gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

RECENT TRENDS IN SUPERALLOYS RESEARCH FOR CRITICAL AERO-ENGINE COMPONENTS

Luc Remy*, Jean-Yves Guedou**

*Centre des Materiaux, Mines ParisTech, CNRS UMR 7633, B.P. 87, 91003 Evry Cedex, France ** Materials and Processes Department, Snecma, Safran Group, 77550 Moissy-Cramayel, France

Abstract

This paper is a brief survey of common research activity on superalloys for aero-engines between Snecma and Mines ParisTecb Centre des Materiaux during recent years. First in disks applications, the development of new powder metallurgy superalloys is shown. Then grain boundary engineering is investigated in a wrought superalloy. Secondly, design oriented research on single crystals blades is shown: a damage model for low cycle fatigue is used for life prediction when cracks initiated at casting pores. The methodology developed for assessing coating life is illustrated for thermal barrier coating deposited on AM I single crystal superalloy.

Keywords: Nickel base superalloys, microstructure, low cycle fatigue, thermal barrier coatings, damage model

1. Introduction

Turboengines for aircrafts represent a major challenge for mechanics and materials (Fig. 1). Major requirements are resistance to long duration operation under severe environmental and thermal - mechanical loading, lightness and reliability. Aero-engine manufacturers have therefore to develop high tech engineering methods, materials and lifing techniques. There is a continued improvement of aircraft engine performance with an increase of turbine entry temperature and pressure ratio. During the last three decades, the increase of Turbine Entry Temperature has been was approximately 15°C per year. This increase in performance is heavily relying on disks and blades in the high pressure turbine that are the most critical parts (Fig. 1). These components are made of nickel base superalloy and the increase in performance has been achieved first through material improvement. These alloys are strengthened by a high volume fraction of"( Nh (fi, AI) precipitates and the continuous trend has been to increase the volume fraction of strengthening precipitates. The content can reach about 0.50-0.55 in powder metallurgy alloys for discs and about 0.7 in cast directionally solidified single crystals for blades. A long cooperation has been developed between SNECMA and Centre des Materiaux, in particular in superalloy research for disks and blades. This has involved alloy development with powder metallurgy alloys for disks, and single crystals for blades, investigation of damage mechanisms in creep, low cycle fatigue (LCF) and thermal-mechanical fatigue (TMF), assessment of constitutive models and damage models over 30 years. This has involved interactions with mechanical engineering groups at SNECMA, ONERA and Mines ParisTech and the materials science group at ONERA.

596 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck a.nd B. Kuhn.

SaM146 Figure 1: Aero-engines and a rotor with disk and blades made of superalloys.

The present paper gives a brief survey of conunon research activity along these lines during recent years. Alloy development and optimization of microstructures will be first shown with the development of improved PM superalloys and the investigation of grain boundary engineering in alloys for disks. Design-oriented research will be then illustrated with the development and assessment of life prediction models for single crystals blades: low cycle fatigue life prediction for the substrate with cracks initiating at casting pores and prediction of spalling of thermal barrier coatings that are now deposited on blades in advanced high pressure stages.

2. Alloy development and optimisation of microstructnres.

2.1. Powder Metallurgy superalloy for disks.

The decrease of NOx and C02 emissions and noise reductions on the short term, and increasing fuel saving levels for optimized acquisition and life cycle costs on the long term are major targets in modem gas turbine engines. These challenges are a major drift for higher overall pressure ratios and turbine entry temperatures. This results in higher temperatures and loadings for long durations for disks in high pressure compressors and turbines.

Today such parts are made of a powder metallurgy (PM) superalloy N18 [1]. This alloy has a high volume fraction of"( precipitates (55%) and hence it has a high "( solvus temperature (ST) and a very narrow solution heat treatment window. In addition it is sensitive to quench cracking and prone to TCP (topologically close-packed) phase precipitation during long time exposure at 650°C.

Therefore a collaborative program was undertaken between Snecma, ONERA and Centre des Materiaux to develop a new PM superalloy with the following specifications: capability of supersolvus solutioning which implies a reduced "( content, higher creep and fatigue resistance with respect to N18 up to 700°C, increased strain hardening, density lower than 8.35 Kg.dm-3 [2]. The "( content decrease has therefore to be counterbalanced by the improvement of strengthening of both matrix and precipitate. A careful balance between Cr, Mo and W was optimized to achieve matrix strengthening and avoiding TCP phases, using the new PHACOMP method; the Co content was used mainly to decrease the solvus temperature. Then in order to strengthen the"( phase, the ratio (Ti+Nb+Ta)/AI (in at%) was increased from 0.57 in alloy N18 to about 1. The content of minor elements was kept in the

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range 150-320 wt ppm for C, 150-200 wt ppm for C, 0.3 wt% for Hf and 600 wt ppm for Zr when added [2].

A total of 22 experimental heats were processed with alloys Nl8 and Rem~ 88. Electrodes of each grade were machined from vacuum induction melted ingots of 3.2 kg. Pre-alloyed powders were produced by the rotating electrode process (REP). The phase transformation temperature (solvus, incipient melting) were measured through differential thermal analysis (DTA) analysis. Powders were extruded in subsolvus conditions (ST-25°C). Screening tests consisted in metallography, tensile tests and creep tests. From 22 alloys, 5 remain after these tests and two of these were chosen for more detailed analysis [2].

Alloy Ni Co Cr Mo w A1 Ti Nb Hf B c Zr Nl8REP Bal. 14.9 12.4 3.8 - 9.1 5.1 - 0.13 0.09 0.07 0.02 SM048 Bal. 14.9 12.3 3.6 4.0 3.2 4.4 0.8 0.3 0.01 0.03 - SM043 Bal. 12.2 13.3 4.6 3.0 2.9 3.6 1.5 0.25 0.01 0.015 0.05 Table 1. Chemical composition ofalloys (wt%)

Two alloys were selected SM043 and SM048, and provided by Aubert & Duval (composition is given in Table 1). Powders were produced using an argon atomization industrial facility (Aubert & Duval). They were then processed using the industrial route as for Nl8 PM alloy: sieving (<53 Jl.lll), container filling and hot isostatic pressing (1150°C). Two bars were extruded at Snecma quite below the solvus temperature. The amount of"( phase was computed as 43% and 48% in alloys SM043 and SM048 respectively. Alloy density was 8.34 and 8.31 Kg.dm·3 respectively (instead of 8.00 for Nl8). Complete solutioning was possible due to the decrease in solvus temperature.

stress MPa 10000 f--- • SM043-&so·c f--- A. SM043-7So•c • SM048-650"C r--- e SM048-750"C N18-650"C

1000

- r- .. -..::: 1--.

100 10 100 1000 10000 Time,h

Figure 2: Creep times to 0.2% strain for new PM superalloys compared with alloy N18.

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~ Quasi-cleavage Pseudo-stress amplitude, MPa ,. 1-400 r-+-il-+tttt11r-+-t-t+H+Hr---i SM043 SM048 N18 1200 +--+-t-ffittt+l--t-+i-+tH+t----i I • 100) +-"""tod-ffittt+l--t-+i-+tH+t----i

Inclusion -

-' (". 10000 1(XX)OO 100)00) Number of cycles ' Figure 3: Low cycle fatigue at 650°C of new PM superalloys compared with alloy NIB. Initiation sites by quasi-cleavage or at ceramic inclusions (SM043 ).

SMO alloys with a medium grain size (45-60 ~)exhibit similar tensile properties with fine grain size N18. However tailored chemistries for matrix strengthening, and medium grain size result in a major improvement in the creep times to rupture and even more in creep times to 0.2% creep strain (Fig.2). This yields a 100°C increase in temperature capability. Fatigue properties are of major importance for disks, but PM superalloys are known to be very sensitive to inclusions [3, 4] which requires good fatigue crack growth resistance [5] for damage tolerance and to increase inspections intervals. Only a limited number of LCF tests have been completed on smooth specimens under strain controlled conditions (zero to maximum strain, at 0.5Hz). At 650°C SM043 exhibits a higher life than N18 (average): crack initiation often occurs in grains along quasi-cleavage facets and less often on ceramic inclusions (Fig.3). These alloys seem to have a reduced sensitivity to inclusions as compared to alloy N18 but this has to be confirmed on a large number of specimens. Fatigue crack propagation remains fairly good, nearly as good as for alloy N18 [2, 6].

The LCF resistance is the major criterion for disk life. Therefore the alloy SM043 was selected as the new superalloy for disks. An extensive study in underway to optimize the microstructures and fatigue properties, both LCF and fatigue crack propagation of this alloy.

2.2. Evaluation of Grain boundary engineering in alloys for disks.

2.2.1. Basics of grain boundary engineering As illustrated in the previous paragraph, the conventional way to improve alloy properties is mostly through the optimization of alloy chemistry, and then that of microstructure parameters such as grain size, or precipitation. This is mostly achieved through the choice of appropriate heat treatments. In recent years grains boundary engineering (GBE) has been proposed as an alternative way to improve alloy properties by optimizing the strength of the grain boundary network taking advantage of the crystallography of grain boundaries [7].

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Grain boundaries are classified using the coincidence site lattice (CSL) model [8]. Each boundary is given an index I: defined as the inverse of the density of lattice sites common to adjacent grains (taken as an integer). Exact orientations are rare but within small deviations boundaries exhibit properties close to these "special orientations". In f.c.c. polycrystals boundaries with I:<29 are often taken as "special boundaries" as opposed to random [7]. Then an analysis of the network in terms of connectivity can be made.

The peculiar importance of twin boundaries has been recognized, since they occur frequently in alloys with low stacking fault energy [9]. Annealing twins are characterized by planar interfaces and a 60° rotation around a <111> and are associated with the coincidence index D. Their planar boundaries have a low diffusion rate and hence migration rate. Further multiple twins can also form and have the same properties of stability. In particular second 2 order twins (noted I:9. or I:3 ) can form from the encounter of two twin boundaries on different octahedral planes within the same original crystal. By the same mechanisms higher order twin boundaries can form I:3° where n is an integer (1,2,3 ... ). Thus a high density of non-coherent grain boundaries of type I:3° can be found in alloys with low stacking fault energy.

Figure 4: Recrystallisation sequence computed using the venex theory (after [ 11] ).

Twin boundary trace Twin plane

Figure 5: Sketch showing the incorporation ofannealing boundary in the recrystallisation scheTM (after [ 12]).

The influence of twinning can be simulated on the evolution of the grain boundary network. The GBE process typically consists in straining followed by subsequent annealing. The

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energy stored in the deformation stage is the driving force for the nucleation and growth of new grain during the re-crystallization phase. The kinetics of grain boundary network evolution depends on the stored energy, temperature and duration of the annealing treatment. The migration of grain boundary assumes a minimisation of the global energy, stored energy in each grain, interfacial energy as a function of crystallographic misorientation between grains. Vertex models are used for re-crystallisation [10] (Fig.4). An adaptation has been proposed recently [11, 12] to include the description of annealing twins (Fig.5).

At high temperature creep as well as low cycle fatigue crack propagation are known to favor inter-granular fracture, especially in superalloys [13]. Therefore if one reduces the number of random grain boundaries and increase the proportion of special boundaries such as l:3n, increased mechanical resistance at high temperature should follow. Some recent experiments under hold time creep fatigue crack growth have shown that the crack was growing preferentially along random boundaries [14].

2.2.2. Application to nickel base superalloys

The application of such a concept to Ni alloys is attractive since they have often a strong tendency to form annealing twins due to their moderate stacking fault energy. This is the aim of the ongoing project so-called ORGANDI and supported the french National Research Agency involving Snecma-Safran, CEA, Turbomeca-Safran, Aubert & Duval, ONERA, CEMEF and Centre des Materiaux (Mines-ParisTech).

~-'· t'. ' ...... ~ • ·---.-·•------/ ,. ·- I 11:

Uttr---~-~---~---~--,

tJ2 Ut £ t

Figure 6: Evolution of the fraction of twin boundaries/to grain boundaries and of the grain size ofU720 Uwith torsion strain (after [15]).

601 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

However superalloys with a high volume fraction of y' precipitates can only be hot-worked whereas most literature investigations used room temperature deformation. The applicability of GBE is investigated using a french grade of U720 Ll. A thermo-mechanical route has been chosen in accordance with industrial needs i.e. to achieve a determined final shape from an initial billet; Compression and torsion were used to finish with a strain of 0.8 [15, 16]. A prestrain of 0.5 was used to cancel billet thermo-mechanical history. The remaining strain was applied in several steps at a temperature below the solvus of secondary precipitates (coherent with matrix, 20-100nm in size); therefore boundary migration is largely inhibited by the distribution of primary precipitates 1-5 J.l.l11 in size, incoherent with matrix and located at grain 1 1 2 1 boundaries. Two strain rates (10" s· and 10" s" ) were investigated and 30 min final heat treatment was carried above the solvus temperature. The proportion of twin boundaries increases rapidly up to a maximum after 0.12 strain and then decreases to a saturation level, as does the grain size when dynamic re-crystallisation occurs (Fig.6). A very large proportion of :E3 boundaries can thus be achieved, much higher than in the standard processing route (as received condition): more special boundaries are created during the super-solvus annealing treatment in the microstructures retaining high dislocations density after straining, i.e. those do not show dynamic re-crystallisation and especially those deformed at high strain rates.

Therefore this work shows that grain boundary engineering is clearly a way to improve the mechanical properties of superal1oys for disks. Creep and fatigue testing is currently under way to compare the properties of GB engineered alloy with standard processing route.

3. Design oriented research for blades: life assessment for increased durability.

3.1./ntroduction.

There is a continued improvement of aircraft engine performance with an increase of turbine entry temperature and pressure ratio. During the last three decades, the increase of blade maximum temperature has been approximately 15°C per year (Fig. 7).

Superalloy

Figure 7: Variation of maximum blade temperature with years and evolution of materials and cooling technologies (left); Turbine blade coated by TBC and scanning electron microscope observation of various layers in the TBC (right).

602 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T . Beck and B. Kuhn.

A limit to alloy development for blades is now achieved with incipient melting temperature around 1350°C and volume fraction of 0.7 of y' Ni) (Ti, AI) precipitates in nickel based alloy single crystals. Recent improvements were achieved combining alloy development and cooling technologies, using hollow components [17]. Thermal - Barrier coating are now introduced to keep average component temperature to a suitable level.

2.2. Low cyck fatigue life in single crystals based on micro-crack growth from pores.

A predominant source of failure in nickel based single crystal superalloys remains, however, the internal pores resulting from the solidification process [18, 19]. A damage model based on the propagation of micro-cracks originating at casting defects has been developed for single crystal (SX) turbine blades, operating under thermo-mechanical creep-fatigue conditions [20- 23]. The model basically used a process zone concept, using a characteristic microstructure element modeled as a cube of side L. The volume L3 corresponds to that of a secondary dendrite of SX superalloy. A probabilistic analysis is carried out in order to determine the number and size of casting pores present in a given volume of material. Another peculiarity of Ni-based SX superalloys is their constitutive and damage anisotropy. The anisotropy of these face-centered cubic materials is therefore considered in this model in the following way: stresses applied to a given volume of material are resolved onto all octahedral ({ 111} <011>) and cubic ({001 }<011>) slip systems.

amoyO 100000 Nfcalc: • <001> ••...,.. - ...... amaxO 8 <001> W • bore 10000 b.<001>7o..,.·bero X <001> 4 11.,..- cooled lKW --led L eh..,.-be,.. 1000 +411> 411opeo-bero •

4 olopeo • bero

Em w 100

T 10

Nfexp 100 'C 600 'C 950 'C 1 I 10 100 1000 10000 100000

Figure 8: Sketch showing the competition between initiation at sub-surface and initiation at internal pores for low cycle fatigue of single crystals. Comparison between model predictions and experiments for various TMF cycles and specimen orientations.

Finite element analyses of blades are made using a crystallographic viscoplastic model [24, 25] where constitutive equations are written at the level of these slip systems. The damage equations are then applied on slip systems themselves using resolved shear and normal

603 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

stresses on every slip system. Material parameters chosen in the model are obtained from uni­ axial material data generated on <001> and <111> specimens. Material properties thus obtained are considered to be intrinsic properties of octahedral and cubic slip systems respectively. Any other combination of loading and crystal orientation can then be reduced to shear and normal stresses applied on active slip systems.

In reality, the pores are of random irregular shapes but for the sake of simplicity the encircling diameter of pores was measured. The distribution of material pores is modeled on the basis of three types of pores (Fig. 8) i.e. mean size pores uniformly distributed over the entire volume, the largest pore within a given volume and the largest pore in the sub-surface zone. This latter is responsible for a single surface crack initiation. Depending upon the applied stress range, the growth of the dominant crack nucleating from the largest surface pore or the largest internal pore can take place in globally elastic or plastic material behaviour. Therefore two different crack propagation regimes are used in parallel. One is denoted here as Basquin's crack propagation regime and the other one as Tomkins' crack propagation regime. Basquin's relation caters for stress controlled crack propagation regime under globally elastic material behaviour. In terms of resolved shear stresses, it can be written as in eqn (1 ):

---·da _L ~~:t.-- r (1) dN n s 2· 'tt

where M is constant, n is the total number of slip systems considered in calculations, ~'t s is the resolved shear stress range on a given slip system and 'tr is critical fatigue shear strength of the material on either cubic or octahedral slip system [21-23]. Tomkins' equation allows calculating crack growth rate under generalized plasticity condition in polycrystalline materials [21]. This equation is therefore rewritten using shear stress and strain components on slip systems:

-=da L B · a and B = [sec( -·tr --•~i' )-1 J ·~"' (2) dN ' ' ' s 2 2·-r/ Is

~Ys is the applied plastic (or viscoplastic) shear strain range on a particular slip system s and "a" is the current length of the dominant crack [21-23].

Creep damage growth is described using the classical continuum damage mechanics approach of Rabotnov. The interaction between oxidation and creep-fatigue damage is introduced assuming that localized oxidation reduces critical strength of material in an area ahead of the dominant crack. Critical strength of material in such an area affected by oxidation is thus lower than in non-oxidized material. This procedure has already been described in some detail for polycrystalline superalloys in cast and wrought forms [see e.g. 21]. One can estimate the size of the area embrittled by simple oxidation performing crack growth tests on non-oxidized compact tension specimens as well as tests on pre-oxidized specimens. The crack length, over which crack growth rates for pre-oxidized or non-oxidized specimens are different, gives direct access to the embrittled zone size. The depth of the embrittled zone It is supposed to be proportional to the oxidized depth Io. (with a constant ratio). The kinetics of oxidation in the

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interdendritic zone of single crystal superalloys, as in polycrystalline material [13] follows a power relation with time and is triggered by plastic strain [22, 23].

Tests lifetime under various cycle shapes, loading conditions and material orientations have been calculated using the same set of equations and model parameters. Different examples of model applications can be found elsewhere for LCF lifetimes under a wide range of test frequencies, temperatures, dwell times, load orientation effects as well as TMF on volume elements tests [2I, 23]. This has been applied to simple structures [20, 2I] and components with good success. Fig. 8 shows the TMF lifetime prediction capabilities of the model, for different TMF test cycles applied to AMI single crystals specimens of various orientations. These cycles have different shapes and lead to different lifetime values under the same applied strain range. The graph shows the comparison of experimental and calculated values. A good correlation is achieved. This work is nonetheless completed currently by an investigation of LCF at sharp notches to investigate the situation in the immediate vicinity of pores, and anomalous regimes of short crack growth can occur under some conditions, depending upon local plasticity and oxidation [26].

2.3. Coated single crystals : Thermal Barrier Coatings.

Blades and vanes in aero-engines are protected against oxidation and corrosion using aluminide coatings. HPT (High Pressure Turbine) blades of SNECMA aircraft engines were first made of AMI single crystal superalloy with CIA coating. This coating is elaborated by, at first, a chromium deposition and then a nickel aluminide (NiAl) deposition on the substrate surface, through Chemical Vapor Deposition (CVD) process. The presence of these coatings is often ignored in component design. Coating resistance is usually assessed using oxidation or cyclic oxidation tests on different kinds of specimens. Under service conditions however, coatings are experiencing cyclic oxidation under stress (due to centrifugal load) as well as thermal transients. Thermo-mechanical creep - fatigue using hold times under load is especially appropriate to assess damage of coated systems under realistic conditions [27]. The industrial need is to know at what time the NiAl tank, which provides the resistance to oxidation in the turbine aggressive environment, will be completely destroyed.

Advanced blades are now coated with a thermal barrier. Fig.7 shows cooled blades used by SNECMA with a TBC that is manufactured by electron beam physical vapour deposition (EBPVD) of yttria partly stabilized zirconia on a platinum modified nickel aluminide bondcoat. EBPVD confers a well-known columnar structure to the partly stabilized zirconia that improves its resistance to transverse cracking. A small layer of alumina is present from the initial condition at the interface between partly stabilized zirconia and bondcoat. The bond coat (BC) inserted between the blade and the ceramic protects the substrate against oxidation by the growth of an alumina layer, the thermally grown oxide (TOO), on its surface. The bond coat also accommodates the substrate strain to prevent ceramic rupture [27-30]. The purpose of Thermal Barrier Coating is to insulate the blade from outside environment and particularly to reduce heat transfer between hot gas flow and the blade alloy. Thermal barrier coatings (TBC) are used in advanced turbines to increase the turbine entry temperature when the temperature capability of the alloy comes to a limit. A strong thermal gradient occurs in the thin outer layer of ceramic, which reduces metal temperature particularly during thermal transients and during steady operation too. This results in a significant increase in component life. The TBC coated blade is actually a multi-material component. In this case coating

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integrity is a key issue in the life of component: coating spallation results in a large temperature increase in alloy substrate, which is detrimental for the overall component integrity.

Many publications investigated the mechanisms leading to spalling of TBC systems. The most important factor is the growth of thermally grown oxide (TOO) at the interface between the ceramic layer (zirconia partly stabilised with yttria) and the bondcoat. Detailed damage mechanisms are dependent on the nature of the bond coat and alloy substrate, NW and AM 1 in the present study. The mechanisms of oxidation as well as the influence of segregating species like sulfur were investigated in our group, as well their influence of TBC adherence to substrate [29-33]. More information on the present system can be found elsewhere [27-30]. Damage of the TBC-substrate interface is a delamination process viewed as a sequence of nucleation, growth and coalescence events of interfacial cracks. Macroscopic buckling near room temperature can result and then lead to catastrophic failure and spalling off the TBC.

Stress, MPa 350 Blade cycle 30'Yo 300 .... cyclic oxidation 250

200 150

100 so

Of--lllo.llli::!--.....,..--f-...... -:2'""-r~---f--t-=-----, -50 -100 Temperature, •c

Figure 9: Thermal-mechanical fatigue cycle used to simulate service loading for a CVD nickel aluminide coating deposited on AMJ single crystals, using a 5 min hold time at maximum temperature; micro-spalling of alumina scale on specimen surface after 8(){) cycles.

Therefore a general methodology has been proposed to develop damage models that applicable to coatings. The need for such an approach is triggered by the search for longer lives in civil applications, where oxidation and aging problems become a major issue. Basically there is a need to develop an extensive data basis using cyclic oxidation tests that are commonly in industrial practice. Relevant thermal-mechanical fatigue tests [27, 34] are added to identify interactions between loading and oxidation effects, which might occur in service. Figure 9 shows a thermal-mechanical fatigue cycle, with a 5 min hold time at maximum temperature, used to simulate service loading at critical areas on single crystal blades that are protected by an aluminide layer. More severe damage is generated during the TMF cycles than during conventional cyclic oxidation tests, that results in surface rumpling and micro-spalling of alumina layers; this favours re-oxidation events and accelerates aluminium depletion of the NiAI coating [27]. Micromechanical models [36, 37] can be of special use as well as detailed observations at fine scales to identify major mechanisms of

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damage accumulation. However it is essential to use interrupted oxidation tests on specimens and to estimate damage evolution using observations or mechanical tests. This methodology has been used for alumina-former coating as well as for TBC.

Cyclic oxidation tests that combine oxidation during high temperature exposure and thermal fatigue during thermal transients are used first. Figure 10 is an example of the effect of maximum temperature and of cycle period on the life to TBC spalling off. Compressive tests were used after different exposure times to static (or cyclic) oxidation [29, 30] to investigate the reduction of the resistance of TBC to spalling off induce by increased exposure times (Fig.ll ). Observations using scanning electron microscope shows that de lamination occurs fust before buckling and catastrophic spalling as usually admitted [39]. On a local scale rumpling of the TOO is observed which is triggered by thermal cycling as reported for aluminide, and thermal-mechanical fatigue [27]. While small pores nucleate preferentially at the TOO-boncoat interface, this rumpling favours the nucleation of cracks at the TOO­ zirconia interface [40].

116 1200 Temperature, oc 115 1000 114 800 113 800 ...=E 112 111 thokl2

A life prediction model for TBC systems has been developed, taking into account a local scale where damage takes place at interfaces and a macroscopic scale where buckling leads to the TBC spallation. A post-processor is built assuming that in plane strain of the substrate is imposed to the multi-layer system composed of bondcoat, TGO and zirconia topcoat. Zirconia is described using as transversely isotropic elasticity, TOO was assumed elastic and bondcoat obeys a Norton type viscoplastic behaviour, with hardening. The total thickness of the TOO layer was described though a parabolic equation, identified from thickness measurements. The damage model is then identified assuming two major contributions: one due to oxide growth Dox. and the other due to the development of rumpling and thermal cycling Drvmp/btg· Rumpling is observed for cyclic loading involving viscosity of the bondcoat at high temperature. Therefore an engineering description of this effect was made using viscous strain of bondcoat e. cum and oxide thickness to simulate the damage induced by thermal transient:

dDox = f(hox) (4)

= ) dDI'IIlfi/Jlito8 g(h.,,t;"' ,N111 (5)

607 gth Liege Conference : Materiala for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

where N1h represents the number of thermal cycles experienced. The final spallation of the system was identified from buckling of the delaminated area, considered as a circular blister. The critical strain with clamped boundary conditions, £.:,;1, is thus predicted by the classical buckling theory for elastic analysis [39]:

Ecrit =1.22{ ~cJ {1-D) (6), D is total damage, Ro is determined from a spalling test applied in the initial condition (when D is equal to 0). This critical condition is applied to the strain tensor experienced by the TBC. The model is tested for different thermal cycles used in the experimental database. Compression tests conducted to spalling are used to identify damage parameters of the model after different thermal-mechanical conditions giving rise to different oxidation conditions and oxide thicknesses. Figure 11 (left) shows for instance the evolution of macroscopic compression strain to spalling with isothermal oxidation times. A good fit is actually achieved between model and experiment. Predictions are finally shown for cyclic oxidation tests (Fig.11 right) where the number of oxidation cycles predicted is in good agreement with experimental data. The developed model is able to represent oxidation effect for long time exposure at high temperature as well as thermal fatigue life.

Oxidised/virgin Critical strain ratio I o.e ._..,__ ~ o.a I ~ I 0.7 ~ 0--- 1-1 o.a n +- ! I ... jl t h •.. I 0.3 I j •..0.2 -- -· - I I I f-t 0 +~t- ~ ~ 0 200 400 ... 1000 ... NexpiNexpO 10 Time at 11oo•c, h

Figure 11: Thermal cycle and influence of maximum temperature (and hold time) on the lifetime to TBC spalling off on AM1 superalloy during cyclic oxidation.

4. Conclusions

For disk application, several ways seem promising. A new PM superalloy has been developed with improved creep and fatigue resistance. Grain boundary engineering seems to be applicable to wrought alloys but still needs proper assessment. For blades made of single crystal superalloys, a lifetime model has been identified using micro-crack growth from casting pores. A methodology proposed for the degradation of coating seems promising for thermal barrier coating spalling.

Acknowledgements The financial support of Snecma, French Defence Agency, European Community and French National Research Agency is gratefully acknowledged. The authors are indebted to many collaborators who have contributed to various aspects of this work within Snecma, Mines ParisTech and at ONERA.

608 gtn Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

References

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609 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

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610 gth Liege Conference : Materials for Advanced Power Engineering 2010 edited by J. Lecomte-Beckers, Q. Contrepois, T. Beck and B. Kuhn.

[34] A.Peichl, T. Beck, 0. Vohringer, Behaviour of an EB-PVD thermal barrier coating system under thermal-mechanical fatigue loading, Surface and Coating Technology, 162, 113-118, (2003). [35] T. Beck, R. Herzog, 0. Trunova, M. Offerman, R.W. Steinbrech, L. Singheiser, Damage mechanisms and lifetime behavior of plasma-sprayed thermal barrier coating systems for gas turbines- Partll: Modeling, Surface and Coatings Technology, 202, 5901-5908, (2008). [36] M. Caliez, F. Feye1, S. Kruch, J-L. Chabocbe, Oxidation induced stress fields in an EBPVD thermal barrier coating, Surface and Coatings Technology, 157, 103-110 (2002). [37] D.S. Balint, J.W. Hutcbinson, An analytical model of rumpling in thermal barrier system, Journal ofthe Mechanics and Physics of Solids, 53,949-973, (2005). [38] V.K. Tolpygo, D.R. Clark:e, Temperature and cycle-time dependence of rumpling in platinum-modified diffusion aluminide coating, Scripta Materialia, 57, 563-566, (2007). [39] I.W. Hutcbinson, M.D. Thouless, E.G. Liniger, Growth and configurational stability of circular buckling-driven film delaminations, Acta Materialia, 40, 295- 308, (1992.) [40] C. Courcier, V. Maurel, L. Remy, A. Phelippeau, Damage based life prediction model for EBPVD thermal barrier coatings , LCF6, Sixth International Conference on Low Cycle Fatigue, September 8-12, 2008, Berlin, DVM, 155-160, (2008).

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