On the failure mechanisms of thermal barrier coatings with diffusion aluminide bond coatings

Materials Science and Engineering A 394 (2005) 176–191

On the failure mechanisms of thermal barrier coatings with diffusion aluminide bond coatings I.T. Spitsberga,∗ , D.R. Mummb , A.G. Evansc a

GE Aircraft Engines, Materials and Processes Engineering Department, 1 Neumann Way, MD M 69, Cincinnati, OH 45215, USA b Department of Chemical Engineering and Materials Science, University of California, Irvine, CA 92697-2575, USA c Materials and Mechanical Engineering Departments, University of California, Santa Barbara, CA 93106, USA Received 6 August 2004; accepted 15 November 2004

Abstract The mechanisms governing the failure of multi-layer thermal barrier systems based on Pt-modified nickel aluminide bond coats and electron beam deposited thermal barrier coatings (TBCs) have been studied. The primary experimental variable is the morphology of the bond coat surface prior to application of the TBC, at constant multi-layer chemistry. The durability of these systems in a furnace cycle test has been measured and compared. The failure mechanisms, as well as the thickening of the thermally grown oxide (TGO), have been characterized for each of the surface morphologies. The major findings are that the durability is enhanced by removing imperfections on the surface of the bond coat, as well as by surface pre-treatments that diminish subsequent TGO thickening and by incorporating a reactive-element in the substrate that strengthens the bond coat upon inter-diffusion during manufacture. These effects are consistent with the expectations of a TGO instability mechanism, driven by a combination of growth and thermal expansion misfit strains in the TGO. The grain structure of the bond coat also affects failure through its influence on the TGO instability sites. © 2004 Elsevier B.V. All rights reserved. Keywords: Thermal barrier coatings; Durability; Interface morphology; Thermally grown oxide; Displacement instability; Plastic anisotropy; Pre-oxidation

1. Introduction Many recent articles have described the role and benefits of thermal barrier systems in gas turbines. These systems comprise the tri-layer: bond coat, thermally grown oxide and thermal barrier [1–6]. While much has been learned about these layers and about their evolution with exposure, several fundamental issues remain to be understood. One broad challenge is to develop a comprehensive, mechanistic understanding of coating durability, at a level sufficient to develop mechanism-based models. Such models should prove useful in guiding practical approaches to fabricating coatings with improved thermo-mechanical performance, as well as for life prediction. Present evidence suggests that there are at least five different failure mechanisms [4], dependent on the bond coat chemistry and microstructure and the method used to ∗

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0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.11.038

deposit the thermal barrier layer. The mechanism to be addressed in this article is that operative in a system comprising an electron beam physical vapor deposited (EB-PVD) thermal barrier coating (TBC) and a single-phase, Pt-modified βNiAl bond coat (PtNiAl), otherwise referred to as a platinum aluminide bond coat. Even in this one system, the dominant failure mechanism depends upon whether the operating conditions are predominantly isothermal or cyclic. The latter is studied here. When this system is subject to thermal cycling, the thermally grown oxide (TGO) undergoes a displacement instability, manifest as local penetrations of the TGO into the bond coat at periodic sites along the interface, as previously observed by Ruud et al. [7]. The instability progresses on a cycle-by-cycle basis (Fig. 1) [8,9]. The downward displacements induce normal strains in the superposed TBC that cause cracks to form, extend laterally and, eventually coalesce to induce failure by large-scale buckling (LSB) [4,8,9]. The comparative area fractions of materials exposed

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after delamination implies that cracks extend with similar facility in the TBC, through the TGO, and at the two bi-material interfaces (TBC/TGO and TGO/bond coat). A more detailed discussion of crack path selection, and its implications for durability, follows. The penetration of the TGO into the bond coat must be accommodated by visco-plastic flow from the base of the penetration toward the interface. Accordingly, the yield/creep strength of the bond coat must be important (Fig. 2). The TGO increases in length, as well as thickness, at the instabilities [9]. Additionally, the bond coat experiences phase transformations. A spatially non-uniform, isothermal β → γ  phase change occurs with an associated volume change [10,11]. The β-phase undergoes a martensite transformation upon thermal cycling with a volume strain of ∼0.7%. [12]. Models characterizing aspects of this instability have been developed [8,13–18]. They reveal that many different material properties (over 20) influence the rate of undulation. The present study designs a series of experiments to test

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these ideas on failure mechanisms and to provide approaches for enhancing durability through morphological and microstructural improvements. Previous analyses have indicated that grit-blasting pre-treatments may be detrimental through incorporation of impurities and contaminants that accelerate the oxidative process [19]. In this study, systematic approaches are developed to isolate the implications of various grit-blasting procedures on the initial morphology of the bond coat surface. Additional factors affecting TBC life also emerge.

2. Experimental protocol 2.1. Materials The experiments were conducted on single-phase (modified) β-NiAl bond coat systems, provided by several manufacturing sources. These were made by electrolytically

Fig. 1. Failure of a TBC system, driven by a cyclic instability in the thermally grown oxide layer [9]: (a) schematic illustrating the thermo-mechanical response of (Ni,Pt)Al-based thermal barrier coatings, indicating the development of a displacement instability, the associated crack nucleation and growth in the TBC, and crack coalescence leading to coating spallation; (b) cross-sectional images detailing the micro-structural and morphological evolution of the TBC system at various stages of life. The imperfections that ratchet and the resulting crack nucleation in the TBC are evident after only 34% of the cyclic life. Both grow with further cycling.

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Fig. 1. (Continued ).

depositing Pt onto a superalloy, then using chemical vapor deposition (CVD) to introduce Al, and finally heat-treating to create the β-phase by inter-diffusion. The thickness, as well as the Al and Pt content of the bond coats varied slightly, depending on the manufacturing source, with the typical ranges being 50–60 ␮m for the thickness, 17–21 wt.% and 25–28 wt.% for Al and Pt, respectively (see Section 5.3 for more details on the chemical composition). The TBC was a standard yttria-stabilized zirconia deposited by EB-PVD. Most series of measurements were conducted using Ren´e N5 single-crystal substrates (developed and commercialized by GE Aircraft Engines, Cincinnati, OH). In one series, an alloy containing reactive elements was used in place of N5. This

Fig. 2. Predictions of a model of the ratcheting instability [8]: (a) a ratcheting diagram which uses the initial imperfection aspect ratio,  = Ao /L, as the bc abscissa
and the stress ratio,  ≡ σo /σy as the ordinate. It shows the range of , space in which ratcheting can be expected for a given growth strain per cycle; (b) a plot of predicted ratcheting rates as a function of the lateral growth strain per cycle, computed for various values of the bond coat yield strength, σybc [12].

alloy has a composition similar to N5, but with enhanced reactive element additions. The nominal composition of Ren´e N5 is: 8 wt.% Co, 7% Cr, 2.0% Mo, 5% W, 7% Ta, 3.0% Re, 6.2% Al, 0.2% Hf, and Ni (balance). 2.2. Surface treatments After forming the bond coat, but before depositing the TBC, the surfaces were subjected to a range of treatments (Table 1), resulting in the following five series of experiments:

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Table 1 Summary of experimental conditions Processing method and substrate

Surface preparation procedure(s)

Identifier

Process 1/N5 Process 2/N5

As processed As processed

Series A

Process 1/N5

Series B

Process 2/N5

Grit blast (medium grit) Grit blast (fine grit)

Process 2/N5

Polished

Series C

Process 2/N5

Polished + grit blast (coarse grit) Polished + grit blast (medium grit) Polished + grit blast (fine grit)

Series C

Process 2/N5

Grit blast (coarse grit) + pre-oxidation

Series D

Process 2/RE-Sub

Grit blast (coarse grit)

Series E

four performed with N5 substrates and one with a reactive element containing substrate similar to N5. (i) For series A, the bond coat was retained in its asmanufactured state, with continuous ridges present where the grain boundaries intersect the surface [20,21]. Using bond coat specimens from various manufacturing sources (with slightly different process conditions) resulted in two such surface conditions, differing in the bond coat grain size (and, therefore, in the size and separation of the ridges), as well as the initial surface chemistry. The system with the larger grain size is referred to as “process 1” and that with the smaller grain size, “process 2”. (ii) In series B, the surface was given a “grit-blast” treatment, typical of that used in commercial practice. That is, alumina particles with coarse, medium or fine grit were propelled at the surface, at high velocity, to remove and distort a thin surface layer. (iii) Series C experiments were conducted on surfaces that had been mechanically polished to flatten the surface by removing the ridges. In one variant, series C , polishing was followed by grit blasting with alumina particles with fine, medium or coarse grit. (iv) For series D, the surfaces were grit blasted and heattreated in a controlled atmosphere of low partial pressure oxygen at temperatures greater than 2000 ◦ F (1093 ◦ C) to form a thin (approximately 0.1–0.3 ␮m) TGO of αAl2 O3 before depositing the TBC. Pure α-Al2 O3 was formed in all cases, as verified by TEM and A¨uger profiling techniques. (v) Series E used RE-containing substrates, with surfaces prepared using the baseline (grit-blasting) treatment. 2.3. Testing and characterization approach In all cases, test specimens comprising 25 mm diameter disks were used. After preparing each of the five series, the surfaces of the bond coat were characterized using the

atomic force microscope (AFM) as well as by optical and scanning electron microscopy (SEM). In some cases, A¨uger microscopy was used to establish the chemistry of the near surface region. The TBC was then deposited and the thermal cycling tests performed. “Standard” furnace cycle test (FCT) conditions were used. This test comprised 1 h thermal cycles between ambient and a temperature greater than 2025 ◦ F (1121 ◦ C), with a rapid heat up (∼1 min), a 45 min hold at temperature, and a 14 min cooling period. The number of thermal cycles is reported in this paper as the fraction, f, of the nominal coating life, Nf (for example, N0.75 indicates 75% of the average cyclic life). Failure is defined as a condition wherein more than 20% of the surface area exhibits spalls. For the series A, C, D, and E, at least 10 specimens were tested. For the series B, a large population was used: combining data generated for the present study with GE Aircraft Engines historical test data. Some tests were interrupted and cross-sections made using procedures described elsewhere [9]. After sectioning, the specimens were comprehensively characterized using combinations of field-emission-gun scanning electron microscopy (FEG-SEM), energy dispersive X-ray spectroscopic mapping (EDS), electron microprobe wavelength dispersive spectroscopy (WDS) and Auger spectroscopic analysis. The thickness of the TGO was determined from the crosssections, either at various fractions of life, or recorded at failure. The procedures used in preparing samples for crosssectional examination and characterization are critical to the correct assessment of damage evolution and the appropriate identification of coating failure mechanisms. Because of this criticality, the sample preparation procedure is elaborated with a particular focus on two aspects: (a) The choice of the resin used to embed the sample for grinding and polishing is paramount. The shrinkage encountered upon curing of common thermoset resins induces stresses of sufficient magnitude to delaminate the

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TGO/bond coat interface: with misleading inferences regarding failure mechanisms. Low viscosity, slow curing resins, specifically designed to minimize shrinkage, are essential. Even then, the temperature of the exothermic cure reaction must be carefully controlled to minimize the curing/shrinkage rate. (The necessity for stress avoidance is exacerbated by the fact that a mode I stress intensity is induced by the shrinkage and that the mode I toughness of the metal/oxide interface is low relative to the mode II toughness). (b) The procedures used for final polishing prior to examination are also crucial. Standard procedures generally involve final polishing on a soft cloth with colloidal sil-

ica. The authors have noted a vast difference in the interface integrity when the final polish is conducted with diamond lapping films and pH-neutral water, instead of commercial colloidal silica solutions, which are not pHbalanced. That is, there is greater tendency toward interface de-adhesion when the colloidal silica solution is used. This tendency is attributed to the susceptibility of this interface to stress-corrosion. 3. Surface topologies Representative optical and SEM images of the different surfaces are summarized on Fig. 3a–h. AFM analysis of the

Fig. 3. Pairs of optical and scanning electron microscope (SEM) images illustrating the surface morphology resulting from the different bond coat surface preparation conditions: including as-processed, polished and grit-blasted surfaces. Each optical image is acquired at equivalent magnification, and similarly for the SEM images.

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Fig. 3. (Continued ).

associated roughness is summarized on Fig. 4a–h. The asprocessed bond coat surfaces exhibited the grain boundary ridges described by others [20,21] (Figs. 3a,b and 4a,b), but with two different grain sizes achieved through process modifications. The grain boundaries are generally located at the base of the ridges, on one side. The “shadow” lines reflect the migration of the grain boundaries during the coating growth process. Note that ridges are not in evidence at all grain boundaries and that there is “etchpit” contrast on some of the grains. The AFM measurements (Fig. 4a,b) indicate that the ridges are sometimes facetted. On average, process 1 coatings have ridges with height ∼3 ␮m and width ∼20 ␮m. For process 2, the

ridges are much smaller, with height ∼0.5 ␮m and width ∼3 ␮m. The surfaces formed by coarse grit blasting are indicated on Figs. 3c and 4c. After blasting, the surface of the process 1 coating is indistinguishable from process 2 materials, even when viewed at high magnification. There is no evidence of the original surface ridges. The surface comprises a continuous landscape of ‘cuts’ and imprints caused by the impacting alumina particles. The peak-to-valley roughness amplitudes are in the range of 5–8 ␮m, based on large-area AFM measurements. The corresponding surfaces after blasting with intermediate and fine grit (Figs. 3d,e and 4d,e) reveal that the ridge

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Fig. 4. Atomic force microscopy (AFM) plots of the measured surface profiles resulting from the surface preparation conditions detailed in Fig. 3.

structure from processing is partially retained. That is, while the ridge height has been diminished from the as-processed level, the ridge structure is still clearly in evidence. Superposed on the diminished ridges is a pattern of imprints and cuts from the alumina impacts. The latter are diminutive versions of the pattern observed with the coarse alumina. The surfaces created by polishing are revealed on Figs. 3f and 4f. The ridges have been removed and the surface has a near-optical quality finish. Fine scratches remain after polishing. Note that some of the grains give optical contrast, indicative of polishing relief governed by the hardness anisotropy of the β-grains. Blasting after the surfaces had been polished (Figs. 3g,h and 4g,h) gave the same pattern of cuts and imprints noted above, but now superposed on a planar surface. Accordingly,

the coarse grit resulted in a surface indistinguishable from that produced by the grit-blasting treatment applied to the as-processed surface (with appearance as in Figs. 3c and 4c): whereas the surfaces treated with finer media differed from the corresponding as-processed and blasted surfaces, because of the absence of the remnant ridges.

4. Durability tests Test results for series B (Fig. 5) revealed that variations in the alumina particle size used for blasting did not influence the durability, within typical batch-to-batch variability, even though the amplitude of the blasting-induced roughness changed by a factor 3 (see Section 3). Moreover, specimens

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Fig. 5. Failure lifetimes for TBC systems, series A through series E, relative to a ‘baseline’ system with a standard grit-blast surface.

with as-deposited surfaces, series A, exhibited similar durability (Fig. 5). (For some conditions used to process the bond coat, the durability was appreciably lower due to surface contamination effects discussed below). Series C specimens, with fully polished bond coat surfaces, exhibited cyclic life 2 to 5 times greater than that for the baseline. It will be shown that this large increase coincides with a mechanism change. Partially polished surfaces gave intermediate durability. Surfaces that had been polished and then grit-blast, series C , exhibited a duality in life. Those subject to coarse particle blasting were comparable to the baseline (series A treated with coarse blasting media): whereas specimens experiencing medium and fine particle treatments gave enhanced life, comparable to the as-polished series C. The pre-oxidation treatment, series D, also caused a systematic increase in durability, by a factor of 2 to 3 relative to the baseline, without a mechanism change. It is concluded from these four series that, for the same material system, the durability is strongly affected by the surface morphology and chemistry present before the TBC is deposited. Changing the substrate to the reactive element-containing alloy, series E, resulted in durability 3 times greater than the baseline with N5 substrates, even though the surface treatments were identical. Moreover, it will be demonstrated that failure occurred by a different mechanism (relative to the baseline behavior). The remainder of this article will attempt to elucidate the mechanistic basis for these effects.

series (Fig. 6), reveal that, while the growth rates are about the same for series A and B systems, they are appreciably lower for the pre-oxidized, series D, specimens, as well as the polished, series C samples, although the latter exhibit much greater variability. All such measurements are based on samples incorporating Ren´e N5 substrates. At failure, the average thickness of the TGO was 5–6 ␮m for series A, B and D, and about 8 ␮m for series C. The TGO in series E specimens (with reactive element containing substrates) contained thickness heterogeneities associated with entrained reactive element based oxide (REoxide) precipitates (Fig. 7). Such microstructural features are analogous to the Y-based precipitates found in TGO layers

5. Characterization 5.1. TGO thickness The trends in TGO thickness with hot time, examined using selected specimens representative of each experimental

Fig. 6. Thickness of the TGO layer determined from cross-sections either at various fractions of life (filled symbols), or recorded at failure (open symbols), for materials prepared using the surface preparations detailed in Fig. 3.

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Fig. 7. Cross-sectional image of the TGO/BC interface at failure for a series E specimen. Note the reactive-element based oxide precipitates entrained in the TGO layer, and the associated morphology at the metal/oxide interface.

grown on NiCoCrAlY bond coatings [22]. At failure, the average thickness of the TGO was 7→9 ␮m. Note that, despite the entrained RE-oxide precipitates the overall TGO growth rate is nearly identical to that for the precipitate-free TGO in the baseline materials. 5.2. Isothermal phase changes Regions of the bond coat that transform from β → γ  during cyclic exposure are apparent on cross-sections (Fig. 8a–c) generated at various fractions of life. The γ  domains have the same topological characteristics noted by others [6,9–11]. They originate as “triangular” regions, which enlarge with time-at-temperature (Fig. 8), typically occurring at the intersections between the β-grain boundaries and the interface with the TGO (though not in evidence at all such grain intersections). Most of the transformed domains connect with narrow (0.1 ␮m) γ  zones along the bond coat grain boundaries. The distribution of γ  indicates that the structure and microchemistry of the grain boundaries play a controlling role in the transformation to γ  during cyclic oxidation. Note, however, that there is no discernable difference in the thickness of the TGO located above β and γ  domains. In series C specimens, the β → γ  transformation occurring at the intersections between the TGO and the bond coat grain boundaries initiates at shorter times, and the γ  -domains are more distinctive (see Fig. 8c). 5.3. Micro-chemistry Prior to TBC deposition, A¨uger analysis of the asprocessed bond coat surfaces revealed that (for some manufacturing processes) the surface contaminants Ca and S were present in concentrations up to 1.7 atomic %. For other processes, these contaminants were at levels below the XPS resolution limit. Tests performed on the former, with as-processed surfaces, resulted in deficient durability: failing prematurely

Fig. 8. Cross-sectional images showing the relationship between transformed γ  domains and the displacement instabilities along the interface: (a,b) large-area survey at an intermediate stage of life in samples prepared via standard grit-blasting procedures; (c) γ  domain formation at an earlier stage of life in a sample where the bond coat had been polished prior to TBC deposition.

by delamination along the TGO/bond coat interface. These tests have not been included in Fig. 5. The results reported are those based on processing procedures wherein Ca and S are not present in detectable quantities.

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Table 2 Composition of the γ  and β-phases of the bond coat as a function of the number of cycles Phase

Percentage of life

Ni

Cr

Co

Ta

W

␥

50 100

55.45 58.81

2.99 3.51

5.50 5.75

5.43 4.40

1.89 1.99

9.56 9.74

18.8 14.7

␤

0 15 50 90 100

46.80 44.23 47.02 49.91 51.50

3.38 3.85 3.95 3.98 4.01

4.08 4.26 4.55 4.41 4.64

0.54 0.99 076 0.72 0.60

0.06 0.24 0.21 0.23 0.17

17.1 15.08 14.86 14.72 16.05

26.8 30.7 26.9 27.1 22.0

Fig. 9. Distribution of aluminum in the bond coat as a function of hot time, measured for three separate specimens with a standard grit blast surface. Note that, in all cases, the Al decreases to about the same level after approximately 10 h, substantially prior to failure.

Micro-chemical analysis by electron microprobe (Table 2) of cross-sections produced at various fractions of life has revealed substantial Ta in the γ  , as well as moderate amounts of W: a consequence of diffusion up from the substrate. These elements, which are known to be γ  stabilizers, are present in only trace quantities in the β-phase. Accordingly, because of the uniformity of the TGO thickness, these elements must have minimal effect on the TGO growth mechanism. The Al concentration in each phase appears to be insensitive to distance from the TGO and, on average, decreases to a level consistent with that at the γ  /βphase boundary after about 10 h at the peak temperature (Fig. 9). 6. Failure mechanisms 6.1. Synopsis Because of the multiplicity of observations made on each of the experimental series, a synopsis is presented first, followed by pertinent details. Five factors encapsulate the predominant effects of morphology, surface chemistry and grain structure on the cyclic durability measured with Ren´e N5 substrates.

Al

Pt

(i) Initial non-planarities act as imperfections that initiate displacement instabilities [7]. Removal of these imperfections, by polishing, inhibits instability formation [20,21], changing the failure mechanism to interface delamination, which happens at TGO thickness (h0 = 8 ␮m) greater than the baseline case (h0 = 6 ␮m). For instabilities to form, the imperfections must exceed a critical size, estimated as Ac ≈1 ␮m. For imperfections exceeding Ac , shape is more important than size. (ii) The location and spacing of the dominant instabilities, as well as their growth rates, are about the same for a wide spectrum of initial imperfections: ranging from intrusions caused by grit blasting to ridges present after manufacturing. The spacing of the dominant instabilities correlates with the bond coat grain structure. In many cases, the instabilities originate where triple grain junctions in the bond coat intersect the TGO. (iii) In cases where failure is driven by the TGO displacement instability, spallation occurs whenever the downward displacement of the dominant imperfections equals about H/2, where H corresponds to the initial thickness of the β-phase region within the bond coat. (iv) System durability is dominated by cracks evolving in the TBC and at the TBC/TGO interface. Coalescence along the TGO/bond coat interface constitutes only a small fraction of the life. (v) The interface between the TGO and the bond coat degrades with exposure due to kinetic processes, such as the segregation of impurities. This degradation has negligible effect on life when the TGO instability dominates, but becomes important when the instabilities are suppressed and failure occurs by edge delamination mechanisms. The measurements performed on specimens with reactive element containing substrates have another implication. (vi) By changing the properties of the bond coat, the TGO instability is suppressed, even in the presence of surface imperfections, resulting in a transition to failure by edge delamination. The consequence is greater durability, and a capacity to sustain a thicker TGO. It is not known whether the effect is caused by changes in strength of the bond coat or by a reduction in the martensite start and finish temperatures.

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6.2. TGO instabilities Cross-sections of series B and D specimens revealed periodic TGO instabilities that propagate into the bond coat, having the same characteristics described elsewhere [9] (see Fig. 1). They exhibit a variety of morphologies, ranging from circular to square to triangular, dependent upon process conditions for the bond coat and the local bond coat microstructure. The squares and triangles always have corners where a bond coat grain boundary intersects the TGO, indicating that the microstructure of the bond coat plays a prominent role in the development of damage. The TGO has essentially the same thickness at the instability sites as elsewhere along the interface and may be cracked at the corner. At failure, the average amplitude of the largest penetrations was about 20 ␮m, with spacing approximately 60 ␮m. Note that the average thickness of the TGO was similar in all cases, even though series D had durability about twice that for the baseline (Fig. 5). For series A, the TGO instabilities are again in evidence. The shape distortions that accompany the instability, vividly demonstrated on Fig. 10, indicate that the displacements are coincident with the pre-existing ridges. This judgment may be made because the grain structure of the superposed TBC is a strong indicator of the prior bond coat morphology. Namely, ‘pinched-off’ grain columns develop at concave features and ‘fanned-out’ columns at convex segments [9]. The image also demonstrates that the ridges displace upward, consistent with other observations [15], as well as simulations [15]. The TGO remains relatively planar in the zones between the ridges. Series C and E specimens retained a planar TGO up to failure. That is, there are no deeply penetrating instabilities. The only detectable off-planar features are incipient instabilities apparent in series C tests, highlighted on Fig. 8c. 6.3. Failure topology and crack path selection Series A, B and D specimens all failed by large scale buckling (LSB): whereas, series C and E specimens generally failed by delamination of the bond coat/TGO interface. The exceptions were series A specimens with contaminated surfaces, which delaminated prematurely. The spalled surfaces of series A and B specimens exhibited a mixed failure pathway, with regions of TBC, TGO, and bond coat exposed (Fig. 11a and d). The failure of series D specimens occurred entirely at the TGO/TBC interface (Fig. 11b). For series A and B, the fraction of exposed surface comprising bond coat, designated fbc , ranged from less than 0.05 to about 0.4, with no obvious correlation between fbc and the durability. In cases where fbc was relatively large, “oxide islands” remained embedded in the bond coat. These islands had the TBC at the center and the TGO around the perimeter. At a more detailed level, a somewhat larger proportion of series A than series B failures occurred almost exclusively at the TBC/TGO interface and the “islands”, when present, appeared to exhibit greater pla-

Fig. 10. Cross-sectional image of a series A specimen at failure, vividly illustrating development of upward-moving displacement instabilities and their relationship to initial convex features at the bond coat surfaces (grain boundary ridges). Note that the interface between the ridges remains relatively planar.

narity (Fig. 11c). A rationale for this observation is presented below. The cross-sections of series B specimens revealed the characteristics reported elsewhere [9]. Namely, all of the cracks originate either in the TBC or at the TBC/TGO interface at locations above the TGO instability sites (Fig. 12). Their initial growth, either in the TBC or within the TBC/TGO interfacial zone, comprises the majority of the cyclic life. Subsequent growth patterns during the latter stages of life appeared to be quite variable. In some cases, the cracks continued to extend either exclusively in the TBC or at the TBC/TGO interface (see Fig. 1b). In others, they diverted through the TGO and coalesced along the intervening TGO/bond coat interface, particularly when relatively planar. These findings highlight the potentially misleading nature of post-failure surface observations. They also rationalize the insensitivity of the durability to the fraction of bond coat exposed, since failure along the TGO/bond coat interface only

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Fig. 11. Failure surfaces of: (a) grit-blasted series B, (b) pre-oxidized series D, and (c,d) as-received series A samples. The images shown in (c) and (d) illustrate the wide variability in the fraction of exposed bond coat for nominally identical samples.

Fig. 12. Cross-sectional image showing the nucleation of cracks in the intermixed zone of the TBC/TGO interface, above the instability sites. Note the convolution of the TGO and the location of the γ  domains (the darker phase) in the bond coat.

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happens during the final stage of life: whereupon the specific crack trajectory is quite sensitive to the local geometry and to the presence of a buckling-induced tensile (mode I) loading at the crack tip. For series A, the cracking originates in the TBC near the ridges. In the latter stages of life, cracks are more likely to extend along the TGO/bond coat interface than in series B specimens, even though the durability is about the same. Series C and E specimens generally failed by a delamination crack extending almost exclusively along the TGO/bond coat interface: exposing the bond coat and leaving the TGO attached to the TBC (Fig. 13). In some areas, small domains of TGO were left embedded in the exposed bond coat [22]. This mode of failure is governed by de-adhesion of the metal/oxide interface, motivated by the strain energy density in the TGO and TBC. Series C had the duality noted above. Namely, the polished specimens subsequently treated with coarse grit exhibited TGO instabilities, as in series B. Conversely, treatments with finer grits resulted in failure by interface delamination.

Fig. 13. Images of the (a) top surface and (b) cross-section of polished, series C, specimens, illustrating how removal of morphological imperfections suppresses the TGO displacement instability, resulting in edge delamination with cracking occurring at the TGO/BC interface.

7. Correspondence with models Correspondence of the proceeding findings with models of the TGO displacement instability is addressed within the expectations of the undulation mechanism [8,13–18]. Ratcheting simulations reveal that imperfections on the bond coat surface are critically important. They suggest a critical imperfection size, below which the mechanism is suppressed [13–18], and a major role of the aspect ratio above the critical size. The ensuing experimental findings are consistent with these predictions. Namely, removal of imperfections by planarizing the bond coat surface suppresses the instability, leading to a mechanism change and enhanced durability. Small imperfections superposed onto these surfaces by blasting with fine grits (Figs. 3g,h and 4g,h) have insufficient size to reactivate the instability and, thereby, retain greater durability. All surfaces with larger imperfections (series A and B) experience instabilities and exhibit lower durability. Note that the similarity in the size and shape of ridges in series A to the imprints in baseline (coarse grit) series B specimens is consistent with their comparable effects on instability growth [15], elaborated below. Series B specimens treated with either medium or fine grit particles, with remnants of the as-processed surface ridges still present, retain dominant imperfections of about the same size (above critical) and, accordingly, have durability similar to series A. The failure modes in series A and B specimens correlate with two different manifestations of ratcheting. For series B, the concave impressions displace downward (Fig. 1b) [9], resulting in highly localized strains, with cracks that initiate in the TBC above the ratcheting sites, extend laterally and coalesce. Conversely, for series A, the cyclic plastic displacement of the ridges proceeds in an upward direction, away from the interface (Fig. 10) [15]. The observed crack trajectories, and the fraction of exposed bond coat at failure, are dependent upon the specific shape of the ridges present above the bond coat grain boundaries, as illustrated in Fig. 14. If the ridges have a small radius of curvature at the apex, with a relatively smooth transition to the surface away from the ridge, the displacement instability will propagate in the fashion shown in Fig. 14a, with the TBC effectively being pushed away from the substrate. Such displacements cause the intervening interfaces to experience quite uniform tensile vertical strains that increase in magnitude as the displacements accumulate. The spatial uniformity of the strains in these regions facilitates cracking between the ridges, along the trajectory with lowest toughness (which may be the TGO/BC interface for interfaces embrittled by segregation during extended thermal cycling). Conversely, for ridges that approach a hemispherical shape, the root of the ridge acts as a concave morphological feature susceptible to displacement downward into the bond coat (while the apex of the ridge is relatively stable). In this case the displacement instability results in highly localized tensile stresses, above the TGO layer, resulting in cracking within the TBC or at the TGO/TBC interface. The stochastic nature of the ridge shape gives rise to the varied

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spatial correlation between the locations of the instabilities, their spacing and the β-grain boundaries.

8. Failure by edge delamination The incidence of edge delamination can be addressed through assessments of the steady-state energy release-rate, Gss , relative to the effective toughness of the metal/oxide interface, Γ i . For TBC/TGO bi-layers, Gss is given by [24,25]: Gss =

M=

(E1 h1 (1 + υ1 )ε1r + E2 h2 (1 + υ2 )ε2r )2 Mκ −   2(E1 h1 + E2 h2 ) 2

E1 E2 h1 h2 (h1 + h2 )[(1 + ν2 )ε2r − (1 + ν1 )ε1r ] 2(E1 h1 + E2 h2 )

6[ε2r (1 + υ2 ) − ε1r (1 + υ1 )] κ = h1

ξ= Fig. 14. Schematics illustrating the effect of the shape of the grain boundary ridges on the propagation of displacement instabilities. In (a), where the ridges have a small radius of curvature at the peak and smooth transitions to the surrounding surface, upward-vectored ratcheting displacements occur resulting in relatively uniform normal strains in the TBC and along the interface between ridges. The consequence is a mechanism whereby, failure may occur at the TGO/BC interface. In (b), where the ridges have a larger radius of curvature at the peak, but have sharp transitions to the surrounding surface, the instability propagates downward resulting in more localized normal stresses that preference crack nucleation above the TGO.

crack trajectories evidenced by the mixed failure observed in Fig. 11d, and may explain the lack of correlation between the fraction of exposed bond coat and the durability. Note that ridge shapes representative of Fig. 14a and b are both present in the AFM measurement shown in Fig. 4a. The evolution of the TGO morphology is often strongly dependent upon the microstructure of the bond coat (see Fig. 15). In their simplest form, the instability models do not predict the strong association between the predominant instabilities and the bond coat grain structure. However, a spatial correspondence is clearly present [11]. However, by starting with geometric imperfections at the junctions between the bond coat grain boundaries and the TGO, the models predict a more rapid development of groves at the grain boundaries than of undulations within the grains [23]. The grain structure has another, secondary, role related to the preference for γ  formation at β-phase grain boundaries (see Fig. 8). Instability development is locally suppressed at these domains, as noted in cross-sectional images (Fig. 12), since γ  is significantly stronger than β [9]. Consequently, there is an indirect



1 + h 1 / h2 ξ

(2a)

(2b)

 (2c)

        E1 h1 2 E2 h2 2 h1 h2 + + 4 + 4 +6 E2 h2 E1 h1 h2 h1 (2d)

where h1 and h2 are the thickness of the TGO and TBC layers, respectively, ε1r and ε2r are the residual strains in each layer, and E1 and E2 are the respective plane-strain moduli, Ei = Ei (1 − υi2 ). Equating the steady-state strain energy release rate, Gss , to the appropriate measure of the interfacial toughness, Γ i , and re-arranging the above equations, allows for determination of a critical TGO thickness. The relevant modulus for the TBC is that parallel to the substrate, taken as, E2 ≈ 20 GPa [26]. While this may change with time due to sintering, the value chosen is regarded as representative. The effective interface toughness Γ i , is that appropriate to an interface, degraded by thermal cycling, because of segregation and the accumulation of small interface separations. Estimates made by various measurement approaches suggest a value of order, Γ i ≈ 80–100 J/m2 [22,25,27,28]. The critical TGO thickness is predicted from (2), upon using this range of toughness values and the constituent properties listed elsewhere [19], as hc ≈ 6 ␮m. This thickness is comparable to the TGO thickness measured at failure in series C and E specimens.

9. Some implications The effect of pre-oxidation (series D) on durability has particularly interesting consequences, since at failure, the TGO thickness and the penetration depth of the TGO into the bond coat are the same as in the benchmark systems:

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Fig. 15. Images illustrating the strong correlation between the development of the TGO displacement instabilities and the underlying grain structure: (a) ratcheting occurs within a domain between grain boundaries; (b) each of the primary instability sites is correlated with a bond coat grain boundary.

despite the factor two increase in durability. The incontrovertible implication is that pre-oxidation reduces the TGO growth rate without adverse changes in other properties. The reduced growth rate requires that the diffusivity of reactants in the TGO growth process be reduced by about a factor 4 (assuming parabolic growth). The reduced growth rate is most likely caused by a larger TGO grain size [29]. This is consistent with the fact that transport of the reacting species through the TGO occurs predominantly along grain boundaries. A larger grain size may arise because of differing nucleation conditions during the initial stages of TGO growth, resulting in fewer transport pathways across the TGO layer per unit area. Further characterization is needed to test this assertion. A similar effect is observed with polishing, series C. Again the TGO thickness at failure is similar to that observed with the baseline system, despite a pronounced difference in exposure time (Fig. 6). A microstructure having fewer through-thickness transport pathways would also result in slower lateral oxide growth rates, consistent with the observation that the amplitude of the displacement instabilities (as well as the tortuosity of the TGO layer) is about the same for the series B and series D samples: again, despite the factor two difference in cyclic exposure. Recall that the displacement instability mechanism requires a lateral component of the TGO growth to ‘feed’ the cyclic evolution of the TGO layer morphology, and that the instability displacement rate scales with the overall lateral growth rate [8].

10. Conclusions A study has been conducted of the mechanisms governing the failure of multi-layer thermal barrier systems based on (Ni,Pt)Al bond coats and electron beam deposited thermal barrier coatings (TBC). A major finding has been that the durability is affected by four separate factors. (i) Imperfections on the bond coat surface are critical. Removal of imperfections by polishing suppresses the

TGO instability and enhances durability. The failure mechanism changes to edge delamination. (ii) The TGO thickening rate is important. A pre-oxidation heat treatment that diminishes the thickening rate increases the durability, even though failure occurs subject to the same TGO instability mechanism, at the same final TGO thickness. (iii) The bond coat properties are a factor. Changing the substrate to an alloy that includes reactive elements results in inter-diffusion during manufacture, causing the properties of the bond coat to change. The resultant properties suppress the TGO instability, resulting in superior durability. It is not known whether the effect is due to a strengthening of the bond coat or a reduction in the martensite start and finish temperatures.

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