UTC 2007
`General Electric v. United Technologies
`IPR2016-01289
`
`1
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`

`
`384
`
`CLARKE
`
`
`Turbine
`
`
`
`AirfoilGas,Ternperatu1'eCapability
`
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`COHVECEIOH 4“
`impingement
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`film
`mating \
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`convectimi
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`Coflling
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`Coo1ing+
`EB_PVD TBC
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`
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`wrought
`
`i940
`1960
`
`Llncooled Superalloya
`
`I 1980
`
`Year
`
`Increase in turbine airfoil temperature over the last six decades through com-
`Figure 1
`binations of materials advances and associated developments in cooling techniques.
`Since this diagram was constructed, the shaded region has extended to the present year,
`and the use of uncooled silicon nitride remains for the future.
`
`due to flame out, and as a means to even out local temperature gradients. Indeed,
`in some cases, the use of a TBC has simplified the design of blades by minimizing
`thermal distortions ofthe blade. However, undoubtedly the biggest benefit of TBCS
`has been to extend the life of alloy components in the hottest sections in an engine
`by decreasing their surface temperatures.
`Present day TBCs generally consist of a yttria—stabilized zirconia (YSZ) coat-
`ing deposited onto an oxidation—resistant bond—coat alloy that is first applied to a
`nicl<el—based sup eralloy component (Figure 2). In diesel engine applications where
`the temperatures are usually lower, the YSZ coating is generally applied directly
`onto the alloy. Two main types of coating are in use. For relatively small compo-
`nents such as blades and vanes in aerospace turbines, the coatings can be applied
`by electron—beam physical vapor deposition (EB—PVD). For larger components
`such as the combustion chambers and the blades and vanes of power generation,
`stationary turbines, the coatings are usually applied by plas1na—spraying (PS). In
`many respects, the choice of materials and their production represent a mature
`materials technology. While improvements in their capabilities continue, there is
`a growing realization that new TBC systems will be required for the next genera-
`tion turbines presently being designed. To set the stage for coming developments,
`we first review the selection of materials used in present YSZ coatings, some
`of the new insights that have been gained in understanding how YSZ coatings
`
`2
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`

`
`THERMAI %ARRTER COATINGS‘
`
`“I385
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`5%
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`£3
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`2 E 2
`
`Schematic illustration of a 1‘ SC
`Figure 2
`and the associated bond—coat on a superalloy
`in a thermal gradient.
`
`Superalltéji
`
`PF???‘
`Elmiriliiig Air
`
`fail, and then describe approaches to the development of the next generation TBC
`systems.
`
`PRINCIPAL REQUIREMENTS OF A THERMAL
`BARRIER COATING
`
`The turbine designers’ primary requirement of a TBC is that it have a low ther-
`mal conductivity and, for rotating components, preferably also a low density to
`minimize centrifugal loads. At the materials design level this translates into three
`additional requirements. First, the material must have strain compliance so as to
`withstand the strains associated with thermal expansion mismatch between the
`coating and the underlying alloy on thermal cycling. The use cycle, both the 1nax—
`imum temperature and the times at temperature, of course, varies between aircraft
`and power generation turbines, but nevertheless the coating must accommodate
`the large strains associated with thermal cycling. The need for strain compliance
`is illustrated in Figure 3, where the thermal expansion coefficients of zirconia,
`
`3
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`

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`386
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`CLARKP I LJVI
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` Q3E
`
`NJ3
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`
`
`ThermalExpansionCoefficient(ppmC"')
`
`
`
`
`
` 10
`1000
`
`Thermal Conductivity (WlmK)
`
`Figure 3 The thermal expansion coeflicients and thermal conductivity of
`a range of materials illustrating the differences in thermal expansion and
`conductivity of the principal components in TBC systems.
`
`alumina, and a number of alloys including nickel—based superalloys are cross-
`plotted against thermal conductivity. In the absence of any strain compliance, for
`instance due to a decreased elastic modulus, the large elastic mismatch would gen-
`erate Veiy large stresses and lead to spontaneous failure on cooling. Second, the
`coating material must exhibit thermodynamic compatibility with the oxide, usually
`aluminum oxide, formed on the bond-coat alloy at high temperatures. Third, with
`the continual quest to run engines at higher temperatures and the increasing dif-
`ficulty of increasing the metal temperature, it is increasingly likely that designers
`will seek “prime reliant” coatings, namely ones that can be used with assurance
`that they will not fail. Prime reliant thermal coatings are ones that are necessary
`to prevent the temperature of the metal from exceeding its maximum temperature,
`much in the same way that the tiles on the space shuttle prevent the underlying
`aluminum airframe from being exposed to temperatures in excess of their melting
`temperature on re~ent1y.
`Because Weight is at a premium in aircraft engines, thin coatings with the lowest
`possible thermal conductivity are required. In contrast, in stationary, ground—bas ed
`engines where weight is less of a consideration, a desired temperature drop can be
`
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`TH *RMAlI BARRIER COATINGS
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`387
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`achieved through simply increasing the '1' ’C thickness. In practice, in components
`in both types of engine, the thickness of the T 3C usually is varied from place to
`place to provide the desired thermal insulation.
`Jrosion of the coating by both ingested pa ticles, such as sand, from the op-
`erating environment and particles that come loose from the combustor liners as it
`degrades is a perennial source of concern, especially when the particles are large
`enough to cause impact damage of the coating. In some cases, inborn fine parti-
`cles, primarily dust and sand, melt into the coating as a wetting silicate while it is
`hot and can degrade the coating. These silicates, usually variants of Si-Al-Mg—Ca
`oxides that are the principal elements in sands, are often referred to collectively as
`CMAS.
`
`A recently recognized requirement of many materials exposed to high tempera-
`tures in gas turbines is a long-tenn stability in the presence of steam. This is partly
`a direct result of the generation of water during the combustion process, but in a
`number of designs it is a consequence of the use of steam injection to enhance
`turbine efliciency. Little is known about the effects of long-term exposure to steam
`on turbine materials. However, tests have revealed that many silicon—based com-
`pounds, including SiC, are unstable to the formation of volatile SiO, which results
`in the slow retraction of the material as evidenced by the reduction in thickness
`of components over long operating periods. This active oxidation and evaporation
`phenomenon precludes the use of silicon compounds in coatings unless protected
`by another coating.
`More difiicult to design against are the effects of corrosion, especially airborne
`species and those, such as sulfur and vanadium, in the fuel itself. The majority of
`land-based turbines operate on natural gas, but there is increasing interest in using
`alternative fuels, such as coal gas, that are much dirtier. The consequences ofusing
`such alternative fuels and their effects on coatings are only now beginning to be
`investigated.
`
`THE THERMAL BARRIER COATING SYSTEM
`
`From a materials engineering perspective, it is necessary to consider the TBC as
`an integrated materials system rather than simply a thermally insulating material
`coating on a structural alloy component. A representative cross-section of a com-
`mercial coating, shown in Figure 4, illustrates the multilayered nature of a coating
`after high—te1nperature exposure. There are three principal layers in addition to the
`superalloy and the low—conductivity coating. Between the alloy and the coating is
`the bond—coat, so called because in the initial development stages in producing a
`viable coating, it was found that the superalloy had to be first covered with a bond-
`coat to ensure that the YSZ coating remained adherent upon oxidation. Between
`the bond—coat and the YSZ coating—sometimes referred to as the overcoat—is the
`oxide formed during high—temperature exposure. Finally, during the formation of
`the bond—coat and the YSZ coating as well as subsequently during use, a reaction
`layer forms as a result of inter—diffusion between the bond—coat and the superalloy.
`
`5
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`CI ARI 4
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`T EVI
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`Figure 4 Cross—section of a TBC deposited by electron beam evaporation. Note the
`columnar microstructure of the zirconia coating. The white band is a reaction layer
`formed by interdiffusion during use between the Al-rich bond—coat above and the Ni-
`
`rich superalloy, below. The thermally grown oxide (TGO) is too thin to be discernible
`in this micrograph.
`
`In practice, as mentioned above, there are two distinct types of zirconia coat-
`ings reflecting different approaches to creating the strain compliance essential to
`withstand thermal cycling. The two types of coatings are EB—PVD coatings and
`plasma-sprayed coatings. In EB—PVD coatings, the lateral strain compliance re-
`sults from the columnar structure and inter—columnar gaps produced by rotation of
`the component during deposition. The columnar structure can be seen in the mi-
`crograph of Figure 4. Transmission electron microscopy reveals that the individual
`columns also contain microscopic porosity that reduces the thennal conductivity
`of the coating. In plasma-sprayed coatings, the lateral strain compliance and re-
`duced thermal conductivity is conferred by the incorporation of porosity between
`“splats” of successively deposited material. This porosity is illustrated in Figure 5,
`where the splats in a plasma spray bond coat have preferentially oxidized and
`consequentially appear as dark veins.
`Two major classes of bond—coat alloys have also evolved over the years, but
`both were developed to form an aluminum oxide (o¢—Al2O3) on exposure to air
`at high temperatures. This is important for several reasons. One is that A1203
`is phase compatible with YSZ, ensuring long—term theimodynamic stability of
`the coating. Uncoated, the majority of nicl<el—based superalloys form complex,
`
`6
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`TH ~RMAl %ARRIER COATINGS
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`389
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`
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`Figure 5 Cross—section of a TBC deposited by plas1na—spraying. The plate-
`like porosity is evident in the coating as the dark veins in the center of the
`micrograph.
`
`multilayered nickel oxide, nickel—chromiu1n spinels and chromium oxide, in addi-
`tion to alumina, and these are not thermodynamically stable with YSZ (2). Further-
`more, alumina is usually considered to be the slowest growing high—te1nperature ox-
`ide on account of it having the smallest oxygen diffusivity (3). The rationale for the
`selection ofbond-coat alloys is really a subj ect of another review, but the bond-coat
`has to perform a number of disparate functions. It must provide a bond between the
`deposited TBC and the underlying alloy. In the early days of TBC development, the
`bonding to the alloy was a major concern, particularly plasma—sprayed, hence
`the term bond-coat. Because zirconia is a fast—ion oxygen conductor, the bond-coat
`must also be able to form a protective, stable, and slow—growing oxide to prevent
`oxidative attack of the alloy. As is described below, one of the principal forms of
`failure is associated with failure of the protective aluminium oxide. The bond-coat
`must also have suflicient morphological stability so that on heating and cooling,
`as well as at high temperature, it does not distort and introduce incompatibilities
`that can also cause the introduction of interface defects.
`
`The two classes of bond-coat alloys that have been developed are the platinu1n—
`modified nickel aluminide (PtNiAl) and MCrAlY alloys (M here refers to one or
`more of the elements Co, Ni, and Fe). The selection of these two classes of alloys
`is largely based on their prior use as oxidation— and co1rosion—resistant coatings for
`protecting high—temperature alloys before the advent of TB Cs. For instance, the
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`390
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`CLARKE : LEVI
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`PtNiAl was originally developed as an alternative oxidation—resistant coating for
`protecting alloys at higher—temperature operation than the MCrAlY alloys available
`at the time.
`
`Different methods of applying the bond-coat alloys have been developed largely
`to meet production goals. Typically, PtNiAl bond-coats are formed by first electro-
`depositing Pt onto the superalloy component and then annealing it in an alu1ninu1n—
`rich vapor atmosphere. In this second step, aluminum diffuses into the surface of
`the alloy while nickel diffuses out where it reacts with the aluminum and platinum
`to form the PtNiAl aluminide coating. Depending on the quality of the coating
`required, the aluminum is provided in a pack—process or in a CVD reactor from an
`AlCl3 source. In contrast, the MCrAlY coatings are commonly deposited by one of
`a number ofvariants ofplasma—spraying. These processes are particularly attractive
`for coating large components and are, of course, cheaper than EB deposition. Also,
`as plasma-spraying does not involve a diffusion process, thicker bond-coats can be
`deposited than with the aluminizing process used to form the PtNiAl bond coats.
`It remains uncertain at this time which ofthese coating types is best for different
`applications. In large part this is because it is not yet known which combination
`of materials properties leads to the longest, high-temperature life of the coating.
`To provide the largest reservoir of aluminum one would expect that the thicker the
`bond-coat and the higher its aluminum content the better. However, one would also
`expect that the bond-coat should have as large a yield stress as possible at high tem-
`perature with as closely matched thermal expansion mismatch with the superalloy
`as possible to avoid thermal expansion mismatch stresses on thermal cycling.
`
`FAILURE MECHANISMS
`
`Investigation of the ways in which present YSZ coatings fail has provided consid-
`erable insight into the underlying mechanisms that limit coating life. Although, as
`with failure analysis in other areas of complex material systems, there are many
`complications, the findings nevertheless point toward methods of producing coat-
`ings that can withstand longer lives at temperature and higher use temperatures.
`One of the chronic problems is that the life of present TBC coatings invariably
`shows a wide distribution, with a high proportion of the population clustered about
`a median value but with a significant proportion failing at much earlier times.
`There is substantial circumstantial evidence to suggest that many of the TBC
`failures are associated with the oxidation of the bond-coat (4). Indeed, a number
`ofmanufacturers are believed to use an oxidation criterion as a basis for predicting
`average life. One such criterion is the combination of time and temperature to lead
`to a critical thickness of the TGO1. Another, embodied in the Coatlife software,
`is an aluminum depletion criterion based on the combined time and temperature
`
`1The concept of a critical thickness, of the order of 6 mn at ll00“C, appears to have
`originated from observations of coatings that failed under isothermal testing conditions.
`
`8
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`TH * RMAL BARRIER COATINGS
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`391
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`for the concentration of aluminum at the bond—coat surface to fall below a critical
`
`value. In the case of MCrAlY bond-coats, the rationale for this is that when the
`Al concentration falls to ~23 a/o, aluminum oxide is no longer the thermodynamic
`preferred phase and other oxides, notably spinels, form (5). These other oxides
`do not form such a protective scale, and consequently the alloy oxidizes faster. In
`addition, the formation of these oxides is associated with an increase in volume
`that can be disruptive and possibly have lower fracture energies, although this has
`yet to be unequivocally demonstrated. Nevertheless, there are reports that when
`the bond—coat is porous, and at low—oxidation temperatures, failure follows such
`aluminum depletion (6).
`Although related to the oxidation behavior of the bond—coat, neither the con-
`cept of a critical thickness nor aluminum depletion can account for the wide dis-
`tribution in failure lives, especially under thermal cycling conditions. Indeed, in
`the majority of materials examined after failures above about l000°C, the alu-
`rninum concentration, although depleted somewhat, has not fallen to the critical
`Value (7). Similarly, the short—lived coatings have failed before the TGO thickness
`has reached the thickness of its counterparts that have shown the longest lives.
`Together these findings indicate that failure occurs due to extrinsic factors arising
`during oxidation.
`The prevailing mode of failure is one in which part of the coating buckles and
`spalls away from the alloy, typically on cooling down to room temperature (8, 9).
`A typical buckling failure, in this case nucleated from the edge of a test coupon, is
`illustrated in Figure 6. Such buckling and subsequent spallation is a common mode
`of failure of all films and coatings under compression, generally associated with
`the development of compressive residual stresses in the coatings as a result of the
`difference in thermal expansion coefiicient between the coating and the underlying
`alloy. The mechanics of the failure by buckling of a thin, elastically isotropic
`film under compression from a flat surface is well understood (10), provided an
`unbonded region of a critical size, db, exists at the interface (Figure 7). For a fixed
`film thickness and residual stress, the stress at which buckling will occur is given
`by the relation:
`
`/E 48
`
`0 :. ——
`db
`
`h
`
`2
`
`.
`
`1
`
`.
`
`Thin film buckling is entirely analogous to the standard Euler buckling condition
`ofa column—a bifurcation phenomenon. The striking feature ofthis relation is that
`the flaw size depends linearly on the thickness of the film.2 Since even the thinnest
`of TBC is over l00 um thick, the critical size to which an interface flaw must grow
`before buckling can occur can be several millimeters. As interface separations of
`this large size are not usually present after coating, one of the major unresolved
`
`2Because they are designed to have in—plane strain compliance, TBCs are not usually
`isotropic elastic solids. The buckling condition is modified from that in Equation l
`to
`account for the elastic anisotropy.
`
`9
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`

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`392
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`CLARKE E T
`
`(Top) Incipient buckling of a TBC coating viewed under reflected light.
`Figure 6
`(Bottom) The failure surface revealed by spallation of the TBC consists of a mixture
`
`of local failure between the TGO and the bond—coat (appearing dark) and in the TBC
`itself (light regions).
`
`questions is how interface separations first form and then grow to such a large
`size. Such progressive failure consisting of nucleation of local interface separation
`and their subsequent growth has indeed been observed (1 l). Recent mechanics
`calculations have shown that interface perturbations from flatness can decrease
`the critical size at which buckles can initiate and then grow in size to form a spall
`
`10
`
`10
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`

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`TH *RMAll %ARRI * R COAIINGS
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`393
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`After;
`
`Substrate
`
`
`Figure 7 Schematic illustration of the buckling of a compressed film above
`
`a pre—existing defect of diameter db.
`
`(12). Nevertheless, localized flaws must first initiate and then grow for spallation
`failure to occur. Understanding the nucleation of these flaws, their growth, and
`progressive linking together is essential before realistic models for predicting life
`can be developed.
`Insight into the formation of flaws comes from microstructural examination of
`coating cross-sections after high—temperature exposure but prior to spallation. Four
`examples are shown in Figure 8, each from a YSZ TBC—coated PtNiAl bond—coat
`(13). In each case, the coating was deposited conformally onto the surface of a
`flat bond-coat so that the coating/bond—coat interface was initially flat and intact.
`As three of the inicrographs illustrate, the surface of the bond—coat roughens
`and separations form with the TBC even though the bottom surface of the TBC
`remains flat. These separations are the interface flaws that progressively grow in
`size and link together with adjacent ones to allow buckling and spallation (9, ll).
`Roughening is more pronounced with thermal cycling but also occurs, albeit more
`slowly, on isothermal exposures (14). The growth ofthese separations with thermal
`cycling can now be monitored by luminescence piezospectroscopy, as described
`
`11
`
`11
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`

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`mM..u
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`W!) ‘H Jsmrn
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`12
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`12
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`

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`TH *‘RMAlI BARRIER COATINGS
`
`395
`
`in the next section, which suggests that it has the potential to be used as a viable
`N )J:’ tool (1 1, 15).
`The micrographs in Figure 8 raise the question as to the underlying mechanisms
`responsible for the observed roughening. At least two new mechanisms have now
`been identified that can lead to such roughening. The roughening has been at-
`tiibuted to a “ratcheting” phenomenon motivated by the lateral compressive stress
`in the growing TGO and facilitated by thermal cycling (9). Measurements indicate
`that as the TGO grows in thickness with oxidation, it also concurrently develops
`a compressive stress (l6, 17). If it were free to expand it would decrease its coin-
`pressive stress but because it is attached to the bond-coat, the only way in which
`it can decrease its elastic strain energy is by undulating (Figure 9). In this way, its
`length increases and it remains attached to the alloy. This undulation requires the
`alloy to deform to accommodate the undulation, and the oxide must also deform
`concurrently. According to the ratcheting mechanism, this accommodation is by
`plastic deformation of both the TGO and bond-coat during thermal cycling. As
`the lateral growth of the thickening oxide continues during the high-temperature
`portion of the thermal cycles, it continues to generate compressive stress that is
`relaxed by ratcheting during the thermal cycle so the process is ongoing. Many of
`the essential features of the mechanism have been substantiated by finite element
`computations (18) and are consistent with observations of the increase in length
`of the TGO as the surfaces roughen.
`Another new mechanism shown to cause roughening is the surface displace-
`ment associated with volumetric changes in the bond-coat as aluminum depletion
`occurs. This roughening is illustrated in Figure 10, together with etched cross-
`sections revealing the presence of both y’ and ,6 phases in the bond-coat (14).
`After aluminizing and after YSZ deposition, the PtNiAl bond-coat is chemically
`homogeneous and has the ,8—NiAl (B2) crystal structure. After high—temperature
`exposure, the initially flat bond-coat is rumpled and etching reveals that the bond-
`coat has partially transformed to y ’ —Ni3Al. In addition, the remaining fl —NiAl phase
`regions often have the characteristic lath structure of a maitensite. These two ob-
`servations can be understood as being the result of aluminum depletion from the
`bond-coat and concurrent enrichment of nickel from the underlying superalloy, a
`classic example of interdiffusion. This change in composition is illustrated using
`the binary NiAl diagram in Figure 11. As aluminum is depleted, the average com-
`position of the bond-coat becomes increasingly enriched in nickel until reaching
`
`< F
`
`igure 8 Cross-section of four TBCs illustrating different forms of the local sep-
`aration between the TBC and the bond-coat after thermal cycling. In each case, the
`coating was deposited conformally on the bond-coat so interface separations such as
`these indicate that the underlying bond-coat has changed its surface morphology dur-
`ing high—temperature exposure and thermal cycling. In example ([9), no separation has
`occurred and it exhibits the longest life.
`
`13
`
`13
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`

`
`396
`
`CI ARK *
`
`L *‘Vl
`
`
`
`Schematic illustration of how an initially flat but compressed film (left)
`Figure 9
`can lower its elastic strain energy by rumpling (right). The amplitude of rumpling
`is enhanced by thermal cycling and can cause interface separation if a superimposed
`coating cannot deform to follow the displacements of the film.
`
`
`asdemsitédiii
`
`I after cyclic oxirlationf, x
`“j,1200”C,“5Bx_1hr
`r
`a
`
`Figure 10 Microstiucture of an initially flat as-aluminized bond coat after 50 1-11
`cycles at l200°C: (a) surface rumpling; (b) cross—section showing a rather uniform
`oxide layer and strong surface undulations (y’—phase is revealed by etching); (c, (1')
`optical micrographs showing etched cross—section before and after cyclic oxidation.
`Dark areas on the optical images coirespond to the ;3—phase, whereas the y’-phase in
`the coating appears white.
`
`14
`
`14
`
`

`
`TH RMAT ARKIER COATINGS
`
`397
`
`Composifort in wt.% Ni
`
`0
`
`was
`
`20
`e
`
`e
`
`40
`
`u
`
`so
`
`so
`
`100
`.
`
`1800
`
`1460
`
`1200
`
`inC
`Temperature
`
`Al
`
`Composition in at.% Ni
`
`Ni
`
`Figure 11 The Ni—Al pseudo—binary phase diagram illustrating the compositional
`range of the ,8—NiAl phase and the direction of the change in composition as the bond-
`coat is depleted of Al by interdiffusion and selective oxidation.
`
`the single—phase boundary at which point fiirther depletion leads to the formation of
`1/—Ni3Al. (At even later times, the composition can extend into the y region of the
`phase diagram.) Martensitic structures within the ;3—NiAl phase field also form as
`the phase boundary is approached. Whereas the martensite start temperature, Ms,
`of the pure ,8—NiAl compositions is known to be generally around room tempera-
`ture to 300°C (20), the additional Pt, Co, and Cr present in the PtNiAl bond-coat
`increase the Mg temperature, and Hemker et al. have reported Mg temperatures of
`~600°C (21).
`Substantial progress has been made in the past few years in understanding some
`ofthe mechanisms that lead to flaw initiation and growth during use (9, 14, 18, 22).
`These provide the basis for developing life—prediction models, but nevertheless a
`number ofunanswered questions remain. For instance, the micrographs in Figure 8
`were obtained from nominally the same superalloy, with the same bond-coat and
`the same YSZ coating all made by the same manufacturer in the same process
`manner. This difference in interface separation and roughening is particularly
`marked in this figure, but it does suggest that even small, but as yet unidentified,
`concentrations of dopants can have a large effect on life (23).
`Insights gained in the past few years into some of the important processes
`occurring within the coating during use are summarized in Figure 12. Essentially,
`
`15
`
`15
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`

`
`398
`
`CLARKE I LEVI
`
`L
`
`'
`
`L
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`""* L
`,
`rireepsandjalsne
`densifies by eintering
`L
`L
`A Phase lransformaticun
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`Q; .
`Ira TGO during uxldattiszni
`Emile? interface
`
`pmsétien a A
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`JA‘
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`Local mughnees grows.
`by growth strain induced
`Qxgde gmws in
`ratcheting and bond
`Wckmsa amj
`carat vclume decreases
`expands Iaterauy
`
`Elxidatioh causes
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`Al-depleticzn of
`I‘_)()nd—{;u:aa?(
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`N
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`LLLL
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`_N
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`i, W’ etc
`
`Localized Densification
`Due to [3——;s3:‘Trans'forn1afinn
`
`baweeu
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`Figure 12
`
`Schematic summary of the concurrent processes occurring in the bond-
`
`coat, TGO and TBC, during use at high temperatures. The complexity in failure times
`and failure modes is believed to reflect the competition between hese individual pro-
`cesses.
`
`the TBC system is one that evolves with time at temperature and its evolution
`depends in detail on not only the temperature but also on the thermal cycle history
`and heating and cooling rates, as well as composition.
`
`NON—DESTRUCTIVE EVALUATION
`
`As the design of coatings shifts to a philosophy of prime reliance, the ability to
`non—destructively monitor the coating, identify defects, and evaluate its remaining
`life becomes more important. In addition, there is a growing economic pressure to
`defer maintenance and replacement of parts until really necessaiy. (The costs are
`staggering: The cost, in replacement electricity alone, oftaking a power generation
`turbine out ofoperation can be ofthe order of $1 million a day. The cost ofreplacing
`a single, first—stage turbine blade can also be very high. Depending on its size the
`cost can be as high as $10,000.)
`In the case of large area separations, several millimeters to centimeters, infrared
`imaging provides a direct means of visualizing incipient coating failure provided
`the blade can be accessed. Usually, though, a coating has failed by the time the
`separation reaches such a large size, and thus methods of identifying damage at
`an earlier time, and hence smaller size, are required. No single solution appears
`practical at this stage, although laser-induced acoustic sounding (from Lasson
`
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`THERMAL BARRIER COATING S
`
`399
`
`Technologies, personal communication) and higher—spatia resolution imaging us-
`ing polarized scattered light have shown promise in detectng deliberately created
`interface flaws (25). An alternative method is one that utilizes piezospectroscopy,
`the strain—induced shift of luminescence and Raman lines. When illuminated with
`
`a laser having an appropriate wavelength, luminescence from the aluminum oxide
`TGO fonned by oxidation on the bond—coat can be detected through the thickness
`ofthe TBC (26, 27). The luminescence spectrum, from Cr3+ ions incorporated into
`the alumina TGO as it grows, is sensitive to the stress state in the TGO. (Because
`zirconia is a wide band—gap material, it is transparent in the visible so illu1nina—
`tion in the blue or green, e.g., as an argon ion laser, can penetrate to the TGO,
`and the stimulatedluminescence, which is in the red, is transmitted back through
`the coating.) The key to the use of photoluminescence as an NDE tool is that the
`frequency of the luminescence lines shifts with mean stress and the shape of the
`luminescence lines is a direct measure of the stress distribution within the region
`ofthe TGO probed by the laser. As described in detail elsewhere, the luminescence
`lines can be deconvoluted and the degree of damage and local interface separation
`evaluated (1 1, 15). Since this can be performed using focused lasers and is also
`nondestructive, piezospectroscopy—based methods show particular promise as a
`NDE tool for evaluating coatings. An example is illustrated in Figure l3a,b, that is
`a comparison of spectral data from two coatings. In Figure 13a, there little change
`in spectral shift with thermal cycling, the TGO remains flat and does not fail within
`the time of the experiments (curve A in Figure 130) In Figure 131), a continuous
`change in frequency (curve B in Figure 13c) is shown that exhibited substantial
`spectral broadening and failed after several hundred cycles, an average value but
`shorter than the other.
`
`SINTERING AND DENSIFICATION
`
`One of the concerns in developing a reliable and robust coating is how the coating
`changes during use at high temperatures. By analogy with the behavior of other
`porous ceramics, it might be expected that the coating will densify by the reduction
`in surface energy associated with the excess surface area of the pores. This process
`is commonly referred to as sintering. There are two principal concerns. One is
`that densification inevitably increases the elastic modulus and thereby decreases
`the strain compliance of the coating. The other is that densification decreases
`the volume fraction of porosity which, in tum, causes the thermal conductivity
`to increase. Although the kinetics of densification of bulk, free—standing zirconia
`ceramics is well characterized, evaluation of these effects in actual coatings is
`not. In part this is because of the peculiar geometry of the porosity in plasma-
`sprayed and electron-beam deposited coatings. In a bulk ceramic, porosity tends
`to be spherical in shape, whereas it is plate—like in plasma—sprayed coatings and
`columnar-like in EB~PVD TBCs. These differences in pore shape and topology
`are compounded by the fact that the densification of a TBC is constrained by the
`presence of the underlying alloy, which has not yet been investigated in detail.
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`TBCS might be quite different from one another and also from bulk zirconia
`ceramics on account of the constraint imp