High temperature creep degradation of the AM1/NiAlPt/EBPVD YSZ system Fanny Riallanta,b, Jonathan Cormiera, Arnaud Longuetb, Xavier Milheta and José Mendeza a
Institut Pprime, CNRS – ENSMA - Université de Poitiers, UPR CNRS 3346, Département Physique et Mécanique des Matériaux, ISAE-ENSMA - Téléport 2, 1 avenue Clément Ader, BP 40109, 86961 Futuroscope Chasseneuil Cedex, France b Snecma – SAFRAN group, Site de Villaroche, Rond-Point René Ravaud, 77550 Moissy-Cramayel, France
Corresponding author:
[email protected]
Abstract The failure mechanisms of a NiAlPt/Electron Beam Physical Vapor Deposition (EB-PVD) Yttria-Stabilized-Zirconia (YSZ) Thermal Barrier Coating (TBC) system deposited on the AM1 single cristalline substrate have been investigated under pure creep conditions in the 1273 K1373 K (1000°C-1100°C) temperature range and for durations up to 1000 hours. Doubly tapered specimens were used allowing the analysis of different stress states and different accumulated viscoplastic strains for a given creep condition. Under such experiments, two kinds of damage mechanisms were observed. Under low applied stress conditions (i.e. long creep tests), microcracking is localized in the vicinity of the Thermally Grown Oxide (TGO). Under high applied stress conditions, an unconventional failure mechanism at the substrate/bond coat interface is observed due to large creep strains and fast creep deformation, hence leading to a limited TGO growth. This unconventional failure mechanism is observed despite the interfacial bond coat/top coat TGO thickening is accelerated by the mechanical applied stress beyond a given stress threshold.
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I. Introduction
The enhancement of gas turbines performances is largely dependant on the turbine entry temperature [1]. As a result, various solutions have been gradually introduced in the engine design over the last four decades such as the development of advanced turbine materials (superalloys), new cooling concepts (e.g. film cooling), novel combustor designs and thermal barrier coatings (TBC). A TBC system is composed of different layers deposited on a load-bearing superalloy substrate including a bond coat (BC) which acts as a protection against high temperature oxidation and corrosion, a top coat (TC) thermal insulator and the thermally grown oxide (TGO) at the BC/TC interface.[1, 2] They are currently used for civil and military applications in order to protect structural parts such as high pressure turbine blades and vanes or combustor liners from oxidation and to decrease the substrate temperature. Because of the very different damage processes encountered during service operation of the TBC systems (e.g. oxidation of the BC,[3] BC rumpling,[4, 5] TC sintering,[6] TC erosion,[7] CMAS attacks,[8] …), original engine makers are often conservative in the design of coated blades, leading to numerous and expensive close inspections and maintenance of engines. Many authors have developed physical and numerical predictive models for TBC failure mainly based on cyclic oxidation and/or compression tests which are thought to be the most representative of the in-service damage processes.[9-13] The latter are based on the assumption that TBC systems failure occurs during cooling from a high temperature operating stage, leading to compressive stresses in the TBC system.[14] However, in order to take the whole advantage of TBC systems, it is necessary to get a better understanding of the damage and failure mechanisms and of their possible interplays. Indeed, even if it is widely admitted that a first estimation of the
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TBCs life under “close-to” operation conditions can be obtained using cyclic oxidation testing, it supposes that the TBC spallation results from thermally driven processes, without any contribution of the macroscopic stresses arising from the centrifugal forces. In fact, turbine blades are subjected to complex thermomechanical loadings, including dwell times during some ratings of a complex civil or military flight. A more realistic approach to the problem would require considering the contribution of viscoplasticity to the damage processes of the TBC system and to the oxidation kinetics of the BC. While considering the contribution of creep to the damage mechanisms of TBC’s systems may appear as a non sense since these coating were developed to limit creep damage by decreasing the metal temperature, some engine ratings may induce a non-negligible contribution of viscoplasticity in some component sections close to stress concentrators or insufficiently cooled areas (e.g. close to platforms).
To address this issue, tensile creep tests were performed on a standard Snecma TBC system (AM1/NiAlPt/YSZ) in the 1273 K-1373 K (1000°C-1100°C) temperature range, and for durations up to 1000 hours (either interrupted or up to failure). A special attention has been paid to the nucleation of microcracking at each interface in the TBC system, especially at the BC/TC interface close to the TGO and at the substrate/BC interface, as a function of time. To get a better understanding of the observed failure mechanisms, the TGO thickening has also been characterized as a function of the local stress magnitude.
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II. Experimental procedures A. Materials The TBC system used in this study is composed of an YSZ top coat (i.e. 93 wt pct ZrO2 and 7 wt pct Y2O3) and a NiPtAl BC (composition is provided in Table I). This TBC system was deposited on the AM1 single crystal superalloy substrate. The AM1 is a first generation nickel based superalloy whose composition is also given in Table I and whose solution heat treatment prior to the TBC deposition is 3h at 1573 K (1300°C), gas fan cooling. The BC was deposited by APVS (Aluminium Phase Vapor Snecma, the Snecma process for aluminisation) at Snecma (Châtellerault plant, France) after an initial sand blasting of the single crystalline samples. The TC was then deposited by an Electron Beam Physical Vapor Deposition (EB-PVD) process at the Ceramic Coating Center (Châtellerault, France). Their thicknesses were close to 70 µm and 160 µm respectively. Samples were submitted to a final heat treatment of 16 hours at 1143 K (870°C) followed by air quench to optimized the γ/γ’ microstructure of the substrate. During the TC deposition process and subsequent straining, a thermally grown oxide develops at the BC/TC interface. In this system, it is a pure α-alumina whose initial thickness after the TC deposition is approximately 0.5 µm. The typical as-received microstructure is shown in Figure 1. The specimens used in this study were doubly-tapered with diameters ranging from 2.5 mm (in the samples minimum section) to 4 mm, and with a 14 mm effective length (see Figure 2). In addition, a smooth cylindrical sample (3 mm in gage diameter and 14 mm in gage length) was used to assess the relevance of the TBC coating system damage processes observed using doubly tapered samples. Prior to the coating deposition, samples were carefully hand polished to remove residual stresses introduced during machining and to reduce roughness (Ra < 1 µm). The coating
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was deposited on the gage length of the samples, as well as on the fillets sections with the heads. The geometry of the sample allows the analysis of different stress states and different accumulated viscoplastic strains for a given creep condition. Stress free (i.e. without any applied mechanical stress) oxidation kinetics were established using doubly-tapered samples which were thermally annealed for different durations at different temperatures.
B. Creep Tests
Creep tests were performed using a specific set-up developed in the laboratory. They were performed under constant load. Heating was ensured by a resistive furnace and displacements were measured using a contactless Linear Variable Differential Transformer extensometer. Two specimens can be tested simultaneously at temperatures up to 1873 K (1600°C). The temperature heterogeneity along the gage length was less than 2°C. In the following of the article and for sake of simplification, each sample will be characterized by the maximal substrate nominal stress σmax achieved in the smallest bearing section of the doubly-tapered sample. Creep tests were performed using a σmax in the 120 - 200 MPa stress range, a majority of the experiments being carried out under a maximal nominal stress of 120 MPa at 1273 K (1000°C), 1323 K (1050°C) and 1373 K (1100°C) (Table II). The tests durations ranged between few hours up to a thousand of hours. Non-failed specimens were cooled under load to prevent any stress free damage development in the TBC system during cooling down to room temperature and any TC spallation. Failed specimens were rapidly cooled down to room temperature after failure in order to freeze the microstructure (at 30 K/min in the 1323 K/1373K – 1023 K (1050/1100°C – 800°C) temperature range, and then slower through conduction in the grips and natural convection).
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Moreover, after creep testing, samples were protected from vapor water using a hydrophobic spray in order to avoid any “desktop spallation” before metallographic preparation.[15]
C. Microscopic observations
Prior to observations, the samples were mounted in an electrical conductive resin. They were subsequently cut along the tensile axis and mechanically polished without any water as lubricant. The different layers of the system were observed after creep or oxidation testing using the backscattered imaging mode of a JEOL JSMTM 7000F scanning electron microscope (SEM). For each micrograph, the distance between the center of the sample and the center of the picture was carefully measured in order to determine the local substrate nominal stress for each TGO thickness. Image analyzes were performed using the Visilog® software. Specific algorithms developed in the laboratory were used to characterize the oxide thickness evolution as a function of time, of temperature and of the theoretical substrate nominal stress. The TGO thickness was measured perpendicularly to the local medium plane of the oxide, i.e. taking into account the local curvature of the oxide. At least 10 sections were analyzed for each sample, and 5 measurements of the TGO thickness were performed per image.
III. Results A. TGO thickening
Figure 3 shows the evolution of the TGO thickness as a function of the nominal applied stress for a creep test performed at 1273 K (1000°C) under a maximum nominal stress of 120 MPa, and interrupted after 600 hours. Each single measurement is the mean value of the TGO thickness per
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section while error bars correspond to the standard deviation. A relatively large scatter is observed for each section. This is clearly visible in Figure 4 where a progressive roughening of the oxide layer is observed with the increasing creep time. This roughening mainly results from a wrinkling process. From Figure 3, it is not observed any impact of the applied stress on the TGO thickening at 1273 K (1000°C). The evolutions of the TGO thickness at 1323 K (1050°C) and at 1373 K (1100°C) for different local nominal stresses are plotted in Figures 5(a) and 5(b) respectively. For sake of clarity, the error bars have been removed. The largest error bar among all single measurements is given for reference. In Figure 5(a), a very limited impact of the local mechanical stress on the oxidation kinetics at 1323 K (1050°C) can be observed, especially for short oxidation times (< 200h). The influence of the applied stress can only be noticed for the longest oxidation duration and for local stress above 100 MPa. Under such conditions, the TGO thickening is more pronounced. This effect is even more obvious when considering the stress free oxidation kinetics superimposed in Figure 5(a). In contrast, oxidation kinetics at 1373 K (1100°C) show a more pronounced influence of the nominal applied stress on the TGO thickening, whatever the mechanical loading (Figure 5(b)). This is particularly obvious for the transient oxidation regime of fast scale growth. By superimposing the stress free oxidation kinetics in Figure 5(b), an enhancement of the oxide scale growth during the first hundreds of hours is clearly observed. This is particularly noticed for applied stresses above 60 MPa, i.e. for conditions where the plastic flow of the AM1 substrate is non-negligible.[16] Theses plots are in agreement with previous results established by Seo and al. using another TBC system deposited on a polycrystalline substrate for which an accelerated TGO thickening under tensile creep load was observed.[17]
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The impact of the mechanical applied stress on the TGO growth has been investigated in details using SEM observations focused on the morphology of the interfacial oxide. Figure 6(a) is a high resolution observation of the TGO after 50 hours in creep at 1373 K (1100°C)/σmax = 140 MPa (interrupted creep test). Inhomogeneities in the dense and continuous Al2O3 oxide layer (highlighted in Figure 6 by arrows) can be observed. These differences in TGO contrast were systematically observed for interrupted creep tested samples after short durations, typically below 100h, whatever the temperature and the stress applied. These TGO imperfections are very similar to Y2O3/YAG particles as identified by Evans et al.[18]
B. Damage processes in the TBC system
The failure mechanisms of the TBC system involved in both long and short creep tests were also investigated using SEM observations. They were focused both in the vicinity of the TGO and at the BC/substrate interface in order to detect the preliminary failure stages (which will also be termed in the following of the article as the micropropagation stages). Figure 7(a) shows the first stages of the TBC system failure close to the lowest bearing section for a long creep test. Decohesions located at the TGO/BC interface are observed after 1000 hours at 1273 K (1000°C)/σmax=160 MPa while the TGO is still totally adhesive to the TC. After 1168 hours at 1323 K (1050°C)/σmax=130 MPa, the first stages of microcracking are slightly different. Evidences of delamination are observed at both interfaces of the oxide layer and two types of alumina are observed (see arrows in Figure 7(b)). These two types of alumina result from a reoxidation of the BC surface after cracking of the TGO upper layer. The upper layer is presumably α-Al2O3 while the lower one should be θ-Al2O3 due to a shorter oxidation time.[19] For these two tests, the substrate/BC interface seems hardly affected by the mechanical loading
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since no microcracks could be observed in the vicinity of this interface. The only noticeable microstructure evolution close to this interface, is the precipitation of the µ phase due to the interdiffusion between the substrate and the BC, in good agreement with the literature.[20, 21] The damage mechanisms located close to the TGO are typical of long interrupted creep tests performed under low applied stress where a sufficient oxidation time led to the development of a thick interfacial oxide layer. For creep tests led up to failure, the damage mechanisms of the TBC system are clearly different. It has been evidenced that the delamination is located at the BC/substrate interface, more precisely, at the BC/interdiffusion (IDZ) interface, as observed for creep tests performed at 1373 K (1100°C) up to failure (see Figure 8). For sake of simplification, this interface will still be denominated as the “BC/substrate interface” in the following of the article, although there is the presence of an evolving IDZ with creep time below the initial substrate surface. Actually, after failure, the damage development is quite limited in the vicinity of the TGO. Segmentation of the TGO and of the TC is observed resulting from the bending of the upper part of the failed coating (TC + TGO + BC). This bending is likely to derivate from the relaxation of the stored strain (elastic and plastic) energies. Figure 9 shows a similar failure mechanism after a creep test performed using a smooth sample at 1373 K (1100°C)/σmax=160 MPa up to failure. The damage at the substrate/BC interface was also observed for this smooth sample, hence demonstrating that this failure mechanism is not an effect of the specimen geometry but really related to a typical failure mode in these conditions. A creep test at 1323 K (1050°C)/σmax=160 MPa interrupted just prior to failure also demonstrates a similar damage processes located at the substrate/BC interface (Figure 10). It is observed that the delamination starts in the lowest bearing section and propagates along the substrate/BC
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interface. Moreover, it is observed in Figure 11 that, under conditions involving high creep strain rates and resulting in short creep lives, the very first stages of microcracking at the substrate/BC interface start at pores or oxidized corundum particles located at this interface. Such microstructural defects, initially located at the BC/substrate interface after the coating deposition, and likely to evolve during creep deformation for the pores, are therefore critical in controlling the failure mechanism under these conditions.
IV. Discussion
A. TGO thickening behavior
It has been observed in this study that the TGO thickening is affected by the mechanical loading if a critical applied stress and/or temperature is reached (Figure 5). This result is quite in good agreement with previous results by Seo et al. using the APS YSZ/CoNiCrAlY/IN738LC system.[17] It must also be pointed out that other authors did not notice any impact of the mechanical loading of the TGO thickening under pure isothermal conditions using the René N5/MCrAlY/APS YSZ system.[22] Hence, the key question to answer is what could be the origin(s) of such a mechanical dependence? Assuming that the substrate imposes its plastic strains to the TBC system, especially to the BC and to the TGO,[11,
14]
it is interesting to correlate the TGO thickening to the substrate creep
strain rate. Such a plot is performed in Figure 12 for a creep test up to failure at 1373 K (1100°C)/σmax=120 MPa. A good correlation between the macroscopic plastic displacement rate
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(mainly resulting from the substrate inelastic deformation) and the TGO thickening is observed. During the first fifty hours, the macroscopic creep strain rate decreases. Accordingly, the oxide scale grows very rapidly and the growth progressively slows down along with the strain rate. During the secondary creep stage where the strain rate is almost constant, the oxidation rate is also relatively constant. In addition, it was observed a stress threshold below which no impact of the applied stress was on the TGO growth (Fig. 5). Such a stress threshold is understood as the minimum stress inducing a local creep elongation of the substrate, and then, an impact on the alumina scale thickening. To explain such a plastic rate dependence of the oxidation kinetics, it is usually hypothesized that a modification of the oxidation kinetics of metallic materials under mechanical loading can originate from the (micro-)cracking of the oxide scale.[23-25] In fact, the first step of the fast oxide growth coincides with the large macroscopic creep strain at the beginning of the test. During this stage, the oxide scale micro-cracking should take place to accommodate the fast substrate deformation and the oxidation kinetic increases.[26] Thereafter, during the secondary creep stage, the slower strain rate helps reducing the nucleation rate of micro-cracks in the TGO promoting a decrease of the overall oxide growth kinetics. Nevertheless, we did not clearly detect microcracks after unloading samples exhibiting no damage (i.e. no cracks along one of the TC/TGO or TGO/BC interface) (see e.g. Figure 4), even close to defects in the TGO scale, as those observed in Figure 6. Indeed, contrary to what is observed under thermal cycling conditions, for which the TGO behavior is alternatively plastic during high temperature dwell times and brittle during the cooling stages, the continuous high temperature exposures used in the present experiments allow a continuous viscoplastic behavior of the TGO.[27,28] Hence, fast re-oxidation of potential micro-
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cracks nucleating in the TGO, due to the viscoplastic flow of both the substrate and the BC, is likely to occur. The effect of the mechanical loading on the TGO thickness could also originate from two other reasons. A modification of the diffusion kinetics of the chemical species under applied stress can first be invoked. Indeed, an applied stress is known to enhance the diffusion kinetics,[29- 32] and in the present case, a higher dislocations density in the BC could even enhance the migration of chemical species involved in the formation of the TGO (Al especially) through pipe diffusion. Secondly, the creep loading can also affect the theta to alpha transformation of the Al2O3 scale,[19, 33, 34] a transformation that is known to affect the oxidation kinetics.[19] Such a possible effect would require deeper investigations to be elucidated. However, from the present experiments, all these three assumptions are plausible and none of them seems to be the most likely one to explain the impact of the applied stress on the TGO growth. Experiments under compression could also provide additional insights to understand the observed effect of the tensile mechanical loading on the TGO growth.
B. Failure mechanisms of the TBC system
In the present study, it is observed that despite a faster TGO thickening, decohesions develop along the BC/substrate interface under fast creep rate conditions (Figures 8, 9 and 10). Under slow creep rate conditions, the damage develops at one of the TGO interfaces (Figure 7(b)). Hence, upon progressive oxidation of the BC, a progressive transfer of the critical locations (i.e. locations where cracks will preferentially develop) from the substrate/BC interface to the BC/TGO or TGO/TC interface is observed.
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This is the first time, to the authors’ very best knowledge, that such a failure mechanism at the substrate/BC interface is observed within a TBC coating system. One of the reasons for such a delamination process at the BC/substrate interface is the reduction in the load-bearing section due to the creep deformation of the substrate. This reduction in the bearing section is probably accompanied by an increase in the radial and tangential stresses magnitudes due to a locking effect of the columns of the TC, stresses which cannot be infinitely relaxed by the viscoplastic flow of the BC. In addition, under such conditions of high creep rate, microstructural defects at the substrate/BC interface are critical in controlling the nucleation of cracks because they act as stress concentrators. Indeed, it is observed that the damage always nucleates from interfacial pores and residual sand-blasting corundum particles. Furthermore, plastic strain rate is so fast that the stress relaxation due to creep processes is probably very limited (or even impossible) to avoid local decohesion, leading to this unconventional failure mechanism (see Figures 8 and 9). It is also worth mentioning that during creep deformation under such high temperature conditions, the IDZ beneath the initial substrate surface is evolving due to cross diffusion processes between the BC and the substrate. This interdiffusion can lead to the condensation of pores at the BC/IDZ interfaces, but also, to an evolution of the local mechanical properties both within the BC and the IDZ, as already observed previously.[21, 34-36] This point will be discussed subsequently in this section. For longer creep tests, the TGO is likely to thicken enough so that it becomes the weakest point of the TBC system. Under these conditions of low applied stress, the failure mechanism is classical for a TBC system: spallation occurs due to both the local rugosity and the local (growth) stresses.[3,
18]
Since the present experiments were performed under isothermal
conditions, a rumpling process cannot be responsible for such an increase in the local rugosity.[4,
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37]
We may however mention that grain boundaries sliding (due the BC creep flow) is likely to be
the origin of this rugosity development, as observed in Figures 4(a) and 6. Using all our results, a schematic illustration of the main failure mechanism domains encountered in our TBC system has been established in Figure 13. Based on finite element simulations not shown in this article,[16] substrate/BC decohesion has been estimated to occur for creep strains above a 0.8%-1.0% local creep strain threshold. Upon decreasing the applied stress and/or temperature, the critical creep strain at which substrate/BC delamination occurs increases due to the thickening of the TGO resulting in the displacement of the weak point of the system in the TGO vicinity. This transition in the main decohesion initiation site (dashed area in Figure 13) is quite difficult to delineate in a creep strain – time plot since both the local stresses at each interface and the oxidation kinetics are affected by the stresses applied macroscopically, the evolution of the IDZ and the relaxation processes occurring in the course of creep deformation. As an illustration, the two types of damage mechanisms were observed for the sample interrupted after 312 hours at 1323 K (1050°C)/160 MPa (Figure 10).
A more convenient way for interpreting the change in the failure localization can be performed using an energetic approach. Let’s consider Gc-BC/SX and Gc-TGO the interfacial toughnesses at the substrate/BC and at one of the TGO interfaces respectively, and GBC/SX and GTGO the driving forces for the TBC system spallation at the substrate/BC and at one of the TGO interfaces respectively. Such spallation driving forces exist to relax the stored elastic-plastic energies in the coating system. Considering our tensile situation, these driving forces for spallation have the following generic expression:[4, 38, 39]
G i = α i × σ 2 × h i [1]
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where σ is the local mechanical stress, hi the thickness of the layer i and αi, a material constant inversely dependent on the Young’s modulus of the considered layer. Neglecting the differences in local stresses, it can be assumed reasonably, particularly at the beginning of the creep tests, that GBC/SX > GTGO based on: -
the relative thicknesses of the BC and the TGO (hBC > hTGO),
-
and their respective Young’s moduli.
The situation of substrate/BC spallation in case of fast creep deformation conditions is schematically illustrated in Figure 14(a) where the GBC/SX curve intersects the Gc-BC/SX evolution curve earlier than the GTGO/Gc-TGO one (see the red cross in Figure 14(a) highlighting failure). Indeed, under such conditions, the interfacial toughnesses in the TGO vicinity are assumed to be higher than at the substrate/BC interface due to the presence of processing defect in the form of pores and residual sand-blasting corundum particles. Conversely, for longer tests, even if the driving forces for spallation evolve more slowly, the decrease of the interfacial toughness close to the TGO is much more pronounced than the decrease of Gc-BC/SX, leading to a change in the localization of the fatal decohesion (see the blue cross in Figure 14(b)). Still, the kinetics of the degradation of the interfacial toughnesses Gc-BC/SX and Gc-TGO with time are open to debate and would require a full characterization on the way they evolve under the action of environment (i.e. thickening of the TGO), the sintering of the TC, the ratcheting of the BC for Gc-TGO,[38, 39] and the interdiffusion processes, phase changes, precipitation of TCP phases for Gc-BC/SX.[12, 35, 40, 41] The evolution of Gc-TGO has been placed according to available references from the open literature.[18, 39, 42]
No data is available for the evolution of the substrate/BC interfacial toughness as a function
of time, and the relative position of both curves in Figure 14 have been chosen to scale with experiments. The microstructure evolutions close the IDZ/BC interface (in the IDZ and in the
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BC) would require a deeper investigation to quantify how it modifies the local mechanical properties and, consequently, the evolution of Gc-BC/SX. Actually, Figure 14 is just a schematic illustration summarizing and “averaging” all the underlying physical mechanisms responsible for a transition in the localization of the TBC system failure under creep conditions from the substrate/BC interface to the TGO.
V. Conclusions
The failure mechanisms of the AM1/NiAlPt BC/EB-PVD YSZ TBC system have been investigated during high temperature creep tests performed using doubly-tapered samples in the 1273 K-1373 K (1000 – 1100°C) temperature range. It has been observed that for long creep tests (i.e. low applied stress), the failure mechanism of the TBC system is mainly driven by the decohesion of the TGO layer with the BC and/or the TC. Despite an acceleration of the TGO thickening with an increased applied stress, it has been observed that under fast creep rate conditions, interfacial delamination is mainly localized at the substrate/BC interface resulting from the nucleation and growth of voids or microcracking from sand-blasting corundum particles. This unconventional failure mechanism for the TBC system derives mainly from strain incompatibilities at the substrate/BC interface and negligible oxidation effects at the TC/BC interface.
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ACKNOWLEDGEMENTS
The authors are grateful to Snecma – SAFRAN group for FR PhD grant, for sponsoring this work and for providing the samples. Florence Hamon and Florent Mauget (Engineers in the Physics and Mechanics of Materials Department at the Pprime Institute) are gratefully acknowledged for their help and for fruitful discussions in conducting several experiments. This work is conducted under the French program “PRC Structures Chaudes’’ involving Snecma-SAFRAN group, Turbomeca-SAFRAN group, ONERA, CEAT-DGA and CNRS Laboratories (Mines Paris Tech, Institut Pprime–ISAE-ENSMA, LMT-Cachan and CIRIMATENSIACET). Financial support by the French Ministry of Transportation (Direction des Programmes de l’Aviation Civile) is also gratefully acknowledged.
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40. J. A. Haynes, B. A. Pint, Y. Zhang and I. G. Wright, Surf. Coat. Tech., 2008, vol. 203, pp. 413-6. 41. M. Gell, J. Eric, V. Krishnakumar, K. Mccaron, B. Barber, Y.-H. Sohn and V. K. Tolpygo, Surf. Coat. Tech., 1999, vol. 120-121, pp. 53-60. 42. JR. Vaunois, Thermal Barrier Coating III Conference, M.J. Maloney, U. Schulz, D. Rickerby, R. Darolia, O. Lavigne, H. Murakami and H. Guo, Eds., ECI, Irsee, Germany, 2011.
20
Tables Captions
TABLE I. AM1 and BC compositions (wt. pct).[11] TABLE II. Creep test conditions
21
AM1 BC
Min Max
Ni bal. bal. bal.
TABLE I. AM1 and BC compositions (wt. pct).[11] Co Cr Mo W Ta Al Ti 6 7 1.8 5 7.5 5.1 1 7 8 2.2 6 8,5 5.5 1.4 4.2 5.5 0.2 0.1 1 31 0
C
Fe
0.01
0.2
Pt
3.9
22
Temperature (K) 1273 (1000°C)
1323 (1050°C)
1373 (1100°C)
TABLE II. Creep test conditions Maximal nominal stress Test duration (hours) σmax (MPa) 0 600 120 600 160 1000 50 100 120 200 300 1168 160 312 24 50 100 120 210 307 436 140 50 160 42 160 (smooth sample) 28 2 200 4
Failure ? No No No No No No No No No No No No No No Yes No Yes Yes No Yes
23
Figures Captions Fig. 1 - Typical as received TBC system (a) and magnification on the BC/TC interfacial TGO (b). Fig. 2 - Sample’s geometry (all dimensions in millimeters). Note that only the doubly-tapered gage length and the fillets with the heads of the samples were coated. Fig. 3 - Evolution of the TGO thickness as a function of the applied stress after 600h at 1273 K (1000°C). Error bars corresponds to +/- the standard deviation. Fig. 4 - Evolution of the TGO thickness and morphology as a function of time at 1373 K (1100°C) for two different initial nominal stresses. Fig. 5 - Oxide thickness evolution as a function of the creep test duration for different nominal stresses at 1323 K (1050°C) (a) and 1373 K (1100°C) (b). Note that the arrows represent the maximum difference between each single measurement at a given stress level and considering the entire database. Fig. 6 - TGO inhomogeneities after 50 hours at 1373 K (1100°C)/σmax = 140MPa. Fig. 7 - BC/TGO delamination after 1000 hours at 1273K(1000°C)/σmax = 160 MPa (a) and after 1168 hours at 1323 K (1050°C)/σmax = 130 MPa (b). White arrows highlight the TGO/BC delamination along with some micro-cracks propagating within the BC. Microcracking is observed at both the TC/TGO and TGO/BC interfaces in (b). Fig. 8 - Failures of the TBC system after 4 hours in creep at 1373 K (1100°C)/σmax = 200 MPa (a) , 42 hours at 1100°C/σmax = 160 MPa (b) and 436 hours at 1373 K (1100°C)/σmax = 120 MPa (c). Fig. 9 - Failure of the TBC system on a smooth sample creep tested up to failure at 1373 K (1100°C)/σmax = 160 MPa. Fig. 10 - Failure of the TBC system after 312 hours of the 1323 K (1050°C)/σmax = 160 MPa. Fig. 11 - Failure of the TBC system after 42 hours at 1373 K (1100°C)/σmax = 160 MPa (a) and microcracking initiation at (b) pores and at (c) sand-blasting corundum particles. Fig. 12 - Oxidation kinetics at 1373 K (1100°C)/σmax = 120 MPa as a function of the macroscopic displacement rate evolution. Fig. 13 - Schematic illustration of the different damage mechanisms domains during viscoplastic solicitations of a TBC system.
24
Fig. 14 - Schematic illustrations of the evolutions of the interfacial fracture toughnesses at the BC/SX (Gc-BC/SX) and at one of the TGO interfaces (Gc-TGO), and of the driving forces for spallation at the BC/SX or TGO interfaces for a high (a) and low (b) plastic strain rate creep test. Crosses highlight failures.
25
(a)
(b)
TC
TGO BC
Substrate
100 µm
10 µm
Fig. 1 - Typical as received TBC system (a) and magnification on the BC/TC interfacial TGO (b).
26
Fig. 2 - Sample’s geometry (all dimensions in millimeters). Note that only the doubly-tapered gage length and the fillets with the heads of the samples were coated.
27
4 Thickness (microns)
3,5 3 2,5 2 1,5 1 0,5 0
20
40
60
80
100
120
140
Stress (MPa)
Fig. 3 - Evolution of the TGO thickness as a function of the applied stress after 600h at 1273 K (1000°C). Error bars corresponds to +/- the standard deviation.
28
24h
100h
307h
σmax = 120 MPa 5 µm
σmax = 60 MPa
Fig. 4 - Evolution of the TGO thickness and morphology as a function of time at 1373 K (1100°C) for two different initial nominal stresses.
29
6
120 MPa 80 MPa 40 MPa
Oxide thickness (µm)
5
(a)
100 MPa 60 MPa stress free oxidation
4 3 2 1
max scatter
0 0
100
200
300
400
500
Test duration (h)
6
200 MPa 60 MPa stress free oxidation
Oxide thickness (µm)
5
120 MPa 30 MPa
(b)
4 3 2 1
max scatter
0 0
100
200 300 Test duration (h)
400
500
Fig. 5 - Oxide thickness evolution as a function of the creep test duration for different nominal stresses at 1323 K (1050°C) (a) and 1373 K (1100°C) (b). Note that the arrows represent the maximum difference between each single measurement at a given stress level and considering the entire database.
30
2 µm
10 µm
Fig. 6 - TGO inhomogeneities after 50 hours at 1373 K (1100°C)/σmax = 140MPa.
31
(a)
10 µm
(b)
10 µm
Fig. 7 - BC/TGO delamination after 1000 hours at 1273K(1000°C)/σmax = 160 MPa (a) and after 1168 hours at 1323 K (1050°C)/σmax = 130 MPa (b). White arrows highlight the TGO/BC delamination along with some micro-cracks propagating within the BC. Microcracking is observed at both the TC/TGO and TGO/BC interfaces in (b).
32
(a)
100 µm
100 µm
(b)
(c)
100 µm
100 µm
Fig. 8 - Failures of the TBC system after 4 hours in creep at 1373 K (1100°C)/σmax = 200 MPa (a) , 42 hours at 1100°C/σmax = 160 MPa (b) and 436 hours at 1373 K (1100°C)/σmax = 120 MPa (c).
33
200 µm
Fig. 9 - Failure of the TBC system on a smooth sample creep tested up to failure at 1373 K (1100°C)/σmax = 160 MPa.
34
100 µm
Fig. 10 - Failure of the TBC system after 312 hours of the 1323 K (1050°C)/σmax = 160 MPa.
35
(a)
100 µm
(c)
(b)
20 µm 2 µm Fig. 11 - Failure of the TBC system after 42 hours at 1373 K (1100°C)/σmax = 160 MPa (a) and microcracking initiation at (b) pores and at (c) sand-blasting corundum particles.
36
50 µm
Oxide tickness (µm)
5
0.2E-04
4
0.2E-04
3
dε/dt 0.1E-04 (mm/h)
2 5.0E-06
1 0
0.0E+00 0
100
200
300
400
Test duration (h) Fig. 12 - Oxidation kinetics at 1373 K (1100°C)/σmax = 120 MPa as a function of the macroscopic displacement rate evolution.
37
TC
TGO
T and/or σ
BC
Creep strain
delamination
TC
BC
TGO delamination
~1%
Time Fig. 13 - Schematic illustration of the different damage mechanisms domains during viscoplastic solicitations of a TBC system.
38
Driving force for spallation, G
Driving force for spallation, G
Interfacial fracture toughnesses, GcTGO and GcBC/SX
Interfacial fracture toughnesses, GcTGO and GcBC/SX Gc-TGO
Gc-TGO
(a) Gc-BC/SX
(b) GBC/SX
Gc-BC/SX
GTGO GBC/SX GTGO
Creep time, t
Creep time, t
Fig. 14 - Schematic illustrations of the evolutions of the interfacial fracture toughnesses at the BC/SX (Gc-BC/SX) and at one of the TGO interfaces (Gc-TGO), and of the driving forces for spallation at the BC/SX or TGO interfaces for a high (a) and low (b) plastic strain rate creep test. Crosses highlight failures.
39
