 528.6 17.9 24.2527.0 14.873293.7364.3 438.8 14.3 11.3436.4 13.556 56.0136.9 2100/12.0 616.4 19.0 19.1616.3 14.112 60.0422.5 467.7 l9.1 26 1466.0 11.373273.6374.8 CA 020993~8 l998-06-0l Similarly, CMSX-12B, with Nv3B at 1.80 level and additional chemistry balance as presented in Table 1, provided attractive creep strength at test condition greater than 1900~F, but did not perform quite as well as CMSX-lOGa in lower temperature tests, as indicated by the results reported in Table 20 below.

Table 20 RU~UK~ FINAL CREEP READING TIME IN HOURS
TINE % % % TO REACH
TEST CONDITION HO~RS ELONG. RA t, hours deformation 1.0% 2.0%
1976~F/28.1 ksi 91.7 15.3 17.2 91.214.070 43.9 56.2 72.6 19.4 23.2 72.6 17.396 27.4 36.8 14.1 5.0 1.3 12.7 2.300 8.6 11.9 98.1 16.9 17.6 96.4 13.670 17.8 38.9 108.2 25.2 24.1 108.0 22.794 43.8 S8.7 106.9 24.7 24.2 106.3 21.024 46.1 60.1 104.8 24.0 26.8 104.3 20.094 45.8 S8.7 104.3 26.8 21.4 103.2 22.347 48.6 60.8 1800/36.0515.0 24.7 24.2 513.3 19.468 320.1 358.0 536.4 23.2 21.1 530.8 22.184 318.3 359.5 304.7 13.2 19.9 302.9 12.582 166.0 200.8 1850/36.0262.6 18.4 23.1 262.4 17.660 12.5 142.2 2012/14.51031.3 17.2 18.51029.5 15.113 428.0 703.7 1078.7 15.6 20.01076.7 15.217 704.2 819.2 839.4 14.9 22.8 839.2 9.282 607.6 677.7 836.9 23.2 21.0 834.8 18.024 591.1 658.5 722.0 16.4 21.1 721.9 15.913 170.8 333.6 711.3 14.5 18.8 710.8 12.490 381.9 531.5 711.9 18.3 20.0 711.4 16.201 447.7 530.7 2050/15.0507.5 10.0 10.1 507.2 9.394 70.4 360.4 434.0 17.5 16.8 434.0 13.847 241.7 309.0 2100/12.0487.5 25.3 20.3 486.6 20.986 18.2 224.7 444.9 7.8 11.0 442.2 3.884 347.3 413.6 Alloy composition has the greatest effect on ultimate creep strength. However, some of the variation experienced between alloy derivatives, and particularly for tests exhibiting inconsistent results for a given alloy, can be caused by variation in casting process condition. Casting process thermal gradient variation affects the cast specimen dendrite arm spacing ,~' 2 ~93~8 ~

an~ ultimately, its response to solution heat treatment and primary gamma prime aging treatment. It must, therefore, be recognized that much of the creep-rupture results reported herein may have been generated under non-optimized conditions and may be capable of improvement. Improved casting process control may provide casting microstructures more amenable to solution treatment and study to determine the appropriate primary gamma prime aging treatment to provide the optimum gamma prime particle size, which may result in further mechanical property enhancement.

The CMSX-12C composition was designed to provide a calculated NV3a number of 1.70. The alloy Cr content was designed at 2.8% and Co set at 3.5% aim for this alloy. An attractive Ta/W ratio was maintained while Re content was moderate at 5.8%. The alloy's Al+Ti content was reduced, in comparison to the CMSX-12A and CMSX-12B specimens, to provide improved alloy response to solution procedure.

Similar to the CMSX-lOGa specimen, the CMSX-12C specimen exhibited an improved balance of creep strength for test condition ranging 1800-2100~F, as reported in Table 21 below.

~ABLE 21 ~SX~1ZC ~r~r R~ I~IRE 2 0 ~ 9 3 5 8 TI~E 1~ NOURS
HUPTURE TI~E Z S FI~AL CREEP READIHG TO ~EACH
TEST cO~DmO~ HoURS ELO~G. RAt. hours S J~t~ (ion l.oX 2.0Z
18000F/36.0 ksi 465.2 31.821.0 464.5 30.543 173.0 262.4 518.0 26.1 31.2517.9 24.947 288.1 334.3 U 0.9 28.3 33.6480.0 27.715 239.7 297.5 713.3 30.0 28.0713.2 28.899 455,0 503.7 1850/36.0 237.7 28.2 26.8237.7 27.054 114.4 145.3 ~1.2 22.9 27.3220.7 ~ .491 111.3 135.2 231.7 23.3 24.7231.0 Z.614 121.0 144.7 338.9 26.2 27.0337.5 23.256 216.0 236.3 300.1l - -295.2~ 33.3 33,5 1476/28.1 73.2 20.8 29.1 n.2 17.768 29.3 38.9 79.0 28.1 31.877.4 21.533 31.4 41.4 83.8 21.6 26.582.3 17.860 34.2 43.8 67.6 31.2 29.867.5 24.177 25.5 34.6 113.0~
79-4 30.8 76.2 32.8 68.8 29.3 118.1 26.0 28.0116.2 23.822 49.3 62.0 - - 29.0 ~I~TERRUPTED TESTS) - - Z9.4 - 32.9 1476/18.85 ksi 65.4 218.0 - - 271.9 (IHTERRUPTED TESTS) - - 168.9 - 116.4 240.5 2012/14.51001.8 23.6 20.01000.723.348 249.6 542.8 865.5 20.7 26.1864.8 18.807 418.2 569.3 61.9 267.1 2050/15.0 509.4 13.7 22.3508.012.860 158.1 315.1 546.4 15.6 23.6546.414.044 323.0 404.0 180.8 44.2 240.7 190.9 2100/12.0 404.3 11.2 21.6404.3 8.438 290.1 326.4 321.7 9.5 15.0320.4 7.671 156.6 254.1 545.1 8.2 Z2.1542.2 5.351 236.0 45Z.9 457.4 8.6 23.4455.8 6.612 309.3 380.9 21000F/12.0371.41 14.2 17.1 17500F/50.0446.9~ 16.8 20.4 19760F/18.85 476.6l 19.227.1 459.91 30.6 30 2 19760F/28.1 ksi 120.5~ 24.122.9 99.6~ 25.8 29.4 20500F/15.0 ksi 469.8 - 30.8 U5.4 - ~.~ - - - -20120F/14.5 ksi 638.1 521.8 267,1 61.9 395.7 Machit~d Ftom alade St~ecimens 4 6 CA 020993~8 1998-06-01 with improved casting process controls, this specimen has shown the following 1.0~ longitudinal creep strengths, as reported in Table 22 below.

Test Condition Time to 1.0~ Strain Hrs 1800~F/36.0 ksi455 2100~F/12.0 ksi309.3 Both alloys provide similarly greater rupture strength than CMSX-4 alloy at condition to 1976~F. Respective improvements to metal temperature capability are reported below in Table 23.

Approx. Strength Advantage TemperatureRelative to CMSX-4 1800~F 40~F
1850~F 45~F
1976~F 43~F

Based on 1.0~ creep strength, the respective approximate advantages are:

1800~F +46~F
1850~F +60~F
1976~F +55~F

Note that the comparison is not density corrected.

CA 020993~8 1998-06-01 For test temperature above 1976~F, the test results indicate that the CMSX-lOGa and CMSX-12C specimens provided slightly lower strength than CMSX-4 alloy. The reduction in strength advantage for these alloys is believed to be the result of TCP phase formation. To address this issue, the alloys CMSX-lOGb, CMSX-lOL, CMSX-12Ca, and CMSX-12E, are designed with N:3B number as low as 1.50 (see Table 1) to provide greater phasial stability, and effect much improved high temperature creep strength while maintaining most of the creep advantage demonstrated for the 1800-1976~F test regime.

The CMSX-lORi and CMSX-12Ri compositions were designed at the 1.91 and 1.92 Nv3~3 levels, respectively. These specimens were subjected to the most extensive testing of properties. They were designed with 2.65~ and 3.4~ respective Cr levels, with other features remaining similar to the aforementioned alloy design considerations. The properties generated for these two materials confirm the overall invention design concept with the other material iterations able to provide similar physical properties and relatively better blends of mechanical properties.

The CMSX-lORi and CMSX-12Ri specimens' respective creep-rupture capabilities are reported below in Tables 24 and 2S.

r- 2 0 9 9 3 ~ 8 C~SX-lO(Ri) CREEP-RUPTURE
TIXE IH HOURS
~UPTURE TI~E X X FI~UL CREEP READIHG TO oE~CH
TEST CO~DITIO~ HOURS ELO~G RA t, hours X defon ation l.OX 2.0S
16750F/75.0 ksi Z 7.3 21.2 33.8 225.4 14.359 5Z.8 131.5 231.6 19.3 31.0 231.3 16.671 51.0 125.1 Z3.4 17.0 22.3 223.3 15.360 68.5 126.6 1750/50.0 425.9 18.3 33.7 425.6 16.047 303.4 334.7 428.0 18.4 29.7 427.3 16.229 309.2 343.0 460.8 17.1 25.7 459.0 15.308 314.7 360.3 1800/36.0 698.5 39.9 34.3 696.8 36.980 492.8 521.5 676.3 28.3 33.3 674.5 27.221 479.0 513.8 692.9 38.5 31.3 692.2 36.494 469.3 504.9 1850/36.0 291.2 34.1 33.1291.1 , 31.774 194.1 210.4 260.0 29.3 32.1 258.8 25.321 170.2 186.4 2n.3 34.5 31.8 271.1 30.940 169.3 187.1 ~850/27.56614.0 52.0 42.0 613.5 50.482 365.8 415.5 576.3 49.7 39.0 575.9 49.183 345.1 368.2 481.1 40.4 35.4 480.7 38.294 309.3 335.4 1976/28.1 76.2 23.5 31.7 75.9 22.130 38.6 46.7 80.5 19.0 26.3 79.8 14.665 44.3 51.3 99.7 26.2 28.1 98.9 23.480 40.4 54.0 41.4 (INTERRUPTED TESTS) 37.0 40.5 1976/18.85 265.6 29.5 35.7 264.7 29.010 158.7 184.8 278.8 51.4 38.8 278.1 46.026 82.0 155.0 139.7 (I~TERRUPTED TESTS) 128.8 100.1 2012/14.5 490.8 40.2 33.5 490.5 37.678 286.5 335.3 447.0 37.0 41.5 445.0 32.814 291.4 319.9 - 113.5 (l~tERRUPTED TESTS) - - 205.7 - - 202.2 2050/15.0 251.9 33.6 35.9250.0 25.559100.0 149.5 318.9 27.1 30.0318.2 23.149177.5 221.2 - 181.0 ~ 95.5 (I~TERRUPTED TESTS) - - 34.5 2100/12.0 400.3 17.9 27.2400.1 17.877102.8 225.0 3~2,1 '5.3 22.9361.8 14.986125.7 217.2 389.5 19.9 24.0388.2 19.510 41.1 180.7 CA 020993~8 l998-06-0l CMSX-12(Ri) CREEP-RUPTURE
ku~l ~K~ FINAL CREEP READING TIME IN HOURS
TIME % % % TO REACH
TEST CONDITION HOURS ELONG. RA t, hours deformation 1.0% 2.0%
1675~F/75.0 ksi 209.8 22.3 23.1209.319.958 2.6 46.3 191.4 14.3 17.4 189.712.483 1.6 42.5 189.6 22.0 22.8 188.319.080 1.5 22.3 1750/50.0 448.1 26.7 26.6 447.926.054 302.3 335.5 403.1 19.0 26.9 401.918.566 310.0 290.2 435.0 19.4 26.9 434.418.503 89.1 284.1 1800/36.0 604.5 34.7 29.9 604.334.170 349.4 407.1 583.6 37.0 32.0 581.330.443 391.3 420.6 627.0 25.3 29.7 627.024.417 412.4 455.8 1850/36.0 302.9 33.1 31.3 301.729.034 198.9 215.1 314.4 32.0 27.1 312.727.479 201.4 220.2 1976/28.1 90.0 19.7 29.2 88.516.627 33.9 48.8 91.5 30.3 31.9 90.629.001 37.3 47.9 68.6 35.3 32.2 68.428.869 17.3 27.6 43.7 41.4 (INTERRUPTED TESTS) 38.7 2012/14.5 324.1 31.4 30.8 323.924.403 160.1 207.7 481.4 30.9 31.9 481.129.581 129.9 299.6 551.7 29.9 31.1 549.225.622 304.4 375.5 256.1 182.8 (INTERRUPTED TESTS) 101.5 2050/15.0 243.4 36.1 35.0243.3 20.614143.1 174.2 2100/12.0 374.8 12.1 20.3374.7 11.743166.6 280.4 463.6 15.4 25.9463.3 13.594245.7 363.3 488.0 20.3 25.9487.1 19.55025.7 118.9 The method and results of W and Re microstructural segregation investigation undertaken on fully solutioned and partially solutioned CMSX-12Ri test specimens are reported in Table 26 below. The investigation indicated that it is desirable to minimize the amount of microstructure-contained residual eutectic and that for fully solutioned specimens, the solution treatments developed for the invention are successful in minimizing elemental segregation, which is important in attaining optimized mechanical properties and microstructural stability.

' 20993S(8 .

Alloy: CYSX-12 Ri Test Specimen: 3t8" Dismeter Solid Bar Specimen Condition: Fu~ly Solutioned Solutioned ~ith 2.0X Residual Eutectic ~na~yses Hethod: ~icroprobe Analyses Random array of 350 points across a section st right ang~es to the growth direction Seven line scans, 51~ apart, 50 paint analyses per ~ine The standard deviation of the ~ and Re measurements are the measure of homogeneity Results:
aSX-12 Ri Standard Deviations ~ Re Ful~y Solutioned 0.27 0.50 2X ResiduaL Eutectic 0.36 0.90 Comcarison Typical CMSX-4 0.57 0.60 Table 27 below reports results of burner rig hot corrosion test undertaken with the CMSX-12Ri specimen. The measurements were taken at the bar location which experienced the maximum attack, i.e., 1652~F location, with the results showing the DS MAR M 002 alloy experienced approximately 20X more metal loss than the CMSX-12Ri specimen. Visual observation showed a similar result for the CMSX-lORi alloy. Both CMSX-lORi alloy and CMSX-12Ri alloy showed similar resistance to attack as CMSX-4 alloy based on visual specimen review at 60, 90 and 120 hours.

2099~8 HOT CORROSIOII
METHOO
Burner Rig 1742~F tffO~C) 2 p~n salt, standard fuel lleasurements taken at point of maxin~n attack ~hich vas at 1652~F (900~C~
t Measurements reported ~ere taken at the average mininun diameter of useful metal RESULTS
90 Hour Test Post Test Metal Loss AlloyInitial Dia. Useful Dia. Per Side DS ~lar M 002 6.88 nn 5.14 mn .87n~n ~.034") CMSX-12Ri6.86 mn 6.78 mn .04 n~n ~.0016") Table 28 below reports the results of cyclic oxidation tests undertaken at 2012~F with March 1 gas velocity. The CMSX-12Ri specimen was similarly resistant to oxidation attack at 2012~F, however, it was not as good as CMSX-4 at approximately 1886~F
exposure.

' ~ 209g35~
'~ TABLE 28 Cyclic Oxidation Test 15 Mirlute Cycles to 2012~F ~1100~C), Cooled To Ambient Bet~een Cycles Yach 1 Gas Velocity 89 Hours Total ~ith 77 Hours at 2012~F

CMsX-12 Ri RESULT: at 1100~C Approx. 0.1 mm loss per side for every 300 cycles Approx. 0.1 mm loss per side for every 380 cycles at 1030~C CMSX-12 Ri Approx. .105 mm loss per side after 355 cycles Approx. .03 mm loss per side after 355 cycles CMSX-12Ri elevated temperature tensile data is reported in Table 29 below, while the results of impact tests are reported in Table 30 below. The CMSX-12Ri elevated temperature impact strength minimum is similar to CMSX-4 and its maximum occurrin~ at 1742~F, is better.

- - --' '~ 2033~5~3 ~

TE~SILE DATA
CMsx-12 Ri Alloy Test RA
Temp UUE0.1X ~ld 0.2Z ~ld UTS Elong %
~F ksi ksi ksi %
1382 2.3~ 150.0 160.8 188.7 13 14 1382 2.3~ 153.6 165.1 190.0 13 15 1562 6.2~ 136.5 130.8 152.3 27 24 1562 6.2~ 135.0 128.9 160.1 25 23 1742 5.6~ 92.7 89.2 125.3 24 30 1742 5.6~ 99.9 106.2 129.2 24 32 1922 3.8~ 69.5 74.3 104.1 19 36 1922 3.8~ 72.4 77.6 106.0 19 36 IMPACT DATA
CMSX-12 Ri 0.35 Snch Diameter Plain Cylindrical Specimens rest Temperature, ~F

:MSX-12 Ri ~1 only) 26 J 20 J 60 J 32 J

~MSX-4 (Ave. of 4) 26 J 21 J 42 J 45 J

Further Im4act Property Comparison CMSX-2 -- Min. Impact Strength 16.5 Joules SRR 99 -- Min. Impact Strength 20 Joules ~-- 2099358 (-The results of CMSX-12Ri low cycle fatigue tests undertaken at 1382CF and 1742~F test conditions, with R = O, are reported in Table 31 below. The data indicates that CMSX-12Ri performance is similar to CMSX-4 at 1382~F condition, while the alloy exhibits approximately 2.5 times the typical CMSX-4 life at 1742~F
condition.

La.J CrCLE FATIGUE
CHsX-12 Ri Al loy R = O (zero to maximum stressing) 1382~F (750~C) 1742~F (950~F) PE~K STRESS PEAK STRESS
ksi ~HPa) Cyclesksi (HPa) Cycles 142 ~980) 8686 110~760) 4427 130 ~900) 11950 ~,9 ~600) 15510 125 ~860) 20810 87 ~600) 37430 119 ~820) >100000 75 (520), 92580 . Performance is similar to asx-4 at 1382~F (750~C) . Ca~ared to CMSX-4, at 1742~F (950~C) and in the 20000 eycle region, CHSX-12Ri exhibits 2.5 times life or 15X on strength.

Notched low cycle fatigue test results show the CMSX-12Ri is 2-1/2 times better than CMSX-4 out to about 30000 cycles, while at 50000 cycles and above, the alloy performance is similar to CMSX-~. The r~sults cf these tests performed at 1382~F, Kt = 2.û and R=O test condition, are reported in Table 32 below.

NOTCHED LO~ C'~CLE FATIGUE
aSX-12Ri A~loy 1382~F (~0~C), ~t = 2.0, R = O

PEAK STRESS CYCLES
ksi (MPa) 113.13 (780) ~4879 107.33 (740) 9784 95.72 (660) 28470 84.12 (580) 49810 81.22 (S60) 78.32 (540) > 115,000 7S.42 (520) > 115,000 . Results are 2-1/2 times better than CMSX-4 out to about 30000 cycles.
. Results are similar to CMSX-4 at 50000 cycles, and above.

High cycle fatigue test results for the CMSX-lORi alloy are rep~rted in Table 33 below. For 1742~F, 100 Hz, R = O test condition, the alloy exhibited about 2-1/2 times the typical CMSX-4 lives.

L Z . ... ..

2~33~8 (-H I GH CYCLE FAT I WE
CMSX-lORi Alloy 1742~F (ff O~C), 100 Hz., R = O

PEAK STRESS CYCLES
ksi ~MPa) ~Nf) 81.22 ~560) 15.2 x 106 92.82 ~640) 3.59 x 106 104.43 ~ nO) 0.6 x 106 Lives are 2-1/2 times better than CMSX-4 The CMSX-lORi and CMSX-12Ri test data indicates that adequate hot corrosion and oxidation resistance can be achieved with extremely low alloy chromium content. Additionally, extremely good thermo-mechanical fatigue tensile and impact strengths are apparent with the superalloys of this invention.

The results of alloy specimen density measurements are reported in Table 34 below.

209g3S8 ''-' ~ TABLE 34 5~NGLE C2YSTAL ALLOY DENSITY DATA

A~OY ~ha~

aSX-lOA .324 asx- lOB .3254 asx-10C .325 SX-lOD .325 asx- 1 OE .323 asx- lOF .322 asx- lOG .322 asx- lOGa 324 asx- lOH
asx-lol .33222 asx- lol a 327 asx-lOJ 329 asx-lzA .332235 asx-l2B .326 asx-12C .32 asx- 12Ca tl2D) .326 asX-lORi 323 aSX-10 Ri The alloys of this invention are amenable to HIP processing.
Specimens HIP treated as reported in Table 35 below, showed nearly complete pore closure and absence from incipient melting.

20!~93~8 TABT.~ 35 H~P condi~ion 1. Heat Specimens in ~he HIP ~essel to 245~~F ~t mininum Argon p~essure ~app~oximately 1500 psi~ and hold for 4 houxs while main~aining 2455~/1500 psi condition.
2. ~hile mai~taining t~e 2~55~F
operating temperat~re, increase the Argon p~essure over 1 hour to 2~
~si. Soa~ specimens at a455~~/20 ksi conditio~ f~r 4 houxs.

~hile this inventio~ h~s ~een described with respect to particular em~odiments ~hereof~ it is apparent that n~meroUs other forms and modific~tions of this in~ention wil~ be o~ious to ~hose s~illed in the art. The appen~ed ~laims and thi5 invention ~ener~lly should ~e const~d to co~er ~11 such ob~ious forms and modifica~ions which ~re w~thin the true spirit and scope of ~he present in~ention.

Claims (32)

1. A nickel-based superalloy comprising the following elements in percent by weight:

Rhenium about 5.0-7.0 Chromium about 1.8-4.0 Cobalt about 1.5-9.0 Tantalum about 7.0-10.0 Tungsten about 3.5-7.5 Aluminum about 5.0-7.0 Titanium about 0.1-1.2 Columbium about 0-0.5 Molybdenum about 0.25-2.0 Hafnium about 0-0.15 Carbon (Incidental about 0-0.04 Impurity) Nickel + Other balance Incidental Impurities said superalloy having a phasial stability number NV3B less than about 2.10.

2. The superalloy of Claim 1 further comprising the following elements in percent by weight:

Boron max. about 0.01 Yttrium max. about 0.01 Cerium max. about 0.01 Lanthanum max. about 0.01

3. The superalloy of Claim 1 further comprising the following elements in percent by weight:

Manganese max. about 0.04 Silicon max. about 0.05 Zirconium max. about 0.01 Sulfur max. about 0.001 Vanadium max. about 0.10

4. The superalloy of Claim 1 wherein said superalloy has a phasial stability number NV3B less than 1.85.

5. The superalloy of Claim 1 wherein said superalloy has a chromium content of from about 1.8 to about 3.0 percent by weight.

6. The superalloy of Claim 1 wherein said superalloy has a rhenium content of from about 5.5 to about 6.5 percent by weight.

7. The superalloy of Claim 1 wherein said superalloy has a cobalt content of from about 2.0 to about 5.0 percent by weight.

8. The superalloy of Claim 1 wherein said superalloy has an increased resistance to creep under high stress, high temperature conditions up to about 1975°F.

9. A single crystal article made from the superalloy of Claim 1.

10. The single crystal article of Claim 9 wherein the article is a component for a turbine engine.

11. The article of Claim 10 wherein the component is a gas turbine blade or gas turbine vane.

12. A nickel-based superalloy comprising the following elements in percent by weight:
Rhenium 5.5-6.5 Chromium 1.8-3.0 Cobalt 2.0-5.0 Tantalum 8.0-10.0 Tungsten 3.5-6.5 Aluminum 5.3-6.5 Titanium 0.2-0.8 Columbium 0-0.3 Molybdenum 0.25-1.5 Hafnium 0.02-0.05 Carbon (Incidental 0-0-04 Impurity) Nickel + Other balance Incidental Impurities said superalloy having a phasial stability number NV3B less than about 1.75.

13. The superalloy of Claim 12 further comprising the following elements in percent by weight:

Boron max. 0.01 Yttrium max. 0.01 Cerium max. 0.01 Lanthanum max. 0.01

14. The superalloy of Claim 12 further comprising the following elements in percent by weight:

Manganese max. 0.04 Silicon max. 0.05 Zirconium max. 0.01 Sulfur max. 0.001 Vanadium max. 0.10

15. A single crystal article made from the superalloy of Claim 12.

16. A single crystal casting to be used under high stress, high temperature conditions characterized by an increased resistance to creep under such conditions, said casting being made from a nickel-based superalloy comprising the following elements in percent by weight:

Rhenium about 5.0-7.0 Chromium about 1.8-4.0 Cobalt about 1.5-9.0 Tantalum about 7.0-10.0 Tungsten about 3.5-7.5 Aluminum about 5.0-7.0 Titanium about 0.1-1.2 Columbium about 0-0.5 Molybdenum about 0.25-2.0 Hafnium about 0-0.15 Carbon (Incidental about 0-0.04 Impurity) Nickel + Other balance Incidental Impurities said superalloy having a phasial stability number NV3B less than about 2.10.

17. The single crystal casting of Claim 16 wherein said superalloy further comprises the following elements in percent by weight:

Boron max. about 0.01 Yttrium max. about 0.01 Cerium max. about 0.01 Lanthanum max. about 0.01

18. The single crystal casting of Claim 16 further comprising the following elements in percent by weight:

Manganese max. about 0.04 Silicon max. about 0.05 Zirconium max. about 0.01 Sulfur max. about 0.001 Vanadium max. about 0.10

19. The single crystal casting of Claim 16 wherein said superalloy has a phasial stability number NV3B less than 1.85.

20. The single crystal casting of Claim 16 wherein said superalloy has a chromium content of from about 1.8 to about 3.0 percent by weight.

21. The single crystal casting of Claim 20 wherein said superalloy has a rhenium content of from about 5.5 to about 6.5 percent by weight.

22. The single crystal casting of Claim 21 wherein said superalloy has a cobalt content of from about 2.0 to about 5.0 percent by weight.

23. The single crystal casting of Claim 16 wherein said superalloy has an increased resistance to creep under high stress, high temperature conditions up to about 1975°F.

24. The single crystal casting of Claim 16 wherein said casting has been aged at a temperature of from about 1950°F to about 2125°F for about 1 to about 20 hours.

25. The single crystal casting of Claim 16 wherein said casting is a component for a turbine engine.

26. The single crystal casting of Claim 16 wherein said casting is a gas turbine blade.

27. The single crystal casting of Claim 16 wherein said casting is a gas turbine vane.

28. A single crystal casting to be used under high stress, high temperature conditions up to about 1975°F characterized by an increased resistance to creep under such conditions, said casting being made from a nickel-based superalloy comprising the following elements in percent by weight:

Rhenium about 5.5-6.5 Chromium about 1.8-3.0 Cobalt about 2.0-5.0 Tantalum about 8.0-10.0 Tungsten about 3.5-6.5 Aluminum about 5.3-6.5 Titanium about 0.2-0.8 Columbium about 0-0.3 Molybdenum about 0.25-1.5 Hafnium about 0.02-0.05 Carbon about 0-0.04 Boron about 0-0.01 Yttrium about 0-0.01 Cerium about 0-0.01 Lanthanum about 0-0.01 Manganese about 0-0.04 Silicon about 0-0.05 Zirconium about 0-0.01 Sulfur about 0-0.001 Vanadium about 0-0.10 Nickel balance said superalloy having a phasial stability number NV3B less than about 1.75.

29. The single crystal casting of Claim 28 wherein said casting has been aged at a temperature of from 1950°F to 2125°F for 1 to 20 hours.

30. The single crystal casting of Claim 28 wherein said casting is a component for a turbine engine.

31. The single crystal casting of Claim 28 wherein said casting is a gas turbine blade.

32. The single crystal casting of Claim 28 wherein said casting is a gas turbine vane.