Creep Resistant Superalloys

Creep Resistant Superalloys

Creep Resistant Superalloys The development of nickel-base superalloys has been closely related to the evolution of the gas turbine, particularly for ...

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Creep Resistant Superalloys The development of nickel-base superalloys has been closely related to the evolution of the gas turbine, particularly for propulsion of aircraft (Sims et al. 1987). Indeed, it is arguable that the progressive increase in efficiency of gas turbines over a period of 60 years could not have occurred without the parallel dramatic improvement in the creep resistance and stability of these materials, albeit with very significant developments in engineering technology. Nickel-base superalloys have become the materials of choice for two critical areas of the gas turbine. In the turbine, rotating blades experience a combination of high temperatures produced by the gases exiting the combustion chamber and high centrifugal stresses. Stator blades (nozzle guide vanes) are subject to higher temperatures and gas bending stresses. The efficiency of the gas turbine is a sensitive function of temperature of the gas entering the turbine. The superalloys intended for use as turbine blades have shown a progressive increase of about 6 K per annum in the maximum temperatures at which significant stresses can be tolerated (i.e., in creep resistance) over a 60-year period. This trend is illustrated in Fig. 1 that plots the temperature at which a stress of 150 MPa gives a creep life of 1000 hours as a function of year of introduction of alloys. Figure 1 also indicates how the processing route of the alloys has evolved. The turbine disks to which the blades are attached are subject to much higher centrifugal stresses and lower temperatures than are the blades. Consequently creep deformation and fracture is less of a problem. Rather, the imperative of low engine weight has

Figure 1 Trend in improvement in performance of nickel-base superalloys as a function of their year of introduction: blade alloy temperature for 1000 hours rupture life with a stress of 150 MPa.

required the cross-section of the disk to be minimized, and this has focused attention on the yield strength, rather than creep strength; there has been a progressive increase in yield strength with year for disk alloys. As the yield strength has increased, the fracture toughness has decreased and the materials have become sensitive to defects requiring the development of sophisticated clean-melting procedures. These alloys will not be discussed further in this article. The metal temperatures experienced by modern gas turbine blades ($ 1350 K) approaches 0.8 of the alloy melting point. The retention of creep resistance to such high temperatures has been achieved by a combination of solid solution and precipitation strengthening and by control of the grain size, morphology and orientation. However, by far the most important characteristic of nickel-base superalloys is the ordered Ni (Al,Ti) γh intermetallic phase that $ forms in the disordered γ-nickel matrix by solid state precipitation. The dramatic improvements in creep strength and temperature capability of superalloys have been a result of progressive increases in the volume fraction and solvus temperature of the γh phase. This evolution has had an important influence on the processing of the materials to produce engineering components. In the early blade alloys (e.g., Nimonic 80A, IN115), there was a wide temperature range between the γh solvus and the melting temperature which allowed the alloys to be forged while in the relatively soft, singlephase γ form. From about 1960, alloys (e.g., IN100, MarM200, and IN738) were produced in which the temperature range between γh solution and incipient melting was too small to constitute a practical temperature window for thermal–mechanical processing. Investment casting of turbine blades from these alloys produced components with enhanced creep resistance, but the creep ductility was significantly lower than for the forged components (2–5% compared to 10–15%). This became a serious problem when higher strength alloys were developed where ductilities of 1% were associated with fracture at grain boundaries normal to the direction of maximum tensile stress. This gave problems of thermal fatigue, which is particularly serious for aircraft applications, where the flight cycles, and associated stresses\temperatures experienced by the blades in service, can be very variable. Directional solidification (DS) was developed by Ver Snyder and Guard (1960) to produce turbine blades with columnar grains and few transverse grain boundaries. Creep ductilities of 25–30% were routinely produced in the same alloys that were creep-brittle in the conventionally cast form. In addition, directional solidification produced a sharp f001g crystallographic texture parallel to the axis of the blade. f001g is the direction of minimum Young’s modulus in cubic crystals, and this reduced the thermal stresses allowing the intrinsic creep strength of the superalloys to be 1

Creep Resistant Superalloys more effectively exploited. Few alloys were developed specifically for use in the directionally solidified form, although compositional modifications were introduced, such as small additions of hafnium to suppress the cracking of longitudinal grain boundaries due to solidification stresses. Single crystal superalloys (SX) had been produced with minor modification of the directional solidification process in the 1960s. However, there was no significant advantage of single crystals over DS superalloys for the same alloy composition. Consequently this technology was dormant for some 20 years. From about 1980 (Gell et al. 1980), a series of alloys (CMSX2, SRR99) have been developed for use specifically in the single crystal form that have up to a 20 K advantage over the best DS alloys. The first generation of SX superalloys simplified the alloy compositions by removing the elements intended to produce grain boundary strengthening, which were deemed to be redundant. This increased the solidus temperature, giving an effective heat treatment window for total dissolution of the γh and subsequent control of its morphology. The second generation of SX superalloys (CMSX4, MC2) further increased both the γh volume fraction and the concentration of refractory solid–solution strengthening elements. In particular, addition of about 1% of rhenium has been found to have a significant effect in reducing the creep rate and in suppressing morphological changes in the γh precipitate; however, it also leads to increased microsegregation during solidification and suppresses compositional homogenization during heat treatment. A third generation of SX superalloys (e.g., CMSX10) with higher levels of rhenium and ruthenium have been proposed; these have excellent stress rupture properties, but may be susceptible to phase instabilities (e.g., TCP). All of these types of alloys are currently used in different stages of modern gas turbines. However, since SX superalloys are the current state-of-the-art the remainder of this article will concentrate on their creep properties.

1. Microstructure and Mechanisms The high level of creep resistance and excellent thermal stability of nickel-base superalloys is a consequence of crystallographic similarities between the γ-matrix and the γh-precipitate. The γh has an f.c.c. L1 ordered lattice with similar lattice parameter to the γ #matrix. It precipitates from the solid state, coherent with the matrix, the small lattice parameter mismatch leading to elastic internal stresses, which are relaxed during subsequent plastic deformation by deposition of dislocations at the γ–γh interfaces. The volume fraction of γh in superalloys has increased from $ 20% in the earliest precipitation strengthened alloys (e.g., 2

Nimonic 80A) to $ 70% in the most advanced SX superalloys (e.g., CMSX4). Because of the ordering of the γh lattice, a dislocation with unit Burgers vector in the γ matrix would disturb the crystal symmetry in passing through the γh; it would create an antiphase boundary which would be eliminated and ordering restored by the passage of a second unit dislocation. There is an energy penalty in creating and propagating such dislocation pairs that can cut through the duplex microstructure. The nature of creep deformation in superalloys is very sensitive to temperature, stress and (for SX alloys) crystal orientation. Figure 2 shows typical dislocation distributions observed in single crystal superalloys after creep deformation at various temperatures. At temperatures of 750 mC and below where stresses are high, there is evidence of extensive cutting of the γh. This can lead to fairly homogeneous deformation, particularly for tensile stresses close to f001g, or to deformation twinning for other orientations. In both cases the dislocations in the γh have (111)f211g Burgers vectors. It appears that the deformation associated with γh cutting, which being in the majority phase is very obvious in microscopic examination, is additional to that observed at intermediate temperatures. Ardakani et al. (1999) have shown that the deformation resulting from deformation twinning constitutes about 10% of the observed creep strain. At intermediate temperatures dislocation activity is restricted to the minority γ matrix and has a (111)f110g character. The dislocations bow in the γ channels initially transferring load to γh particles, to which the dislocations are attached, leading to a decreasing creep rate (i.e., primary creep). Deformation is controlled by the rate of climb of the pinned dislocations along the γ–γh interfaces which are (001) planes; the dislocation density increases with increasing plastic strain leading to an increased creep rate (tertiary creep). At high temperatures (1000 mC and above) or at long times at lower temperatures, the γh morphology changes radically from the regular cuboidal distribution to extensive rafts oriented normal to the predominant tensile stresses in the negative mismatch alloys which constitute most current SX alloys. Since the stresses are too low to cause γh-cutting, dislocations are constrained to the extensive γ channels and this is believed to give anomalously low creep rates. Fracture occurs when dislocation pile-ups can cause cutting of the γh. A number of studies has shown that this precipitate rafting reduces the creep strength at intermediate temperatures.

2. Creep Behavior and Constitutive Equations Creep curves of nickel-base superalloys (equiaxed, columnar grained, and single crystal) all show a wide range of shapes for different stresses and temperatures

Creep Resistant Superalloys

g

(a) g

(b)

g

) Figure 2 Transmission electron micrographs of SRR99 after creep testing: (a) f011g at 750 mC\850 MPa showing microtwins; (b) f001g at 950 mC\175 MPa showing deformation restricted to γ; and (c) f001g at 1050 mC\150 MPa showing rafts and dislocation pairs in γh.

as indicated in Fig. 3. There are at least two important features of this creep behavior:

There is no evidence of a significant regime of steady state creep; rather, there is a transition from primary to tertiary creep, with a minimum creep rate occurring at the point of inflection. Primary creep strain is generally greatest in high stress\low temperature tests consistent with load being transferred from the plastically deforming γ to the elastic γh phase. Primary creep strain progressively decreases as the test temperature is increased and the applied stress is decreased. These creep curves are dominated by a progressively increasing creep rate (tertiary) that is indicative of a steady reduction in creep strength. When the minimum creep rate is represented by the Bailey–Norton power law equation, both the stress exponent and the activation energy are much greater than can be justified by current models. This suggests that these materials do not deform by a recovery controlled creep mechanism; rather the deformation is controlled by reaction rate kinetics in which the creep rate as a function of stress is well described by a hyperbolic sine function (often approximated as an exponential function). This behavior is consistent with creep deformation occurring by the viscous motion of dislocations through the minority γ phase and by the progressive increase in dislocation density with creep strain. In these materials, all of the dislocations are mobile. The low dislocation densities mean that the back stress associated with other dislocations is relatively small compared to the barrier to motion presented by the precipitate, and the growing density of dislocations enhances the creep rate. Crystallographic anisotropy presents an additional complexity in single crystal superalloys and a number of anisotropic creep models has been proposed to account for this aspect. These range from the empirical which provide a fit of creep data to a mathematical function having the required crystal symmetry, to approaches which attempt to capture elements of the micromechanisms that are known to operate during high temperature deformation of these materials. For example, Meric et al. (1991) have extended an isotropic visco-plastic model incorporating isotropic and kinematic hardening terms by restricting deformation to specific slip systems. Ghosh et al. (1990) have also proposed a crystallographic-based model of anisotropic creep, later extended by Pan et al. (1997), that uses the formalism of continuum damage mechanics to incorporate features of both primary and tertiary creep. This is an extension of a previous isotropic model of creep that has been successful in accounting for the creep behavior of a range of materials. The minimum requirement of such models is that they should be capable of representing adequately the data from which they have been calibrated. It is unrealistic to expect individual model-generated creep curves to provide a better fit to the experimental creep curves than the intrinsic scatter of the input data. It 3

Creep Resistant Superalloys

700 MPa/750 °C

is of the same order as the intrinsic scatter in the experimental data. Such models, quantified using a database of uniaxial creep tests for simple orientations (e.g., f001g and f111g), have been extended to predict: (i)creepbehavior forarbitraryorientations including strain anisotropy and the associated changes in crystal orientation and specimen shape; (ii) accumulation of strain when the stress and\or temperature varies; (iii)theeffectofcreepinarangeofloadingconditionsincluding stress- or strain-control in monotonic or cyclic loading and stress relaxation; and (iv) the effects of multiaxial stresses.

3. Future Prospects

Figure 3 Typical uniaxial constant stress creep curves for SR99 for different creep test conditions.

Figure 4 Comparison of measured and model-predicted times to achieve 5% creep elongation in CMSX4 at various conditions of stress and temperature.

is more important to test the accuracy of the model predictions over the entire range of temperature and stress for which data are available. Figure 4 compares the model predictions and measurements of times to achieve 5% strain for all the creep conditions for which data are available for CMSX4; the agreement between model representation and experimental measurement 4

The progress in the development of nickel-base superalloys over a period of 60 years or so has been quite remarkable. However, the scope for radical enhancement of the creep resistance of these materials is essentially constrained by the melting point of nickel. The major thrusts in alloy development in recent years that are likely to continue have been on two fronts: Incremental improvements in creep strength have been achieved by solid solution strengthening by a range of refractory additions (e.g., rhenium, ruthenium). The long-term stability of some of the more highly alloyed materials has caused concern, and sophisticated thermodynamic models are being developed to aid alloy development. The potential for intermetallic nickel-base compounds (e.g., NiAl, Ni Al) to provide a step change in creep behavior has not$been fulfilled to date. However, there remains a high level of research activity aimed at solving the problem of brittleness of these materials in order to take advantage of their low densities and relatively high melting points. In the short to medium term, superalloys will remain the material of choice for gas-turbine applications. Effective exploitation will require significant development in two areas: Microsegregation in the modern highly alloyed single crystal superalloys can cause problems in processing. These are likely to become increasingly serious for castings for large industrial gas turbines. A greater understanding of the solidification processing leading to improved yield and quality control is required. The full advantages of single crystal superalloys are probably not yet being exploited because of incomplete appreciation of the anisotropic nature of creep deformation and of how to integrate this knowledge in engineering design. Effective constitutive equations extending steady uniaxial creep to variable and multiaxial loading and compatibility with finite element design methods are required.

Creep Resistant Superalloys See also: Creep Behavior of Materials: A Comparison; Creep-resistant Materials for Steam Turbines; Fatigue and Thermomechanical Fatigue at High Temperature

Bibliography Ardakani M G, McLean M, Shollock B A 1999 Acta Mater. 47, 2593–602 Gell M, Giamei A F, Duhl D N 1980 In: Tien J K (ed.)

Superalloys 1980. ASM, Metals Park, OH, USA Ghosh R N, Curtis R V, McLean M 1990 Acta Metall. Mater. 38, 1977 Pan L-M, Shollock B A, McLean M 1997 Proc. R. Soc. London 453, 1689–715 Meric L, Poubanne P, Cailletaud G 1991 Trans. ASME J. Eng. Mater. Tech. 113, 162–70 Sims C T, Hagel W, Stoloff N S (eds.) 1987 Superalloys II. Wiley, New York Ver Snyder F L, Guard R W 1960 Trans. ASM 52, 485

M. McLean

Copyright ' 2001 Elsevier Science Ltd. All rights reserved. No part of this publication may be reproduced, stored in any retrieval system or transmitted in any form or by any means : electronic, electrostatic, magnetic tape, mechanical, photocopying, recording or otherwise, without permission in writing from the publishers. Encyclopedia of Materials : Science and Technology ISBN: 0-08-0431526 pp. 1845–1849 5