super alloys development

Aerosp. Sci. Technol.3 (1999) 513–523 1999 Éditions scientifiques et médicales Elsevier SAS. All rights reservedS1270-9638(99)00108-X/FLA

Evolution of Ni-based superalloys for single crystal gas turbine bladeapplications

Pierre Carona,* , Tasadduq Khanb

a ONERA, Metallic Materials and Processes Department, BP72, 92322 Châtillon cedex, Franceb ONERA, General Scientific Directorate BP72, 92322 Châtillon cedex, France

Received 19 July 1999; accepted 4 October 1999

Abstract The chemistry of the Ni-based superalloys designed for single crystal gas turbine blades has significantlyevolved since the development of the first generation of alloys derived from columnar grained materials. Theoverall performance of the second and third generations has been significantly improved by the addition ofincreasing amounts of rhenium. However, the problems of increased density, grain defects and microstructuralstability have also become more and more acute and render necessary to carefully control the level of thevarious alloying elements in order to effectively benefit from the high potential of the most recently developedthird generation alloys. 1999 Éditions scientifiques et médicales Elsevier SAS

superalloy / single crystal / rhenium

Résumé Evolution des superalliages à base de nickel destinés aux aubes monocristallines de turbine à gaz.La chimie des superalliages à base de nickel destinés aux aubes monocristallines de turbine à gaz a évoluéde manière significative depuis le développement des alliages de première génération dérivés des matériauxà grains colonnaires. La performance d’ensemble des seconde et troisième générations d’alliages a étégrandement améliorée par l’addition de quantités croissantes de rhénium. Cependant, les problèmes de massevolumique accrue, de grains parasites et de stabilité microstructurale sont aussi devenus de plus en plus aiguset rendent nécessaire le contrôle soigné des concentrations des différents éléments d’alliage afin de pouvoireffectivement bénéficier du potentiel élevé des alliages de troisième génération les plus récents. 1999Éditions scientifiques et médicales Elsevier SAS

superalliage / monocristal / rhénium

1. Introduction

Turbine blades are critical components in both aero-nautical and stationary gas turbines. The engine perfor-mance is closely related to the capability of materials towithstand higher and higher temperatures. Over the past25 years, the turbine blade temperature capability has in-creased significantly. Some of the advances have been

* Correspondence and reprints

achieved through improved alloy compositions, othershave been accomplished by major innovation in process-ing, such as directional solidification or single crystaltechnology of nickel-based superalloys. Although thenickel based superalloys are certainly approaching theirtemperature asymptote, there is still room left for somefurther development in single crystal superalloys, the so-called third generation alloys.

514 P. Caron, T. Khan / Aerosp. Sci. Technol. 3 (1999) 513–523

About twenty years ago, the chemistries of the firstsuperalloys suited for single crystal casting derived fromthat of well-known nickel-based superalloys initiallydesigned for conventional casting, such as Mar-M200.The major chemistry modification as compared to thepolycrystalline materials was the suppression of thegrain boundary strengthening elements C, B, Zr andHf. The alloys’ designers then succeeded in optimisingthe mechanical properties of these alloys, especiallythe creep resistance, by introducing large amounts ofrefractory alloying elements such as W, Ta, Mo. A majorstep was the introduction of Re at a level up to 6 wt.%,but at the expense of the density, castability, cost andmicrostructural stability. The most recent developmentsincluded the addition of more exotic elements such asruthenium and iridium. In some alloys, minor elementssuch as carbon and boron were recently re-introduced inorder to render the material more tolerant to problemsrelated to large angle sub-grain boundaries with theobjective of reducing the number of rejected blades. Itis the purpose of this paper to describe the evolution ofthe chemistry of this class of alloys by focusing on theadvantages and drawbacks resulting from these chemistrychanges.

2. Evolution of the chemistry of the single crystalsuperalloys

2.1. First generation alloys

As detailed intable I the alloying elements of thefirst generation Ni-based single crystal superalloys aremainly Cr, Co, Mo, W, Al, Ti, Ta and sometimes Nb or V.The effects of these alloying elements on the propertiesof superalloys have been extensively studied during thelast forty years and are currently well-known [32]. Cr,Co and Mo partition preferentially to the austenitic face-centred cubic nickel-basedγ matrix where they act

mainly as solid solution strengthening elements. Cr alsoplays an essential role in the hot corrosion resistancesince it promotes the formation of a protective Cr2O3oxide scale. These alloys contain a high volume fractionof strengthening ordered Ni3Al-basedγ ′ phase particleshomogeneously distributed in theγ matrix as near-cubical precipitates (figure 1).

The elements Ti, Ta, Nb and V strengthen theγ ′precipitates by substituting to Al in Ni3Al. Al also playsa fundamental role in promoting the formation of astable Al2O3 alumina surface scale which protects thealloy against further oxidation.

The superalloys for single crystal blades do not gener-ally contain voluntary additions of minor elements suchas C, B, Zr or Hf used as grain boundary strengtheningelements in conventionally cast or columnar grained su-peralloys. As an example, the alloy NASAIR 100 was

Figure 1. Two-phaseγ -γ ′ microstructure in fully heat-treatedAM3 first generation single crystal superalloy.

Table I. Chemical compositions (wt.%) of first generation Ni-based superalloys for single crystal blades.

Alloy Cr Co Mo W Al Ti Ta Nb V Hf Density (g.cm−3) Country Ref.

Nasair 100 9 – 1 10.5 5.75 1.2 3.3 – – – 8.54 USA [37]

CMSX-2 8 4.6 0.6 8 5.6 1 6 – – – 8.60 USA [21]

CMSX-3 8 4.6 0.6 8 5.6 1 6 – 0.1 8.60 USA [21]

CMSX-6 9.8 5 3 – 4.8 4.7 2 – 0.1 7.98 USA [41]

PWA 1480 10 5 – 4 5 1.5 12 – – – 8.70 USA [20]

SRR 99 8 5 – 10 5.5 2.2 3 – – – 8.56 GB [18]

RR 2000 10 15 3 – 5.5 4 – – 1 – 7.87 GB [18]

René N4 9 8 2 6 3.7 4.2 4 0.5 – – 8.56 USA [33,42]

AM1 7.8 6.5 2 5.7 5.2 1.1 7.9 – – – 8.60 F [12]

AM3 8 5.5 2.25 5 6 2 3.5 – – – 8.25 F [24]

P. Caron, T. Khan / Aerosp. Sci. Technol. 3 (1999) 513–523 515

derived from Mar-M247 essentially by suppressing thesegrain boundary strengthening elements [37]. The result-ing increase of the incipient melting temperature from1240 to 1330◦C allowed a complete solutioning of thesecondaryγ ′ precipitates and an almost complete elim-ination of coarse primary interdendriticγ ′ particles byusing a super-solvus high temperature solution heat treat-ment. The resulting increase in creep strength was shownto be equivalent to a 58◦C temperature advantage at hightemperatures and low stresses.

Several alloy compositions were then developed byengine manufacturers, alloy makers or research institutesin order to improve the performance of the single crystalturbine blades (table I). The main goal was to improve thecreep strength which is an important damage process forthese components, together with the thermo-mechanicalfatigue.

Improvement in mechanical properties was thereforeobtained by a careful balance between the various alloy-ing elements. The main differences between the alloysconcern the respective levels of refractory elements W,Mo and Ta which are known to reduce the bulk diffu-sion rate that slows down the coarsening kinetics of thestrengtheningγ ′ precipitates and the diffusion controlledcreep mechanisms such as climb or cross-slip of disloca-tions. These elements also induced strong solid solutionhardening effects due to their large atomic radii comparedto that of Ni or Al.

The CMSX-2 and CMSX-3 single crystal (SC) alloyswere derived from Mar-M247 [21], whereas the low den-sity SC superalloy CMSX6 was derived from the alloyIN6212 [41]. The SC alloys PWA1480 [20] and RenéN4 [33,42] were developed respectively by Pratt & Whit-ney Aircraft and General Electric; the alloys SRR99

and RR2000 were developed by Rolls Royce in UnitedKingdom [18]. The SC alloy AM1, jointly developedby ONERA, SNECMA, ARMINES and TECPHY [12],was selected by SNECMA as a blade and vane materialfor its M88 engine for the RAFALE fighter. The alloyAM3 [24], developed by ONERA, is now used as a SCblade material in the ARRIEL 2 and ARRIUS 2 TUR-BOMECA engines which equip a number of helicopters.

As demonstrated with NASAIR 100, one great advan-tage of the first generation SC superalloys compared toequiaxed conventionally cast (CC) or columnar graineddirectionally solidified (DS) alloys was the possibilityto achieve a homogeneous microstructure using a singlehigh temperature heat treatment which allowed the elim-ination of almost all theγ /γ ′ eutectic interdendritic nod-ules and solutioning of all the secondaryγ ′ precipitates.This solution heat treatment has to be performed withinthe so-called “heat treatment window” (interval betweenthe incipient melting temperature and the solvus temper-ature of the coarse interdendritic secondaryγ ′ precipi-tates). Such an example is provided by the alloy AM3 inwhich theγ ′ phase can be totally solutioned by using asimple treatment for 3 hours at 1300◦C (figure 2). Exceptfor PWA1480 [20] which contains a very high level of Taand consequently large amounts ofγ /γ ′ eutectic phase,the first generation superalloys are therefore easy to ho-mogenise, that allows us to obtain an optimised distrib-ution of fineγ ′ precipitates by applying adequate subse-quent ageing heat treatments. The ageing heat treatmentsperformed on single crystal superalloy blades correspondgenerally to the diffusion coating treatment cycles. Thepioneering work on CMSX-2 has shown that the creepbehaviour of highγ ′ volume fraction alloys can be opti-mised by choosing a precipitation heat treatment leading

(a) (b)

Figure 2. Transverse dendritic microstructure of AM3 single crystals; (a) as-cast; (b) heat-treated for 3 hours at 1300◦C, air cooled.


to cubicalγ ′ precipitates with a mean cube edge of ap-proximately 0.45µm [5]. Since that time alloy designershave attempted to balance these diffusion treatments inorder to achieve optimisedγ ′ size for creep strength.

2.2. Second generation alloys

Several alloy designers showed that a significant im-provement of the creep strength of the single crystalsuperalloys may be obtained by the addition of rheniumat the expense of other refractory elements such as Moor W. Thus, a study carried out on modified Mar-M200SC alloys showed that additions of rhenium substantiallylower theγ ′ coarsening kinetics and result in large nega-tive γ -γ ′ misfits [19]. Atom-probe studies performed onCMSX-2 and PWA-1480 alloys modified by additions ofrhenium also showed the existence of Re atom clusterswithin theγ matrix of these alloys, which is a more po-tent source of strengthening than the conventional solidsolution effect [3,4]. The introduction of 3 wt.% Re toSC nickel based superalloys resulted in a temperaturecapability improvement of about 30◦C [15]. The firstrhenium-containing alloys were labelled as second gen-eration SC superalloys. Typical examples of this class ofalloys are PWA1484 [10], René N5 [43], SC180 [44],CMSX-4 [22] and SMP14 [40]. The chemical composi-tions of these alloys are reported intable II.

However, a rational criterion for the definition of thesesecond generation alloys should be based on the hightemperature creep strength. In these conditions, the MC2superalloy developed by ONERA is now considered as anintegral part of this group, as it shows a high temperaturecreep strength comparable to that of second generationRe-bearing alloys, in spite of the fact that it does notcontain rhenium [6]. It is noteworthy that the absence ofrhenium makes this alloy much cheaper and lighter thanthe other second generation SC alloys.

A major issue of concern in developing second gen-eration alloys, and more specifically Re-bearing alloys,was their microstructural stability, i.e. the propensity toform undesirable topologically close-packed (TCP) brit-tle phases asσ , µ or P phases during exposure at hightemperature. The presence of TCP phases in superalloys

may sometimes induce deleterious effects on the me-chanical properties, such as loss of ductility, decrease ofthe impact strength or sometimes decrease of the creepstrength as demonstrated by Dreshfield and Ashbrook inIN100 [13]. It is generally recognised that superalloyswhich contain high levels of refractory elements such asRe, Mo or W are more or less prone to the precipitationof TCP phases. Re-rich TCP phases have been thus re-ported to precipitate in CMSX-4 [23], whereas W-richrhombohedralµ phase precipitates within the MC2 al-loy under certain conditions [30]. A study performed ona series of experimental Re-bearing alloy showed thatthree types of TCP phases can coexist in the same alloy,rhombohedralµ, tetragonalσ and orthorhombic P, withsimilar compositions, all rich in Re and Cr [11]. It washowever demonstrated that a small amount ofµ phasein MC2 superalloy does not affect its mechanical behav-iour [30], and it has been reported that the Re-rich TCPprecipitates present in CMSX-4 did not detrimentally af-fect creep rupture properties [23]. On the other hand, astudy performed on a series of experimental SC superal-loys with various amounts of Re showed a drastic reduc-tion in rupture life when the amount of Re-richσ phaseprecipitating in these alloys increased [14]. This deleteri-ous effect was attributed mostly to the depletion of theγmatrix of refractory strengthening elements.

As Re partitions preferentially to theγ matrix, thelevels of otherγ -former elements Cr, Co, Mo and Wmust be reduced and carefully balanced in order toavoid the supersaturation of the solid solution, which, inturn, will lead to the precipitation of Re-rich TCP phaseparticles, while maintaining a good balance between themechanical and the environmental properties.

Thus, as the typical level of Cr encountered within thefirst generation SC alloys was around 8 wt.%, it decreasedto 5–7 wt.% in the second generation alloys. It was shownthat this Cr content is sufficient to ensure an acceptablelevel of corrosion resistance and the matrix can keep intosolution high amounts of Re, W and Mo, necessary for anoptimised creep strength.

The role of theγ ′-forming elements Al, Ti, Ta, Nbon the phase stability is also important as their totalamount determines the quantity ofγ ′ phase which is

Table II. Chemical compositions (wt.%) of second generation Ni-based superalloys for single crystal blades.

Alloy Cr Co Mo Re W Al Ti Ta Nb Hf Density (g.cm−3) Country Ref.

CMSX-4 6.5 9 0.6 3 6 5.6 1 6.5 – 0.1 8.70 USA [22]

PWA 1484 5 10 2 3 6 5.6 – 8.7 – 0.1 8.95 USA [10]

René N5 7 8 2 3 5 6.2 – 7 – 0.2 8.70 USA [43]

SC180 5 10 2 3 5 5.2 1 8.5 – 0.1 8.84 USA [44]

SMP14 4.8 8.1 1 3.9 7.6 5.4 – 7.2 1.4 – 9.02 RSA [40]

MC2 8 5 2 – 8 5 1.5 6 – – 8.63 F [6]


close to 70 vol.% in the SC superalloys. An excessiveamount of γ ′ phase renders the matrix more proneto the precipitation of TCP phases due to the higherconcentration of elements which partition preferentiallyto it.

Methods have been developed to predict the propensityof superalloys to the precipitation of TCP phases. Theoldest one is the electron vacancy (Nv) or PHACOMPmethod [35]. The principle is to calculate an averageelectron vacancy numberNv for the alloy matrix, usingspecificNv values for the transition elements, and tocompare this value to a critical one above which the TCPphase is known to precipitate. It should be noted that thecriticalNv value varies with the nature of TCP phase andthe cast alloys, due to the dendritic segregations, havelower critical Nv values than the wrought alloys. Thisis of prime importance in Re-bearing alloys because thiselement segregates strongly to the dendrite cores and isdifficult to homogenise due to its low bulk diffusivity.

A more recent predictive method, known as the NewPhacomp method, has been developed by Morinaga etal. [27]. A parameter Md is calculated which is an averageenergy level of d orbitals of alloying transition elementsin the γ matrix. As in the case of PHACOMP method,an alloy will be prone to the precipitation of TCP phaseswhen the Md value becomes larger than a critical valuewhich depends on the phase type.

The level of alloying elements must also be balancedconsidering other properties which can be of prime im-portance for the application. The environmental proper-ties therefore must not be neglected even if the turbineblades are systematically coated to protect them againstoxidising and corrosive combustion gas. Indeed, theseprotective coatings, aluminides or MCrAlY types, maycrack under thermo-mechanical stresses or could be elim-inated by erosion. The remaining life of the componentwill then depend on the intrinsic resistance of the alloy tohigh temperature oxidation and corrosion attack. Thus,Mo is generally kept at a low level, because this elementis known to have a negative influence on the corrosionresistance of the nickel based superalloys, while alloy de-signers prefer to increase the level of Ta which is morebeneficial for environmental properties. It is also gener-ally considered that Re contributes positively to the hotcorrosion and oxidation resistance partly due to its effecton the diffusion of reactive elements.

The castability of single crystal components made ofRe-containing superalloys must be also assessed care-fully because it was demonstrated that it can be influ-enced negatively by excessive additions of Re and W.Thus, the occurrence of small chains of equiaxed grainsreferred to as freckles was observed in experimental SCsuperalloys containing high levels of Re and W [31].These freckles form due to convective instabilities result-ing from solute partitioning that lowers the density of theliquid in the mushy zone. The dendrite cores are enrichedin Re and W and the freckles form from the low den-

sity interdendritic liquid enriched in Al, Ta and Ti. Someisolated spurious grains can also form during directionalsolidification and the potential for nucleation and growthof these grain defects was also demonstrated to increasewith the levels of Re and W. The nucleation of spuriousgrains is due to the fragmentation of secondary and ter-tiary dendrite arms under the action of the interdendriticfluid flow generated by density differences. This tendencycan however be counterbalanced by increasing the levelof Ta which is rejected to the interdendritic areas, there-fore decreasing the density inversions. The alloys whichwere most resistant to freckling were therefore high in Taand low in Re and W [31].

An other drawback due to Re addition is that itincreases the density as compared to the first generationalloys, which is a drawback for the rotating parts ofaeronautical engines. The density of these alloys are inthe range 8.7–9 g.cm−3 whereas the density of MC2 isonly 8.6 g.cm−3.

2.3. Third generation alloys

More recently, alloy designers tried to improve againthe high temperature capability of the SC blade alloys byincreasing the content of rhenium up to about 6 wt.%.The challenge was to achieve improved creep strength,without increasing the density and by keeping the alloynot too much prone to the precipitation TCP phases. Twotypical third generation alloys are CMSX-10 developedby Cannon-Muskegon [17] and René N6 developed byGeneral Electric [39]. More recent development workconducted by GE was devoted to third generation SCalloys containing also some additions of ruthenium [28].A new generation of SC alloys, a typical example ofwhich is the MC-NG alloy, is developed in France byONERA [8] and the alloys TMS75 and TMS80 weredeveloped in Japan [25]. The chemical compositionsof some of the third generation alloys are reported intable III.

The tendency to the precipitation of TCP phases in thethird generation alloys is a still more important problemthan in the second generation alloys since it is difficultto attain the right balance between the alloying elementspromoting this instability. Once again, the content ofCr was decreased as compared to the second generationalloys in order to keep the alloys less prone to TCP phaseprecipitation. The typical level of Cr encountered withinthe third generation alloys is now between 2 to 4.2 wt.%.

The role of Co on the precipitation of TCP phases isstill very controversial. Whereas Erickson limited the Colevel at 3 wt.% in CMSX-10 claiming that it reducesthe tendency to form TCP phases [17], Walston et al.recommended a high level of Co, 12.5 wt.%, in René N6in order to improve phase stability [39].

Re-rich TCP phase particles have however been re-ported to precipitate in CMSX-10 after high temperatureexposure, the greatest propensity for phase instability oc-


Table III. Chemical compositions (wt.%) of third generation Ni-based superalloys for single crystal blades.

Alloy Cr Co Mo Re W Al Ti Ta Nb Hf Others Density (g.cm−3) Country Ref.

CMSX-10 2 3 0.4 6 5 5.7 0.2 8 0.1 0.03 – 9.05 USA [17]

René N6 4.2 12.5 1.4 5.4 6 5.75 – 7.2 – 0.15 0.05C 8.97 USA [39]

0.004B

0.01Y

Alloy 5A 4.5 12.5 – 6.25 5.75 6.25 – 7 – 0.15 0.05C 8.91 USA [2]

0.004B

0.01Y

TMS-75 3 12 2 5 6 6 – 6 – 0.1 – – J [25]

TMS-80 2.9 11.6 1.9 4.9 5.8 5.8 – 5.8 – 0.1 3 Ir – J [25]

MC-NG Patent pending 8.75 F [8]

Figure 3. Interdendritic carbides and dendrite core Re-rich TCPphase precipitates within a René N6 single crystal fully heat-treated and aged for 200 hours at 1050◦C.

curring within the temperature range 1090–1150◦C [17,39]. A reduction of creep life associated with the precip-itation of these Re-rich TCP phase particles was recog-nised in CMSX-10 for long exposures between 1100 and1160◦C [17]. Re-rich phase precipitation has also beenreported in René N6 [39] but it did not detrimentally af-fect the creep rupture properties of this alloy. The pre-cipitation of Re-rich TCP phases is illustrated infigure 3showing the microstructure within a dendrite core of aRené N6 SC sample melted and cast at ONERA and ex-posed for 200 hours at 1050◦C.

However, a new type of instability was observed byWalston et al. [38] in superalloys containing high levelsof refractory elements and evaluated during the devel-opment phase of René N6. A typical alloy used in thisstudy was Alloy 5A [2] (seetable III). This instability,termed secondary reaction zone (SRZ), was observed be-

neath the diffusion zone of aluminide coatings, in den-drite cores and along some low angle boundaries (LAB’s)away from the coating. These so-called LAB’s delimitsub-grains containing groups of dendrites and with rel-ative misorientations up to 15◦.

These SRZs are described as areas with aγ ′ matrixcontainingγ and P phase needles and are referred toas cellular colonies when present in dendrite cores andsub-grain boundaries. The nucleation and growth of theSRZs is thought to be controlled by local chemicalsupersaturations resulting from the presence of a coatingor from internal segregation (to dendrite core or LBS’ssegregation) and by strain energy introduced by surfacepreparation prior coating, or by misfit strains alonggrain boundaries or betweenγ and γ ′. Rhenium hasbeen demonstrated to be the most potent element fordetermining the alloy propensity to form SRZ [38].

The main concern was thus to evaluate the eventualdeleterious effects resulting from the presence of SRZ be-neath the coating or away from the surface. Walston etal. demonstrated that SRZ beneath the coating can even-tually affect the rupture strength due to the reduction ofcross section [38]. The cellular colonies formed at the lowangle grain boundaries may also induce some reductionof the rupture properties for large misorientations. How-ever, the most detrimental effect was attributed to the cel-lular colonies in the dendrite cores. Some cracks form-ing at the SRZ interfaces indeed induce premature failureleading to a life reduction of up to 70%.

The development of René N6 was thus based upon thework performed on the SRZ containing alloys, with theprimary goal to obtain a microstructure stable with re-spect to sub-coating SRZ and dendritic cellular colonies.This goal was achieved essentially by decreasing the levelof Re and by introducing Mo in order to keep a compa-rable total amount of strengthening refractory elements.Walston et al. pointed out the difference existing betweenthese SRZs and the well known TCP phases [39]. They


suggest that the precipitation of TCP phases assisted byexcess concentrations in Cr and Mo will deplete the ma-trix from these and then reduce the chemical driving forcefor SRZ. They report the formation of SRZs in CMSX-10 beneath a PtAl coating after exposure for 400 hours at982◦C, but only TCP phase precipitation in failed creeprupture bare specimen of CMSX-10 [39].

It is noteworthy that the third generation SC superal-loys containing some additions of ruthenium [8,28] orIr [25] seem to show a reduced tendency to form unde-sirable TCP phases. This beneficial effect is presumablydue to the modification of the partitioning ratios of the el-ements between theγ andγ ′ phase, that will decrease theconcentration of Re in theγ matrix and thereby renderingit less prone to the precipitation of Re-rich TCP phases.These alloys exhibit excellent creep strength, especiallyat high temperatures, but complementary investigationsare needed to evaluate and understand the influence ofRu and Ir on other properties.

A significant difference in the design of the thirdgeneration alloys concerns the levels of minor elementssuch as C, B, Hf and Y. Whereas C and B are notintentionally added within the alloys CMSX-10, TMS-75and -80, Walston et al. [39] choose to reintroduce theseelements in René N6 in order to improve the castabilityand the tolerance to low angle grain boundaries.

The level of carbon was thus set at 0.05 wt.% in RenéN6 for a cleaner melting because it helps in reducingthe oxides, thus improving the castability. In order tominimise the possible deleterious effects of LABs whichare inevitably present in the “single crystal” components,low amounts of C, B and Hf are also added to RenéN6 in order to strengthen the LABs. These additionsimprove the yield of the components, because LABswith misorientations as large as 12◦ are accepted, insteadof the value of 6◦ generally fixed for rejection of thecastings when the grain boundary strengthening elementsare omitted [2].

Yttrium is also added in N6 to improve the adherenceof the Al2O3 protective layer formed at high temperaturewhereas Hf is very often added in the SC superalloys inorder to improve their coatability. Indeed, it was clearlydemonstrated that small additions of reactive elementssuch as yttrium, hafnium, cerium and lanthanum improvethe adherence of alumina scales on the superalloys [36].One of the mechanisms proposed to explain this benefi-cial effect is that these reactive elements tie up sulphur inthe alloy and prevent its segregation to the alloy/oxide in-terface [34]. Residual content of yttrium is however diffi-cult to control within the single crystal component due toreactions between the molten alloy and foundry ceram-ics and yttrium volatilisation during melting and direc-tional solidification [1]. Moreover, the segregation of yt-trium during the directional solidification may promotethe formation of low melting point phases in the inter-dendritic areas which will in turn reduce the tempera-ture of incipient melting as reported for the Y-enriched

Figure 4. Cyclic corrosion behaviour at 850◦C of CMSX-10,MC-NG and René N6 alloys.

version of CMSX-4 [26]. On the other hand, work per-formed at ONERA showed that simultaneous additionsof Si and Hf dramatically improve the cyclic oxidation inair at 1100◦C of the AM1, AM3 and MC2 single crystalalloys at a level comparable to that obtained with Y [7].These additions can be a promising alternative to Y, be-cause additions of Si and Hf are easy to control within thecast components.

It must be pointed out that the hot corrosion resistancewill be difficult to maintain for the low-chromium thirdgeneration SC superalloys because Cr plays a key role inpromoting the formation of a protective chromium richoxide scale at the temperatures where hot corrosion is ac-tive. On the basis of isothermal burner rig corrosion testsperformed at 950◦C with 2 ppm salt ingestion, Erick-son reported comparable corrosion resistance for CMSX-10 and CMSX-4 at least up to 100 hours [17]. How-ever, these corrosion test conditions are not very severeas compared to the real situation in the turbine where theblades are subjected to cyclic thermal stresses and wherehot corrosion damaging may occur at temperatures lowerthan 950◦C. Some cyclic corrosion tests were thus per-formed at ONERA in air at 850◦C with one-hour cyclesand additions of Na2SO4 renewed every 50 hours, on baresamples of René N6, CMSX-10 and MC-NG alloys. Thesalt level deposited on the samples corresponds to thattypically found in a burner rig using kerosene fuel con-taining 0.15% sulphur and a 5 ppm NaCl contamination.The specific weight changes are plotted as a function ofnumber of cycles infigure 4. The CMSX-10 alloy whichcontains only 2 wt.% Cr exhibits a very poor hot corro-sion resistance as compared to René N6 which contains4.2 wt.% Cr, the MC-NG alloy exhibiting an intermedi-ate behaviour. The strong difference in corrosion resis-tance between CMSX-10 and René N6 is perhaps not en-tirely due to the Cr content variation, but these tests showthat the low content of Cr in some third generation SCalloys may not be sufficient to obtain a satisfactory hotcorrosion strength in severe environment. Hot corrosiondip tests using Na2SO4-25% salt mixture performed at900◦C demonstrated that alloys TMS-75 and -80, con-


taining only 3 wt.% Cr, exhibit a corrosion resistanceslightly better than for CMSX-4 [25]. This result hashowever to be confirmed by testing under conditions pre-vailing in aeroengines.

The problem of the formation of grain defects, frecklesand spurious grains during directional solidification ofsingle crystal components, mentioned in the case ofsecondary generation alloys, becomes more acute for thethird generation alloys due to the higher content of Re.It renders in particular more difficult the growth of largeblades for industrial gas turbines [29].

One of the characteristics of the third generation SCalloys is that their high total amount of alloying ele-ments, in particular Ta and W, promotes the formationof a large volume fraction ofγ /γ ′ eutectic. A typicalexample is shown infigure 5 illustrating the as-cast mi-crostructure of a single crystal of the CMSX-10 alloymelted and cast at ONERA. Experience shows that itis impossible to eliminate these eutectic phases by us-ing a single isothermal solution heat treatment as in thecase of first generation SC alloys. It is necessary to ap-ply a complex heat treatment cycle generally including apre-homogenisation heat treatment aiming at increasingthe incipient melting temperature, before performing a fi-nal solution treatment which eliminates the major part orall of the eutectic phases. Complete elimination of theseγ /γ ′ nodules has been thus obtained in CMSX-10 by us-ing the following heat treatment procedure: 1337◦C/3h,ramp up at 3◦C.h−1, 1367◦C/3h/Air cooled (figure 5b).To attain this result, Erickson reported durations of solu-tion heat treatments of 30–35 h, with a peak soak temper-ature of about 1366◦C [15].

The choice to reintroduce minor elements such as C, Band Y which lower the incipient melting temperature inN6 render more difficult the elimination ofγ /γ ′ eutectic

and the solutioning of coarseγ ′ precipitates. This diffi-culty was circumvented by accepting some residual eu-tectic and incipient melting after solutioning the alloy. Inthese conditions, a 20◦C temperature window is claimedfor acceptable solution heat treatment, the optimum beingat temperatures in the range 1315–1335◦C for approxi-mately six hours [39]. The size and amount of such eutec-tic phases must however be carefully controlled becausethey can act as fracture initiation sites and thereby affectthe tensile ductility or the endurance limit in fatigue. It isalso noteworthy that the voluntary addition of carbon ledto the formation of interdendritic blocky (Ti, Ta)C car-bides as shown infigure 3illustrating the microstructureof a René N6 single crystal cast at ONERA, fully heat-treated and finally aged for 200 hours at 1050◦C. As forthe residual eutectic nodules, these carbides may penalisethe ductility and the fatigue properties by promoting earlyrupture.

As Re and W segregate strongly to the dendrites duringthe directional solidification, the benefit of a long solutionheat treatment is to improve the chemical homogeneityof the alloy and hence to prevent phase instabilitieswithin the dendrite cores. This beneficial effect hasbeen demonstrated in CMSX-10 and René N6 [15,39].The duration of the solution treatment must howeverbe optimised, by considering also the cost of this heattreatment.

The high level of refractory elements such as W, Reand Ta can also be a drawback for the third generationalloys, as they induce an increase of the density up tovalues around 9 g.cm−3 (table III). Moreover, high levelsof Re, Ta and eventually Ru and Ir, which are expensivemetals, increase the cost of the alloy, which could be adrawback for cost effective applications in civil engines.

(a) (b)

Figure 5. Transverse dendritic microstructure of a CMSX-10 single crystal; (a) as-cast; (b) solution heat-treated.


Figure 6.Typical temperature advantages over CC superalloys obtained with DS and SC superalloys estimated from stress rupture testsperformed at 982◦C and 248 MPa (from reference [16]).

3. Evolution of the mechanical properties of thesingle crystal superalloys

The constant creep strength improvements resultingfrom the transition from CC superalloys, to DS materials,including second generation DS alloys with 3 wt.%Re, and then to the three generations of SC alloys areillustrated infigure 6where the temperature advantagesare reported for stress rupture tests performed at 982◦Cand 248 MPa.

The major objective of the alloy designers in develop-ing third generation SC superalloys was to improve sig-nificantly the stress rupture strength of this class of al-loys compared to previous first and second generation al-loys. Data on René N5, René N6 [39] and CMSX-10 [9,25] alloys are compared to data produced at ONERA onMC-NG alloy infigure 7where the specific stress-rupturestrengths are reported in a Larson Miller diagram.

The third generation SC alloys exhibit comparablestress rupture strengths significantly higher than for thesecond generation alloy René N5. Thus, Walston etal. [39] claim for a 30◦C benefit obtained with René N6 ascompared to René N5. It is noteworthy that the MC-NGalloy shows a very high creep resistance with a densitysignificantly lower than those of CMSX-10 and René N6,which is a great advantage because it helps to reduce thestresses on the disk, and the weight of the rotor.

The temperature advantage of CMSX-10 over CMSX-4 appears to be equal to 36◦C in the temperature range913–1010◦C. However, the CMSX-10 creep strengthbenefit at 982◦C as compared to CMSX-4 decreases withincreasing exposure times [17], which suggests that theremay be a point beyond which the third generation al-loys would not really be better than the second gener-ation alloys. This analysis was confirmed by Ross andO’Hara [33] who compared the stress rupture life at

Figure 7. Comparative Larson–Miller stress-rupture curves forsecond and third generation SC superalloys.

982◦C of different generations of SC superalloys. It ap-pears that the first generation SC René N4 surpasses thesecond generation alloys in the 7000–10000 hour liferegime at this temperature and data extrapolation indi-cates that this alloy could attain the level of the third gen-eration alloys for rupture life around 20000 hours. On theother hand, some comparative creep tests performed atONERA on various single crystal alloys demonstrate thata Re-free superalloy such as MC2 exhibits stress rupturelives at 1050◦C comparable to that of superalloys con-taining Re contents close to 6 wt.%.

A great advantage of the third generation SC superal-loys is that they maintain a rather high creep resistance attemperature above 1100◦C. Indeed, the stress rupture lifeat 1150◦C and 100 MPa of the MC-NG alloy is over 150hours, whereas the first generation alloys show, typically,a rupture life less than 10 hours.

Contrary to the creep strength, tensile properties ofthe single crystal superalloys are not reported to be


very sensitive to their chemical compositions. The tensilebehaviour of the third generation alloys is thereforecomparable to that of first and second generation alloyswith a peak strength at about 760◦C and a good ductilitywithin the temperature range of interest [17]. However,some significant improvements in low cycle fatigue(LCF) and high cycle fatigue performance have beenreported both for CMSX-10 [16] and René N6 [39] alloysas compared to their predecessors containing 3 wt.%Re. The significant increase of LCF properties observedat temperatures around 950◦C are of prime importancein the case of cooled blades or vanes. Indeed, the thinairfoil walls are subjected to huge thermal strains andstresses due to the thermal gradients occurring withinthe blade during the working cycles of the gas turbineengines, which lead to damaging crack initiation andgrowth processes. It is therefore thought that the fatiguestrength benefit afforded by the third generation alloyscould improve significantly the durability of the singlecrystal components.

4. Conclusions

Significant progress in the improvement of the perfor-mance of the nickel based single crystal superalloys hasbeen made with the optimisation of the chemical com-positions. The addition of rhenium up to 6 wt.% in thethird generation alloys was a key factor for the increaseof the mechanical properties. The development of suchalloys has however taken into account some specific fea-tures such as increase in cost and density, tendency tophase instability and grain defects, which must be care-fully controlled in order to ensure that these alloys willbe effectively exploitable as turbine blade materials. Re-cent development at ONERA of new generation singlecrystal superalloys with additions of both rhenium andruthenium could pave the way for the development of im-proved third generation superalloys with reduced densityand better phase stability compared to third generationalloys containing high levels of rhenium.

Acknowledgements

The authors acknowledge the French Ministry ofDefence for the financial support of a part of the workpresented in this paper and would like to thankJ.-L. Raffestin for single crystal preparation.

References

[1] Aimone P.R., McCormick R.L., The effect of yttriumand sulfur on the oxidation resistance of an advancedsingle crystal nickel based superalloy, in: Antolovich S.D.et al. (Eds.), Superalloys 1992, The Minerals, Metals &Materials Society, Warrendale, PA, USA, 1992, pp. 817–823.

[2] Austin C.M., Darolia R., O’Hara K.S., Ross E., GeneralElectric Company, US Patent 5 151 249, 1992.

[3] Blavette D., Caron P., Khan T., An atom probe investiga-tion of the role of rhenium additions in improving creepresistance of Ni-base superalloys, Scripta Metall. Mater.20 (1986) 1395–1400.

[4] Blavette D., Caron P., Khan T., An atom-probe study ofsome fine-scale microstructural features in nickel-basedsingle crystal superalloys, in: Reichman S. et al. (Eds.),Superalloys 1988, The Metallurgical Society, Warrendale,PA, USA, 1988, pp. 305–314.

[5] Caron P., Khan T., Improvement of creep strength in anickel-base single-crystal superalloy by heat-treatment,Mat. Sci. Eng. 61 (1983) 173–194.

[6] Caron P., Khan T., Development of a new nickel basedsingle crystal turbine blade alloy for very high temper-atures, in: Exner H.E., Schumacher V. (Eds.), AdvancedMaterials and Processes, Vol. 1, DGM Informationsge-sellschaft mbH, Oberursel, Germany, 1990, pp. 333–338.

[7] Caron P., Navéos S., Khan T., Improvement of the cyclicoxidation behaviour of uncoated nickel based singlecrystal superalloys, in: Coutsouradis D. et al. (Eds.),Materials for Advanced Power Engineering 1994 – PartI, Kluwer Academic Publisher, Dordrecht, Holland,1994,pp. 1185–1194.

[8] Caron P., Raffestin J.-L., Navéos S., ONERA, Patentpending.

[9] Caron P., ONERA, unpublished work.[10] Cetel A.D., Duhl D.N., Second-generation nickel-base

single crystal superalloy, in: Duhl D.N. et al. (Eds.), Su-peralloys 1988, The Metallurgical Society, Inc., Warren-dale, PA, USA, 1988, pp. 235–244.

[11] Darolia R., Lahrman D.F., Field R.D., Sisson R., Forma-tion of topologically closed packed phases in nickel basesingle crystal superalloys, in: Duhl D.N. et al. (Eds.), Su-peralloys 1988, The Metallurgical Society, Inc., Warren-dale, PA, USA, 1988, pp. 255–264.

[12] Davidson J.H., Fredholm A., Khan T., Théret J.-M.,French patent N◦2 557 598, 1983.

[13] Dreshfield R.L., Ashbrook R.L., Effects of sigma-phaseformation on some mechanical properties of wroughtnickel-base superalloys (IN 100), NASA TN D-7654,1974.

[14] Duhl D.N., Alloy phase stability requirements in singlecrystal superalloys, in: Stocks G.M. (Ed.), Alloy PhaseStability and Design, Mat. Res. Soc. Symp. Proceedings,Vol. 186, Materials Research Society, 1991, pp. 389–399.

[15] Erickson G.L., A new, third-generation, single-crystal,casting superalloy, J. Metals 47 (4) (1995) 36–39.

[16] Erickson G.L., The development of CMSX-10, a thirdgeneration SX casting superalloy, in: Proceedings of theSecond Pacific Rim International Conference on Ad-vanced Materials and Processing (PRCIM-2), Kyongju,Korea, 18–22 June 1995.

[17] Erickson G.L., The development and application ofCMSX-10, in: Kissinger R.D. et al. (Eds.), Superalloys1996, The Minerals, Metals & Materials Society, Warren-dale, PA, USA, 1996, pp. 35–44.

[18] Ford D.A., Arthey R.P., Development of single crystalalloys for specific engine applications, in: Gell M. et al.


(Eds.), Superalloys 1984, The Metallurgical Society ofAIME, Warrendale, PA, USA, 1984, pp. 115–124.

[19] Giamei A.F., Anton D.L., Rhenium additions to a Ni-base superalloy: effects on microstructure, Metall. Trans.A 16A (1985) 1997.

[20] Gell M., Duhl D.N., Giamei A.F., The development ofsingle crystal superalloy turbine blades, in: Tien J.K. et al.(Eds.), Superalloys 1980, American Society for Metals,Metals Park, OH, USA, 1980, pp. 205–214.

[21] Harris K., Erickson G.L., Schwer R.E., Mar M 247derivations – CM 247 LC DS alloy, CMSX single crystalalloys, properties and performance, in: Gell M. et al.(Eds.), Superalloys 1984, The Metallurgical Society ofAIME, Warrendale, PA, USA, 1984, pp. 221–230.

[22] Harris K., Erickson G.L., Schwer R.E., Frasier D.J.,Whetstone J.R., Process and alloy optimization forCMSX-4 superalloy single crystal airfoils, in: Bachelet E.et al. (Eds.), High Temperature Materials for Power Engi-neering 1990, Part II, Kluwer Academic Publishers, Dor-drecht, Holland, 1990, pp. 1281–1300.

[23] Harris K., Erickson G.L., Sikkenga, S.L., Brentnall W.D.,Aurrecoechea J.M., Kubarych K.G., Develoment of therhenium containing superalloys CMSX-4 and CM 186LC for single crystal blade and directionally solidifiedblade vane applications in advanced turbine engines, in:Antolovich S.D. et al. (Eds.), Superalloys 1992, TheMinerals, Metals & Materials Society, Warrendale, PA,USA, 1992, pp. 297–306.

[24] Khan T., Recent developments and potential of singlecrystal superalloys for advanced turbine blades, in: BetzW. et al. (Eds.), High Temperature Alloys for Gas Tur-bines and Other Applications 1986, D. Reidel PublishingCompany, Dordrecht, Holland, 1986, pp. 21–50.

[25] Kobayashi T., Koizumi Y., Nakazawa S., Yamagata T.,Harada H., Design of high rhenium containing singlecrystal superalloys with balanced intermediate and hightemperature creep strengths, in: Strang A. et al. (Eds.),Advances in Turbine Materials, Design and Manufactur-ing, The Institute of Materials, London, UK, 1997, pp.766–773.

[26] Marchionni M., Goldschmidt D., Maldini M., High tem-perature mechanical properties of CMSX-4+Yttriumsingle-crystal nickel-base superalloy, in: Antolovich S.D.et al. (Eds.), Superalloys 1992, The Minerals, Metals &Materials Society, Warrendale, PA, USA, 1992, pp. 775–784.

[27] Morinaga M., Yukawa N., Adachi H., Ezaki H., Newphacomp and its applications to alloy design, in: Gell M.et al. (Eds.), Superalloys 1984, The Metallurgical Societyof AIME, Warrendale, PA, USA, 1984, pp. 523–532.

[28] O’Hara K.S., Walston W.S., Ross E.W., Darolia R.,General Electric Company, US Patent 5 482 789, 1996.

[29] Parks Jr. W.P., Hoffman P., Karnitz M.A., Wright I.G.,The advanced turbine systems program in the USA, in:Lecomte-Beckers J. et al. (Eds.), Materials for AdvancedPower Engineering 1998 – Part III, ForschungzentrumJülich, Jülich, Germany, 1998, pp. 1789–1805.

[30] Pessah M., Caron P., Khan T., Effect ofµ phase on themechanical properties of a nickel-base single crystal su-

peralloy, in: Antolovich S.D. et al. (Eds.), Superalloys1992, The Minerals, Metals & Materials Society, Warren-dale, PA, USA, 1992, pp. 567–576.

[31] Pollock T.M., The growth and elevated temperature sat-bility of high refractory nickel-base single crystals, Mater.Sci. Eng. B32 (1995) 255–266.

[32] Ross E.W., Sims, C.T., Nickel-base superalloys, in: SimsC.T., Stoloff N., Hagel W.C. (Eds.), Superalloys II, JohnWiley & Sons, New York, USA, 1987, pp. 97–133.

[33] Ross E.W., O’Hara K.S., René N4: a first generation sin-gle crystal turbine airfoil alloy with improved oxidationresistance, low angle boundary strength and superior longtime rupture strength, in: Kissinger R.D. et al. (Eds.), Su-peralloys 1996, The Minerals, Metals & Materials Soci-ety, Warrendale, PA, USA, 1996, pp. 19–25.

[34] Sarioglu C., Stinner C., Blachere J.R., Birks N., PettitF.S., Meier G.H., Smialek J.L., The control of sulfurcontent in nickel-base, single crystal superalloys and itseffect on cyclic oxidation resistance, in: Kissinger R.D.et al. (Eds.), Superalloys 1996, The Minerals, Metals &Materials Society, Warrendale, PA, USA, 1996, pp. 71–80.

[35] Sims, C.T., Prediction of phase composition, in: SimsC.T., Stoloff N., Hagel W.C. (Eds.), Superalloys II, JohnWiley & Sons, New York, USA, 1987, pp. 217–240.

[36] Smialek J.L., Meier G.H., High-temperature oxidation, in:Sims C.T., Stoloff N., Hagel W.C. (Eds.), Superalloys II,John Wiley & Sons, New York, USA, 1987, pp. 293–326.

[37] Strangman T.E., Hoppin III G.S., Phipps C.M., HarrisK., Schwer R.E., Development of exothermically castsingle crystal Mar-M 247 and derivative alloys, in: TienJ.K. et al. (Eds.), Superalloys 1980, American Society forMetals, Metals Park, OH, USA, 1980, pp. 225–234.

[38] Walston W.S., Schaeffer J.C., Murphy W.H., A newtype of microstructural instability in superalloys – SRZ,in: Kissinger R.D. et al. (Eds.), Superalloys 1996, TheMinerals, Metals & Materials Society, Warrendale, PA,USA, 1996, pp. 9–18.

[39] Walston W.S., O’Hara K.S., Ross E.W., Pollock T.M.,Murphy W.H., René N6: third generation single crystalsuperalloy, in: Kissinger R.D. et al. (Eds.), Superalloys1996, The Minerals, Metals & Materials Society, Warren-dale, PA, USA, 1996, pp. 27–34.

[40] Waudby P.E., Benson J.M., Stander C.M., PennefatherR., McColvin G., Advanced high strength single crystalsuperalloy – SMP14, in: Strang A. et al. (Eds.), Advancesin Turbine Materials, Design and Manufacturing, TheInstitute of Materials, London, UK, 1997, pp. 322–338.

[41] Wortmann J., Wege R., Harris K., Erickson G.L., SchwerR.E., in: Proceedings 7th World Conference on Invest-ment Casting, Münich, Germany, 29 June–2 July 1988.

[42] Wukusick C.S., Final Report NAVAIR/N62269–78-C-0315, 25 August 1980.

[43] Wukusick C.S., Buchakjian, Jr. L., General Electric Com-pany, Patent Application #GB 2 235 697 A, March 1991.

[44] Xuan Nguyen-Dinh, Allied-Signal, Inc., US patent 4 935072, 1990.

super alloys development

Documents