- Email: [email protected]
Failure characteristics and mechanisms of EB-PVD TBCs with Pt-modified NiAl bond coats
Author's Accepted Manuscript
Failure Characteristics and Mechanisms of EBPVD TBCs with Pt-Modified NiAl Bond Coats Le Zhou, Sriparna Mukherjee, Ke Huang, Young Whan Park, Yongho Sohn
www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(15)00380-9 http://dx.doi.org/10.1016/j.msea.2015.03.120 MSA32217
To appear in:
Materials Science & Engineering A
Received date: 9 September 2014 Revised date: 4 March 2015 Accepted date: 27 March 2015 Cite this article as: Le Zhou, Sriparna Mukherjee, Ke Huang, Young Whan Park, Yongho Sohn, Failure Characteristics and Mechanisms of EB-PVD TBCs with Pt-Modified NiAl Bond Coats, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2015.03.120 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Failure Characteristics and Mechanisms of EB-PVD TBCs with Pt-Modified NiAl Bond Coats Le Zhou, Sriparna Mukherjee# , Ke Huang&, Young Whan Park%, Yongho Sohn* Advanced Materials Processing and Analysis Center and Department of Materials Science and Engineering University of Central Florida, Orlando, FL, 32816, USA
#
Now with Department of Chemistry, North Carolina State University, Raleigh, NC 27607, USA &
%
Now with Siemens Energy, Inc., Orlando, FL 32826, USA
On Sabbatical from Department of Mechanical Engineering, Pukyung National University, Busan, Republic of Korea
* Corresponding author, Professor, Department of Materials Science and Engineering, University of Central Florida, Orlando, FL 32816, USA; Email: [email protected]; Voice: 407.882.1181
Key Words: Thermal barrier coatings, (Ni,Pt)Al bond coat, failure, microstructure, transmission electron microscopy
1
Abstract Microstructural evolution and failure characteristics/mechanisms were investigated for thermal barrier coatings that consist of electron beam physical vapor deposited ZrO2 - 8 wt.% Y2O3 (YSZ) topcoat, Pt-modified nickel aluminide, (Ni,Pt)Al bond coat, and CMSX-4 superalloy substrate with furnace cycling at 1100°C with 1-hour dwell. Photo stimulated luminescence spectroscopy, scanning electron microscopy equipped with X-ray energy dispersive spectroscopy and transmission electron microscopy were employed to examine the residual stress of the thermally grown oxide (TGO) and microstructural changes. For comparison, (Ni,Pt)Al bond coat on CMSX-4 without the YSZ topcoat was also characterized. The TGO grew faster for the YSZ-coated (Ni,Pt)Al bond coat than the (Ni,Pt)Al coating without the YSZ topcoat. Correspondingly, the β-to-γ’/martensite formation in the (Ni,Pt)Al bond coat occurred faster on the YSZ-coated (Ni,Pt)Al bond coat. However the rumpling occurred much faster and with larger amplitude on the (Ni,Pt)Al coating without the YSZ topcoat. Still, the rumpling at the TGO/bond coat interface caused crack initiation as early as 10 thermal cycles, decohesion at the YSZ/TGO interface, and eventual spallation failure primarily through the TGO/bond coat interface. The magnitude of compressive residual stress in the TGO showed an initial increase up to 3~4 GPa followed by a gradual decrease. The rate of stress relaxation was much quicker for the TGO scale without the YSZ topcoat with distinctive relief corresponding to the cracking at the top of geometrical ridges associated with the (Ni,Pt)Al bond coat. The maximum elastic energy for the TGO scale was estimated at 90 J/m2 at 50% of its lifetime (Nf = 545 cycles). The YSZ presence/adhesion to the TGO scale is emphasized to minimize the undulation of the TGO/bond coat interface, i.e., decohesion at the YSZ/TGO scale accelerates the rumpling and crack-coalescence at the TGO/bond coat interface where the spallation fracture occurs.
2
1. Introduction Thermal barrier coatings (TBCs) are multilayered coatings consisting of a ceramic topcoat, a metallic bond coat and a thermally grown oxide (TGO) that protected the underlying hot component from high temperature. The ceramic topcoat typically consists of 7-8 wt.% Y2O3 stabilized ZrO2 (YSZ) that is processed via electron beam physical deposition (EB-PVD) or air plasma spray (APS). The metallic bond coat is usually diffusion aluminide or MCrAlY (M=Ni, Cr or both) overlay coatings. The TGO layer forms between the ceramic top coat and metallic bond coat during high temperature operation due to oxidation [1]. Among various TBCs, Pt-modified NiAl aluminide coatings topped with EB-PVD YSZ ceramic topcoats are commonly used because of excellent spallation resistance. The incorporation of Pt into NiAl aluminide is believed to improve the oxidation behavior by strengthening alumina scale adhesion, slowing diffusion of detrimental element and retarding the void growth at the interface [2-5]. Understanding the mechanism that governs the failure of TBCs helps in the material selection and design of TBCs. TBC failure mechanisms depend on the type of TBCs and the operational conditions. For the EB-PVD YSZ coated on (Ni,Pt)Al bond coat under thermal cycling conditions, surface rumpling [6,7] or ratcheting [8] of the bond coat can occur, and cause instabilities of the interface between the TGO and bond coat. The rough interface may give rise to out-of-plane tensile stress that leads to the spallation or delamination of the TBCs. Furthermore the strain energy of the TGO accumulates during cyclic oxidation, and failure can occur when the strain energy within the TGO exceeds a certain critical value [9]. In this paper, commercial-production TBCs consisting of EB-PVD YSZ topcoat, (Ni,Pt)Al bond coat, and CMSX-4 superalloy substrate were furnace cycled at 1100°C without temperature gradient within the specimens. The TBC specimens were fabricated so that one side was coated only with bond coat, and the other side was coated with YSZ topcoat with (Ni,Pt)Al 3
bond coat. Thermal cycling lifetime of TBCs was initially assessed. Then, as a function of thermal cycling (i.e., fraction of lifetime), compressive residual stress, phase constituents, and microstructure were examined by photo-stimulated luminescence (PSLS), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Characteristics of spallation failure are documented and discussed with respect to the failure mechanisms.
2. Experimental details All EB-PVD TBC examined were disk-shaped specimens with diameter of 25.4 mm and thickness of approximately 3 mm. One face-side of the disk was coated with 175 ȝm-thick EB-PVD 8 wt.% Y2O3 stabilized ZrO2 (8YSZ) ceramic topcoat on top of a 45 ȝm-thick grit-blasted (Ni,Pt)Al bond coat. On the other face-side of the disk, only the 45 ȝm-thick (Ni,Pt)Al bond coat was coated. The superalloy substrate for all samples was CMSX-4 with the nominal composition, in weight %, of Ni9.0Co-6.5Cr-6.5Ta-6.0W-5.6Al-3.0Re-1.0Ti-0.6Mo-0.1Hf. Thermal cycling tests were carried out using CMTM Rapid High Temperature furnace. Each cycle consisted of a 10-min heat up to 1100°C, a 1-hour dwell at 1100°C, followed by a 10-min forced air quenching to room temperature. Thermal cycling tests were performed until the failure of TBC defined as the spallation of YSZ (more than 10 area %) from the substrate (Nf = 545 cycles). All TBC specimens were periodically removed from the furnace for the measurement of average compressive stress by photo-stimulated luminescence spectroscopy (PSLS) using a RenishawTM System 2000 Ramanscope. Measurements were performed from the topcoat for the YSZ-coated side, and from the ridges and valleys for the bond-coat-only side. Twenty random measurements were carried out to obtain the average value along with the standard deviation. 4
Several TBC specimens, prior to spallation-failure, were also removed from the furnace after thermal cycles of 10, 50, 100, 200, 300, 400, and 500 for microstructural analysis. Morphological evolution of the bond-coat-only surface was examined by ZeissTM Ultra 55 Field-Emission SEM equipped with X-ray energy dispersive spectroscopy (XEDS). Each specimen was first mounted in epoxy, sectioned by low speed diamond saw, metallographically ground, and polished down to 0.25ȝm followed by chemical etching using HCl/HNO3 (5:1) etchant. Cross-section of each specimen was examined by optical microscopy (OlympusTM LEXT OLS 3000 CSM) and SEM with XEDS. Transmission electron microscopy (TEM FEITM Tecnai F30 300 keV) equipped with high angle annular dark field (HAADF) detector was employed to examine the details of the YSZ/TGO and TGO/bond coat interfaces for the as-coated specimen and specimens subjected to 10 and 200 thermal cycles. The TEM foils were prepared via site-specific focus ion beam (FIB FEITM 200) in-situ lift-out (INLO) technique. Bond coat rumpling was quantitatively measured through a series of cross-sectional SEM micrographs. The interface between the TGO and the bond coat was traced and quantified. The tortuosity (L/L0) that is defined as the ratio of undulated TGO/bond coat interfacial length to the initial length (i.e., straight) was determined. By considering the oxide particles between the bond coat and the superalloy substrate as inert markers, the thickness of the bond coat was determined with root mean square (RMS) value as given by:
RMS =
1 n 2 yi − yo ) ( ¦ n i=1
[1]
where n is the number of points, yi is the length between the YSZ topcoat and the particles, y0 is the mean height of the bond coat.
5
3. Microstructural analysis 3.1 As-received thermal barrier coatings Figures 1(a) and 1(b) present, respectively, the HAADF and bright field (BF) TEM micrographs of the as-coated EB-PVD TBCs, where a uniform TGO layer of ~350 nm was observed between the YSZ and grit-blasted ȕ-(Ni,Pt)Al bond coat. The selected area diffraction pattern (SADP) from one of the alumina grains is shown in Figure 1(c), and indicates that the TGO layer primarily consists of the ĮAl2O3. Metastable Al2O3 frequently found and reported [10-13] was not observed to develop on this grit-blasted ȕ-(Ni,Pt)Al bond coat. The crystal structures of the YSZ and the ȕ-(Ni,Pt)Al were also confirmed by SADPs to be tetragonal and B2, respectively, as represented in Figures 1(d) and 1(e). Both HAADF and BF TEM micrographs show that the YSZ/TGO and TGO/bond coat interface does not have any visible decohesion and/or voids. 3.2 Thermal growth of the oxide scale with thermal cycling The average TGO thickness, along with the standard deviations, for the thermally cycled specimens was measured from a series of cross-sectional backscatter electron (BSE) micrographs and plotted as a function of square root of time as shown in Figure 2. The linear relationship in Figure 2 demonstrates that the TGO growth follows the diffusion-controlled parabolic growth. Accordingly the parabolic growth constant Kp was calculated according to the equation:
Y = Yo + K p t
[2]
where Yo and Y are initial and instantaneous thickness of the TGO, and t is the oxidation time. The calculated growth constants for the YSZ-coated and bond-coat-only sides were 0.241 ȝm/h1/2 and 0.175 ȝm/h1/2, respectively. The TGO underneath the YSZ topcoat grew at a faster rate than that exposed to 6
air. This would be a result of higher partial pressure of oxygen with the YSZ at high temperature than that in the ambient atmosphere (e.g., consider oxygen content in YSZ and air). Furthermore, The Al2O3 breaks into the YSZ and forms small equiaxed oxide grains, as will be seen in Figure 4(b). This formation of intermixed oxide zone was pointed out to promote the outward diffusion of Al [14][15], probably due to more grain boundary paths for Al diffusion, which might increase the oxidation rate. Additionally, it is noted that the bond coat of the YSZ-coated side is grit-blasted while the bondcoated-only side is not. Previous research has shown that the impurities introduced by grit-blasting increases the diffusivity of Al and thus the oxidation rate [15].
3.3 Phase transformations in the ȕ-(Ni,Pt)Al bond coat Figure 3 presents cross-sectional optical micrographs that illustrate the phase evolution within the (Pt,Ni)Al bond coat. The grain boundaries are revealed by chemical etching. For the specimen subjected to 50 cycles in Figures 3(c) and 3(g), white-contrast phase start to appear in the bond coat, preferably at the grain boundaries. These were Ȗ’-Ni3Al phases as will be seen later. With oxidation, the Ȗ’-Ni3Al phase grows as shown in Figures 3(d) and 3(h) due to the Al depletion via TGO formation and substrate-bond coat interdiffusion. Figures 3(d) and 3(h) shows that the growth of Ȗ’ phase appears to be faster for the bond coats with YSZ topcoat. Martensite formation in the bond coat is revealed by the lamellar microstructure with etching. Although the bond coat initially consists of the ȕ phase, the formation of martensite can occur when Al content falls below 37 at.% [17]. The martensite lamellae are observed only after 10 thermal cycles on the side coated with the YSZ as shown in Figure 3(b), and after 50 thermal cycles for the side coated only with (Ni,Pt)Al bond coat. In this study, the transformation to both Ȗ’ and martensite was observed earlier for the YSZ-coated side. This is consistent with faster Al consumption from the faster TGO growth observed on the YSZ-coated side, 7
and conforms that, indeed, the outward diffusion of Al for the formation of Al2O3 is faster when YSZ topcoat is present. Two TEM specimens were prepared from the interfacial area between the TGO and the bond coat for specimens with 10 and 200 thermal cycles. The SADP shown in Figure 4(e) confirms the presence of Ȗ’-Ni3Al, which formed due to the depletion of Al in the bond coat, which initially consisted of the ȕ-NiAl solution phase. Figures 4(d) and 4(f) show the microstructure of the martensites, and the corresponding SADP from the boundary area of the martensitic variants. These variants form twinned microstructure with each variant having a thickness of a few hundred nanometers.
3.4 Rumpling and/or racheting in the ȕ-(Ni,Pt)Al bond coat Figure 5 presents morphological evolution on the top surface of the bond coat without EB-PVD YSZ. Initially, the TGO forms a needle- or plate-like morphology, which is frequently associated with the formation of metastable alumina at the beginning thermal cycles [18]. With high temperature oxidation, the TGO transforms to the stable Į-Al2O3 with a network-like morphology, however with some cracks on top of the ridge area only after 50 thermal cycles, as shown in Figure 5(c). As the thermal cycling continues, the ridges become more prominent and the cracks propagate along the ridges as presented in Figures 5(d) and 5(e). Further thermal cycling leads to connection of these cracks along the ridges and some crack nucleation within the flat/troughs region of the TGO as presented in Figure 5(f). Cross-sectional BSE micrographs of the TGO scale on the YSZ-coated side are presented in Figure 6. The separation in Figures 6(e) and 6(f) is due to metallographic preparation. Initially, a uniform TGO layer forms and attaches to the bond coat tightly. The TGO rumples following the 8
contour-evolution of the bond coat. With increasing thermal cycles, some decohesions at the TGO/YSZ and TGO/bond coat, particularly around the grooved regions of the bond coat, are observed as shown in Figure 6(d). Although the asperity of the bond coat remains relatively fixed, many grooves (e.g., inward oxidation) form after prolonged thermal cycles, as shown in Figures 6(e) and 6(f). Figure 7 present the SEM micrographs from the TBC sample failed by 545 thermal cycling. Figure 7(a) shows the exposed bond coat surface, and Figure 7(b) shows the bottom of the spalled YSZ, Figures 7(c) and 7(d) are the cross-sectional BSE micrographs of the YSZ-spalled specimen and spalled YSZ, respectively. Much of the TGO scale remained attached to the spalled YSZ, and that observed on the bare bond coat was mostly located in the groove area as represented in the inset in Figure 7(a). Circled region in Figure 7(c) also demonstrate the residual TGO attached to the bond coat exists in the groove areas. Although decohesion between the YSZ and TGO is quite clear, much of the TGO scale was observed to be attached to the spalled YSZ. Spallation failure with fracture at the TGO/bond coat interface despite a significant decohesion damage at the YSZ/TGO interface has been observed previously [19]. This failure mechanism has been qualitatively described based on degraded ability of the YSZ to suppress the buckling of the TGO scale due to the decohesion damage at the YSZ/TGO interface [20]. A direct view of the rumpling/racheting behavior of the bond coat as a function of thermal cycling is presented by the contour of the TGO/bond coat interface for the bond-coated-only side and the YSZ-coated side as plotted in Figures 8(a) and 8(b), respectively. For each specimen, a distance of more than 2500 ȝm was examined. The RMS and L/L0 for the bond-coated-only and YSZ-coated side were also calculated and plotted in Figure 9 as a function of thermal cycles. The obvious gain in the amplitude of the contour indicates that the bond coat surface rumpled accumulatively with increasing thermal cycles. The amplitude increases faster for the side with bond coat only. Moreover, the bond9
coated-only surface develops both upward and downward contour change, while only grooving is generally observed for the side with YSZ topcoat. This is consistent with bond coat instability for the YSZ-coated side since it is appropriately termed ratcheting that does not require the upward displacement of the bond coat as in the rumpling behavior observed for the bond-coated-only side [8,9,21,22]. The presence of YSZ topcoat would place a rigid restriction on the instability of the bond coat surface, and can cause a preferred formation of grooves only. A small decrease in L/L0 for the YSZ-coated side was observed at the beginning cycles as shown in Figure 9. Although not substantiated, it could be related to the partial recovery of the severe plastic strains induced by grit blasting. The rumpling and/or ratcheting at the TGO/bond coat interface was observed as a significant damage-factor causing the spallation failure, since the growth of the TGO scale on a wrinkled surface can develop out-of-plane stress. Based on curvature, the stress is tensile at the ridge region and compressive at the groove region, which causes separation between TGO/bond coat interface [23]. Figures 4(a) and 4(b) present the HAADF micrographs of the interfaces of the specimens after 10 and 200 thermal cycles, respectively. Figure 4(c) is the BF TEM micrograph of the interface for the specimen with 10 thermal cycles, highlighting the nucleation of cracks at the TGO/bond coat interface. It is recalled from Figure 1 that no cracks were observed in the as-coated specimen. These cracks would coalesce and grow with further thermal cycling, as observed in Figure 4(b) for the specimen with 200 thermal cycles. The separation between the TGO and the bond coat eventually leads to the failure of the TBC. The damage at the TGO/bond coat interface is the primary reason for final failure, while the degradation via decohesion at the TGO/YSZ interface reduces the suppression ability of the YSZ against buckling of the TGO. In addition, impurities introduced by grit blasting may degrade the fracture resistance of the TGO/bond coat interface [16]. 10
4. Compressive residual stress in the thermally grown oxide The average compressive residual stress in the TGO scale was measured by PSLS as a function of thermal cycles as presented in Figure 10. The open circles represent the magnitude of average compressive residual stress in the TGO underneath the YSZ topcoat. The squares and cross represent TGO stress of the flat area and the ridge regions, respectively, for the bond-coated-only side. In general, all three average compressive residual stress increase slightly initially, followed by continues decrease with thermal cycling. A peak magnitude of approximately 3 GPa is observed for all curves, which corresponds to that determined by thermal expansion mismatch. The initial increase in TGO stress can be caused by the TGO growth or the transformation from metastable Al2O3 to stable Į-Al2O3 [24,25]. A gradual release of the TGO stress can be attributed to the surface rumpling or ratcheting of the bond coat [6,8] or creep of the bond coat [26]. The magnitude of average compressive residual stress in the TGO underneath the YSZ topcoat remains slightly higher than that of the TGO developed on the overlay (Ni,Pt)Al bond coat. Faster TGO growth (and thus the higher growth stress [27]), and lower rate of undulation for the TGO/bond coat interface potentially due to the suppression by the YSZ topcoat, may help retain the higher magnitude of compressive residual stress within the TGO underneath the YSZ topcoat. On the other hand, the magnitude of compressive residual stress within the TGO on top of the ridges observed on the bond-coat-only side decreased rapidly after 50 cycles, consistent with formation of cracks along the ridges as seen in Figure 5.
11
5. Discussion Experimental findings can be correlated quantitatively and qualitatively to provide some insight into the failure mechanisms of TBCs examined in this study. The elastic energy accumulated in the TGO scale on cooling is considered as the main driving force for the final TBC failure [28,29]. In a simplified model, the TGO elastic energy per unit area can be defined [28,29] to contain two parts: strain energy within the TGO and TGO/bond coat interfacial energy, as expressed by:
TGO energy(J / m 2 ) =
σ 2 (1− υ ) h L E
(
Lo
)2 + γ M /O (
L 2 ) Lo
[3]
where ı is the magnitude of compressive residual stress in the TGO, Ȟ is the Poisson’s ratio (~0.25), h is the TGO thickness, E is the TGO elastic modulus (~370 GPa), L/Lo is the tortuosity and ȖM/O is the interfacial energy at the TGO/bond coat interface (~7 J/m2) [22]. By substituting the measured values for ı, h and L/L0 from this study, the TGO energy was calculated as a function of thermal cycles as plotted in Figure 11. The first term on the right-hand-side of the Eq. [4], strain energy within the TGO, is much larger in magnitude than the second term in Eq. [3], the TGO/bond coat interfacial energy. With an increase in thermal cycles, the TGO energy continues to increase to about 90 J/m2. After 400 thermal cycles, a sudden decrease in the TGO energy is observed, corresponding to a partial, yet a significant damage in the TGO scale or at the TGO/bond coat interface. With failure, the decrease in the TGO energy would correspond to the relief of the energy stored in the scale. Previous studies indicate that phase transformations in the bond coat can contribute to the rumpling of the bond coat. The ȕ-to-Ȗ’ phase transformation is associated with a significant molar volume increase of approximately 8% to 38% [7], and the martensitic transformation can induce a transformation strain of 0.7% due to a 2% molar volume difference between the martensite and the ȕ 12
phase [30,31]. In this study, the transformation to Ȗ’ and/or martensite was observed earlier for the YSZ-coated side along with the faster depletion of Al due to faster TGO growth. The undulation of the bond coat surface, however, occurred faster for the bond-coated-only side than the YSZ-coated side as seen in Figure 9. Therefore, the undulation of the bond coat surface is strongly influenced by the adherence of the YSZ coating to the TGO scale, and the formation of Ȗ’ and martensites does not appear to be the only factor on the overall rumpling behavior of the bond coat underneath the YSZ topcoat. 5. Summary Microstructural evolution and failure characteristics of EB-PVD YSZ TBCs with Pt-modified NiAl bond coat were examined as a function of furnace thermal cycling at 1100°C with dwell time of 1 hour using PSLS, XRD, SEC/XEDS, and TEM. To examine the influence of EB-PVD topcoat during thermal cycling, the other side of the button specimen was only coated with (Ni,Pt)Al. Findings from this investigation includes: •
Parabolic oxide growth constants were determined, based on the assumption of diffusion-controlled kinetics, to be 0.241 ȝm/h1/2 and 0.175 ȝm/h1/2 for the TGO developed on the YSZ-coated and bond-coated-only sides, respectively. The TGO scale grew faster underneath the YSZ topcoat.
•
The β-to-γ’ and/or martensite transformation in the (Ni,Pt)Al bond coat occurred faster on the YSZcoated (Ni,Pt)Al bond coat, corresponding to the faster Al-depletion associated with the faster TGO growth observed on the YSZ-coated (Ni,Pt)Al bond coat.
•
Rumpling occurred faster and with larger amplitude on the (Ni,Pt)Al coating without the YSZ topcoat. Rumpling caused a significant decohesion at the YSZ/TGO interface with thermal cycling.
13
•
Crack initiation at the TGO/bond coat interface was observed as early as 10 thermal cycles. Despite the decohesion at the YSZ/TGO interface, and eventual spallation failure primarily through the TGO/bond coat interface.
•
The magnitude of compressive residual stress in the TGO showed an initial increase up to 3~4 GPa followed by a gradual decrease. The rate of stress relaxation was much quicker for the TGO scale without the YSZ topcoat with distinctive relief corresponding to the cracking at the top of geometrical ridges associated with the (Ni,Pt)Al bond coat.
•
The maximum elastic energy for the TGO scale was estimated at 90 J/m2 at 50% of its thermal cycling lifetime (Nf = 545 cycles).
Acknowledgment Authors would like to express sincerely appreciation of specimens provided by Mr. Ken Murphy of Howmet Research Corporation, Whitehall, MI, USA, for this study.
References [1] [2]
[3]
[4]
N.P. Padture, M. Gell, E.H. Jordan, “Thermal barrier coatings for gas-turbine engine applications,” Science, 296 (2002) 280-284. B.A. Pint, I.G. Wright, W.Y. Lee, Y. Zhang, K. Pruessner, K.B. Alexander, “Substrate and bond coat compositions: factors affecting alumina scale adhesion,” Materials Science and Engineering A., 245 (1998) 201-211. Y. Zhang, W.Y. Lee, J.A. Haynes, I.G. Wright, B.A. Pint, K.M. Cooley, P.K. Liaw, “Synthesis and cyclic oxidation behavior of a (Ni,Pt) Al coating on a desulfurized Ni-base superalloy,” Metallurgical and Materials Transactions A, 30 (1999) 2679-2687. Y. Zhang, J.A. Haynes, W.Y. Lee, I.G. Wright, B.A. Pint, K.M. Cooley, P.K. Liaw, “Effects of Pt incorporation on the isothermal oxidation behavior of chemical vapor deposition aluminide coatings,” Metallurgical and Materials Transactions A, 32 (2001) 1727-1741. 14
[5]
[6] [7] [8] [9] [10]
[11]
[12]
[13] [14] [15]
[16]
[17] [18] [19]
[20]
[21] [22]
B. Gleeson, N. Mu, S. Hayashi, “Compositional factors affecting the establishment and maintenance of Al2O3 scales on Ni–Al–Pt systems,” Journal of Materials Science, 44 (2009) 1704-1710. V.K. Tolpygo, D.R. Clarke, “Surface rumpling of a (Ni, Pt) Al bond coat induced by cyclic oxidation,” Acta Materialia, 48 (2000) 3283-3293. V.K. Tolpygo, D.R. Clarke, “On the rumpling mechanism in nickel-aluminide coatings: Part I: an experimental assessment,” Acta Materialia, 52 (2004) 5115-5127. M.Y. He, A.G. Evans, J.W. Hutchinson, “The ratcheting of compressed thermally grown thin films on ductile substrates,” Acta Materialia, 48 (2000) 2593-2601. H.E. Evans, “Oxidation failure of TBC systems: An assessment of mechanisms,” Surface and Coatings Technology, 206 (2011) 1512-1521. S. Laxman, B. Franke, B.W. Kempshall, Y.H. Sohn, L.A. Giannuzzi, K.S. Murphy, “Phase Transformation of Thermally Grown Oxide on (Ni,Pt)Al Bond Coat During Electron Beam Physical Vapor Deposition and Subsequent Oxidation,” Surface and Coatings Technology, 177178 (2004) 121-130. A.J. Burns, R. Subramanian, B.W. Kempshall, Y.H. Sohn, “Microstructure of As-Coated Thermal Barrier Coatings with Varying Lifetimes,” Surface and Coatings Technology, 177-178 (2004) 89-96. J. Liu, J.W. Byeon, Y.H. Sohn, “Effects of Phase Constituents and Microstructure in Thermally Grown Oxide on The Thermal Cycling Lifetime and Failure of Thermal Barrier Coatings,” Surface and Coatings Technology, 200 (2006) 5869-5876. N. Mu, J. Liu, Y.H. Sohn, Y.L. Nava, “Long-term Oxidation and Phase Transformations in Aluminized CMSX-4 Superalloys,” Surface and Coatings Technology, 188-189 (2005) 27-34. M.J. Stiger, N.M. Yanar, M.G. Topping, F.S. Pettit, G.H. Meier, “Thermal barrier coatings for the 21st century,” Zeitschrift fur Metallkunde. 1999. 90(12): p. 1069-1078. M.J. Stiger, N.M. Yanar, F.S. Pettit, G.H. Meier, “Mechanisms for the Failure of Electron Beam Physical Vapor Deposited Thermal Barrier Coatings Induced by High Temperature Oxidation,” Elevated Temperature Coatings: Science and Technology III, TMS, Warrendale, OH, 1999, 5165. V.K. Tolpygo, D.R. Clarke, K.S. Murphy, “The effect of grit blasting on the oxidation behavior of a platinum-modified nickel-aluminide coating,” Metallurgical and Materials Transactions A. 2001. 32(6): p. 1467-1478. Y. Zhang, J.A. Haynes, B.A. Pint, I.G. Wright, W.Y. Lee, “Martensitic transformation in CVD NiAl and (Ni, Pt) Al bond coatings,” Surface and Coatings Technology, 163 (2003) 19-24. V.K. Tolpygo, D.R. Clarke, “Microstructural study of the theta-alpha transformation in alumina scales formed on nickel-aluminides,” Materials at High Temperatures, 17 (2000) 59-70. B.W. Kempshall, Y.H. Sohn, S.K. Jha, S. Laxman, R.R. Vanfleet, J. Kimmel, “A microstructural observation of near-failure thermal barrier coating: a study by photostimulated luminescence spectroscopy and transmission electron microscopy,” Thin Solid Films, 466 (2004) 128-136. Y.H. Sohn, B. Jayaraj, S. Laxman, B. Franke, J.W. Byeon, A.M. Karlsson, “The nondestructive and nano-microstructural characterization of thermal-barrier coatings,” Journal of Metals, 56 (2004) 53-56. D.R. Mumm, A.G. Evans, I.T. Spitsberg, “Characterization of a cyclic displacement instability for a thermally grown oxide in a thermal barrier system,” Acta Materialia, 49 (2001) 2329-2340. A.G. Evans, D.R. Mumm, J.W. Hutchinson, G.H.Meier, F.S. Pettit “Mechanisms controlling 15
[23] [24]
[25]
[26] [27] [28] [29]
[30]
[31]
the durability of thermal barrier coatings,” Progress in Materials Science, 46 (2001) 505-553. X.Y. Gong, D.R. Clarke, “On the measurement of strain in coatings formed on a wrinkled elastic substrate,” Oxidation of Metals, 50 (1998) 355-376. Y.H. Sohn, K. Vaidyanathan, M. Ronski, E.H. Jordan, M. Gell, “Thermal cycling of EBPVD/MCrAlY thermal barrier coatings: II. Evolution of photo-stimulated luminescence,” Surface and Coatings Technology, 146 (2001) 102-109. M. Wen, E.H. Jordan, M. Gell, “Evolution of photo-stimulated luminescence of EBPVD/(Ni,Pt)Al thermal barrier coatings,” Materials Science and Engineering A, 398 (2005) 99107. J. Rösler, M. Bäker, K. Aufzug, “A parametric study of the stress state of thermal barrier coatings: part I: creep relaxation,” Acta Materialia, 52 (2004) 4809-4817. E. Schumann, C. Sarioglu, J.R. Blachere, F.S. Pettit, G.H. Meier, “High-temperature stress measurements during the oxidation of NiAl,” Oxidation of Metals, 53 (2000) 259-272. V.K. Tolpygo, D.R. Clarke, K.S. Murphy, “Oxidation-induced failure of EB-PVD thermal barrier coatings,” Surface and Coatings Technology, 146 (2001) 124-131. N.M. Yanar, F.S. Pettit, G.H. Meier, “Failure characteristics during cyclic oxidation of yttria stabilized zirconia thermal barrier coatings deposited via electron beam physical vapor deposition on platinum aluminide and on NiCoCrAlY bond coats with processing modifications for improved performances,” Metallurgical and Materials Transactions A, 37 (2006) 15631580. M.W. Chen, R.T. Ott, T.C. Hufnagel, P.K. Wright, K.J. Hemker, “Microstructural evolution of platinum modified nickel aluminide bond coat during thermal cycling,” Surface and Coatings Technology, 163 (2003) 25-30. M.W. Chen, M.L. Glynn, R.T. Ott, T.C. Hufnagel, K.J. Hemker “Characterization and modeling of a martensitic transformation in a platinum modified diffusion aluminide bond coat for thermal barrier coatings,” Acta Materialia, 51 (2003) 4279-4294. RMS =
1 n 2 ( yi − yo ) ¦ n i =1
[1]
Y = Yo + K p t
[2]
2
TGO energy ( J / m ) =
σ 2 (1 − υ ) h L (
E
16
Lo
)2 + γ M /O (
L 2 ) Lo
[3]
List of Figures Figure 1. (a) High angle annular dark field micrograph and (b) bright field TEM micrograph of the YSZ/bond coat interface including TGO for the as received specimen. Selected area electron diffraction patterns from the (c) α-Al2O3 TGO, (d) tetragonal YSZ and (e) B2 (Ni,Pt)Al bond coat. Figure 2. Thickness of the TGO scale for both the YSZ-coated and bond-coated-only sides as a function of square root of time for oxidation at 1100°C. Figure 3. Optical micrographs highlighting the phase transformations within the (Ni,Pt)Al bond coat as a function of thermal cycling at 1100°C: (a) as-received, (b) 10 cycles, (c) 50 cycles, and (d) 200 cycles underneath the YSZ topcoat; (e) as-received, (f) 10 cycles, (g) 50 cycles, and (h) 200 cycles on the bond-coated-only side. Figure 4. High angle annular dark field micrograph of the interface for specimens with (a) 10 and (b) 200 thermal cycles at 1100°C. Bright field micrograph of (c) the TGO and (d) martensite phase for specimen with 10 thermal cycles. Both (e) Ȗ’-Ni3Al and (f) martensite phases were observed in the bond coat after 10 thermal cycles at 1100°C. Figure 5. High magnification secondary electron micrographs highlighting the TGO scale development on the bond-coated-only surface: (a) in the as-received specimen, and specimens after thermal cycles of (b) 10, (c) 50, (d) 100, (e) 200 and (f) 400 at 1100°C. Figure 6. Cross-sectional backscatter electron micrographs from the YSZ-coated TBC specimens after thermal cycles of (a) 10, (b) 50, (c) 100, (d) 200, (e) 300 and (f) 400 at 1100°C.
17
Figure 7. Secondary electron micrographs of the (a) exposed bare bond coat surface and (b) bottom of the spalled YSZ after 545 thermal cycling at 1100°C. Cross-sectional backscatter electron micrographs of the (c) bond coat and (d) spalled YSZ. Figure 8. Evolution of the TGO/bond coat interfacial contours for the (a) bond-coated-only and (b) the YSZ-coated sides as a function of thermal cycling. Figure 9. Root mean square (RMS) roughness and tortuosity (L/L0) of the TGO/bond coat interface as a function of thermal cycles. Figure 10. Average compressive residual stress in the TGO scale as a function of thermal cycles measured by photostimulated luminescence spectroscopy. Inset shows the first few thermal cycles. Figure 11. TGO energy per unit area determined for the YSZ-coated TBCs as a function of thermal cycling.
18
Figure 1
Figure 2
Figure 3
Figure 4
Figure 5
Figure 6
Figure 7
Figure 8
Figure 9
Figure 10
Figure 11
Failure Characteristics and Mechanisms of EBPVD TBCs with Pt-Modified NiAl Bond Coats Le Zhou, Sriparna Mukherjee, Ke Huang, Young Whan Park, Yongho Sohn
www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(15)00380-9 http://dx.doi.org/10.1016/j.msea.2015.03.120 MSA32217
To appear in:
Materials Science & Engineering A
Received date: 9 September 2014 Revised date: 4 March 2015 Accepted date: 27 March 2015 Cite this article as: Le Zhou, Sriparna Mukherjee, Ke Huang, Young Whan Park, Yongho Sohn, Failure Characteristics and Mechanisms of EB-PVD TBCs with Pt-Modified NiAl Bond Coats, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2015.03.120 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Failure Characteristics and Mechanisms of EB-PVD TBCs with Pt-Modified NiAl Bond Coats Le Zhou, Sriparna Mukherjee# , Ke Huang&, Young Whan Park%, Yongho Sohn* Advanced Materials Processing and Analysis Center and Department of Materials Science and Engineering University of Central Florida, Orlando, FL, 32816, USA
#
Now with Department of Chemistry, North Carolina State University, Raleigh, NC 27607, USA &
%
Now with Siemens Energy, Inc., Orlando, FL 32826, USA
On Sabbatical from Department of Mechanical Engineering, Pukyung National University, Busan, Republic of Korea
* Corresponding author, Professor, Department of Materials Science and Engineering, University of Central Florida, Orlando, FL 32816, USA; Email: [email protected]; Voice: 407.882.1181
Key Words: Thermal barrier coatings, (Ni,Pt)Al bond coat, failure, microstructure, transmission electron microscopy
1
Abstract Microstructural evolution and failure characteristics/mechanisms were investigated for thermal barrier coatings that consist of electron beam physical vapor deposited ZrO2 - 8 wt.% Y2O3 (YSZ) topcoat, Pt-modified nickel aluminide, (Ni,Pt)Al bond coat, and CMSX-4 superalloy substrate with furnace cycling at 1100°C with 1-hour dwell. Photo stimulated luminescence spectroscopy, scanning electron microscopy equipped with X-ray energy dispersive spectroscopy and transmission electron microscopy were employed to examine the residual stress of the thermally grown oxide (TGO) and microstructural changes. For comparison, (Ni,Pt)Al bond coat on CMSX-4 without the YSZ topcoat was also characterized. The TGO grew faster for the YSZ-coated (Ni,Pt)Al bond coat than the (Ni,Pt)Al coating without the YSZ topcoat. Correspondingly, the β-to-γ’/martensite formation in the (Ni,Pt)Al bond coat occurred faster on the YSZ-coated (Ni,Pt)Al bond coat. However the rumpling occurred much faster and with larger amplitude on the (Ni,Pt)Al coating without the YSZ topcoat. Still, the rumpling at the TGO/bond coat interface caused crack initiation as early as 10 thermal cycles, decohesion at the YSZ/TGO interface, and eventual spallation failure primarily through the TGO/bond coat interface. The magnitude of compressive residual stress in the TGO showed an initial increase up to 3~4 GPa followed by a gradual decrease. The rate of stress relaxation was much quicker for the TGO scale without the YSZ topcoat with distinctive relief corresponding to the cracking at the top of geometrical ridges associated with the (Ni,Pt)Al bond coat. The maximum elastic energy for the TGO scale was estimated at 90 J/m2 at 50% of its lifetime (Nf = 545 cycles). The YSZ presence/adhesion to the TGO scale is emphasized to minimize the undulation of the TGO/bond coat interface, i.e., decohesion at the YSZ/TGO scale accelerates the rumpling and crack-coalescence at the TGO/bond coat interface where the spallation fracture occurs.
2
1. Introduction Thermal barrier coatings (TBCs) are multilayered coatings consisting of a ceramic topcoat, a metallic bond coat and a thermally grown oxide (TGO) that protected the underlying hot component from high temperature. The ceramic topcoat typically consists of 7-8 wt.% Y2O3 stabilized ZrO2 (YSZ) that is processed via electron beam physical deposition (EB-PVD) or air plasma spray (APS). The metallic bond coat is usually diffusion aluminide or MCrAlY (M=Ni, Cr or both) overlay coatings. The TGO layer forms between the ceramic top coat and metallic bond coat during high temperature operation due to oxidation [1]. Among various TBCs, Pt-modified NiAl aluminide coatings topped with EB-PVD YSZ ceramic topcoats are commonly used because of excellent spallation resistance. The incorporation of Pt into NiAl aluminide is believed to improve the oxidation behavior by strengthening alumina scale adhesion, slowing diffusion of detrimental element and retarding the void growth at the interface [2-5]. Understanding the mechanism that governs the failure of TBCs helps in the material selection and design of TBCs. TBC failure mechanisms depend on the type of TBCs and the operational conditions. For the EB-PVD YSZ coated on (Ni,Pt)Al bond coat under thermal cycling conditions, surface rumpling [6,7] or ratcheting [8] of the bond coat can occur, and cause instabilities of the interface between the TGO and bond coat. The rough interface may give rise to out-of-plane tensile stress that leads to the spallation or delamination of the TBCs. Furthermore the strain energy of the TGO accumulates during cyclic oxidation, and failure can occur when the strain energy within the TGO exceeds a certain critical value [9]. In this paper, commercial-production TBCs consisting of EB-PVD YSZ topcoat, (Ni,Pt)Al bond coat, and CMSX-4 superalloy substrate were furnace cycled at 1100°C without temperature gradient within the specimens. The TBC specimens were fabricated so that one side was coated only with bond coat, and the other side was coated with YSZ topcoat with (Ni,Pt)Al 3
bond coat. Thermal cycling lifetime of TBCs was initially assessed. Then, as a function of thermal cycling (i.e., fraction of lifetime), compressive residual stress, phase constituents, and microstructure were examined by photo-stimulated luminescence (PSLS), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Characteristics of spallation failure are documented and discussed with respect to the failure mechanisms.
2. Experimental details All EB-PVD TBC examined were disk-shaped specimens with diameter of 25.4 mm and thickness of approximately 3 mm. One face-side of the disk was coated with 175 ȝm-thick EB-PVD 8 wt.% Y2O3 stabilized ZrO2 (8YSZ) ceramic topcoat on top of a 45 ȝm-thick grit-blasted (Ni,Pt)Al bond coat. On the other face-side of the disk, only the 45 ȝm-thick (Ni,Pt)Al bond coat was coated. The superalloy substrate for all samples was CMSX-4 with the nominal composition, in weight %, of Ni9.0Co-6.5Cr-6.5Ta-6.0W-5.6Al-3.0Re-1.0Ti-0.6Mo-0.1Hf. Thermal cycling tests were carried out using CMTM Rapid High Temperature furnace. Each cycle consisted of a 10-min heat up to 1100°C, a 1-hour dwell at 1100°C, followed by a 10-min forced air quenching to room temperature. Thermal cycling tests were performed until the failure of TBC defined as the spallation of YSZ (more than 10 area %) from the substrate (Nf = 545 cycles). All TBC specimens were periodically removed from the furnace for the measurement of average compressive stress by photo-stimulated luminescence spectroscopy (PSLS) using a RenishawTM System 2000 Ramanscope. Measurements were performed from the topcoat for the YSZ-coated side, and from the ridges and valleys for the bond-coat-only side. Twenty random measurements were carried out to obtain the average value along with the standard deviation. 4
Several TBC specimens, prior to spallation-failure, were also removed from the furnace after thermal cycles of 10, 50, 100, 200, 300, 400, and 500 for microstructural analysis. Morphological evolution of the bond-coat-only surface was examined by ZeissTM Ultra 55 Field-Emission SEM equipped with X-ray energy dispersive spectroscopy (XEDS). Each specimen was first mounted in epoxy, sectioned by low speed diamond saw, metallographically ground, and polished down to 0.25ȝm followed by chemical etching using HCl/HNO3 (5:1) etchant. Cross-section of each specimen was examined by optical microscopy (OlympusTM LEXT OLS 3000 CSM) and SEM with XEDS. Transmission electron microscopy (TEM FEITM Tecnai F30 300 keV) equipped with high angle annular dark field (HAADF) detector was employed to examine the details of the YSZ/TGO and TGO/bond coat interfaces for the as-coated specimen and specimens subjected to 10 and 200 thermal cycles. The TEM foils were prepared via site-specific focus ion beam (FIB FEITM 200) in-situ lift-out (INLO) technique. Bond coat rumpling was quantitatively measured through a series of cross-sectional SEM micrographs. The interface between the TGO and the bond coat was traced and quantified. The tortuosity (L/L0) that is defined as the ratio of undulated TGO/bond coat interfacial length to the initial length (i.e., straight) was determined. By considering the oxide particles between the bond coat and the superalloy substrate as inert markers, the thickness of the bond coat was determined with root mean square (RMS) value as given by:
RMS =
1 n 2 yi − yo ) ( ¦ n i=1
[1]
where n is the number of points, yi is the length between the YSZ topcoat and the particles, y0 is the mean height of the bond coat.
5
3. Microstructural analysis 3.1 As-received thermal barrier coatings Figures 1(a) and 1(b) present, respectively, the HAADF and bright field (BF) TEM micrographs of the as-coated EB-PVD TBCs, where a uniform TGO layer of ~350 nm was observed between the YSZ and grit-blasted ȕ-(Ni,Pt)Al bond coat. The selected area diffraction pattern (SADP) from one of the alumina grains is shown in Figure 1(c), and indicates that the TGO layer primarily consists of the ĮAl2O3. Metastable Al2O3 frequently found and reported [10-13] was not observed to develop on this grit-blasted ȕ-(Ni,Pt)Al bond coat. The crystal structures of the YSZ and the ȕ-(Ni,Pt)Al were also confirmed by SADPs to be tetragonal and B2, respectively, as represented in Figures 1(d) and 1(e). Both HAADF and BF TEM micrographs show that the YSZ/TGO and TGO/bond coat interface does not have any visible decohesion and/or voids. 3.2 Thermal growth of the oxide scale with thermal cycling The average TGO thickness, along with the standard deviations, for the thermally cycled specimens was measured from a series of cross-sectional backscatter electron (BSE) micrographs and plotted as a function of square root of time as shown in Figure 2. The linear relationship in Figure 2 demonstrates that the TGO growth follows the diffusion-controlled parabolic growth. Accordingly the parabolic growth constant Kp was calculated according to the equation:
Y = Yo + K p t
[2]
where Yo and Y are initial and instantaneous thickness of the TGO, and t is the oxidation time. The calculated growth constants for the YSZ-coated and bond-coat-only sides were 0.241 ȝm/h1/2 and 0.175 ȝm/h1/2, respectively. The TGO underneath the YSZ topcoat grew at a faster rate than that exposed to 6
air. This would be a result of higher partial pressure of oxygen with the YSZ at high temperature than that in the ambient atmosphere (e.g., consider oxygen content in YSZ and air). Furthermore, The Al2O3 breaks into the YSZ and forms small equiaxed oxide grains, as will be seen in Figure 4(b). This formation of intermixed oxide zone was pointed out to promote the outward diffusion of Al [14][15], probably due to more grain boundary paths for Al diffusion, which might increase the oxidation rate. Additionally, it is noted that the bond coat of the YSZ-coated side is grit-blasted while the bondcoated-only side is not. Previous research has shown that the impurities introduced by grit-blasting increases the diffusivity of Al and thus the oxidation rate [15].
3.3 Phase transformations in the ȕ-(Ni,Pt)Al bond coat Figure 3 presents cross-sectional optical micrographs that illustrate the phase evolution within the (Pt,Ni)Al bond coat. The grain boundaries are revealed by chemical etching. For the specimen subjected to 50 cycles in Figures 3(c) and 3(g), white-contrast phase start to appear in the bond coat, preferably at the grain boundaries. These were Ȗ’-Ni3Al phases as will be seen later. With oxidation, the Ȗ’-Ni3Al phase grows as shown in Figures 3(d) and 3(h) due to the Al depletion via TGO formation and substrate-bond coat interdiffusion. Figures 3(d) and 3(h) shows that the growth of Ȗ’ phase appears to be faster for the bond coats with YSZ topcoat. Martensite formation in the bond coat is revealed by the lamellar microstructure with etching. Although the bond coat initially consists of the ȕ phase, the formation of martensite can occur when Al content falls below 37 at.% [17]. The martensite lamellae are observed only after 10 thermal cycles on the side coated with the YSZ as shown in Figure 3(b), and after 50 thermal cycles for the side coated only with (Ni,Pt)Al bond coat. In this study, the transformation to both Ȗ’ and martensite was observed earlier for the YSZ-coated side. This is consistent with faster Al consumption from the faster TGO growth observed on the YSZ-coated side, 7
and conforms that, indeed, the outward diffusion of Al for the formation of Al2O3 is faster when YSZ topcoat is present. Two TEM specimens were prepared from the interfacial area between the TGO and the bond coat for specimens with 10 and 200 thermal cycles. The SADP shown in Figure 4(e) confirms the presence of Ȗ’-Ni3Al, which formed due to the depletion of Al in the bond coat, which initially consisted of the ȕ-NiAl solution phase. Figures 4(d) and 4(f) show the microstructure of the martensites, and the corresponding SADP from the boundary area of the martensitic variants. These variants form twinned microstructure with each variant having a thickness of a few hundred nanometers.
3.4 Rumpling and/or racheting in the ȕ-(Ni,Pt)Al bond coat Figure 5 presents morphological evolution on the top surface of the bond coat without EB-PVD YSZ. Initially, the TGO forms a needle- or plate-like morphology, which is frequently associated with the formation of metastable alumina at the beginning thermal cycles [18]. With high temperature oxidation, the TGO transforms to the stable Į-Al2O3 with a network-like morphology, however with some cracks on top of the ridge area only after 50 thermal cycles, as shown in Figure 5(c). As the thermal cycling continues, the ridges become more prominent and the cracks propagate along the ridges as presented in Figures 5(d) and 5(e). Further thermal cycling leads to connection of these cracks along the ridges and some crack nucleation within the flat/troughs region of the TGO as presented in Figure 5(f). Cross-sectional BSE micrographs of the TGO scale on the YSZ-coated side are presented in Figure 6. The separation in Figures 6(e) and 6(f) is due to metallographic preparation. Initially, a uniform TGO layer forms and attaches to the bond coat tightly. The TGO rumples following the 8
contour-evolution of the bond coat. With increasing thermal cycles, some decohesions at the TGO/YSZ and TGO/bond coat, particularly around the grooved regions of the bond coat, are observed as shown in Figure 6(d). Although the asperity of the bond coat remains relatively fixed, many grooves (e.g., inward oxidation) form after prolonged thermal cycles, as shown in Figures 6(e) and 6(f). Figure 7 present the SEM micrographs from the TBC sample failed by 545 thermal cycling. Figure 7(a) shows the exposed bond coat surface, and Figure 7(b) shows the bottom of the spalled YSZ, Figures 7(c) and 7(d) are the cross-sectional BSE micrographs of the YSZ-spalled specimen and spalled YSZ, respectively. Much of the TGO scale remained attached to the spalled YSZ, and that observed on the bare bond coat was mostly located in the groove area as represented in the inset in Figure 7(a). Circled region in Figure 7(c) also demonstrate the residual TGO attached to the bond coat exists in the groove areas. Although decohesion between the YSZ and TGO is quite clear, much of the TGO scale was observed to be attached to the spalled YSZ. Spallation failure with fracture at the TGO/bond coat interface despite a significant decohesion damage at the YSZ/TGO interface has been observed previously [19]. This failure mechanism has been qualitatively described based on degraded ability of the YSZ to suppress the buckling of the TGO scale due to the decohesion damage at the YSZ/TGO interface [20]. A direct view of the rumpling/racheting behavior of the bond coat as a function of thermal cycling is presented by the contour of the TGO/bond coat interface for the bond-coated-only side and the YSZ-coated side as plotted in Figures 8(a) and 8(b), respectively. For each specimen, a distance of more than 2500 ȝm was examined. The RMS and L/L0 for the bond-coated-only and YSZ-coated side were also calculated and plotted in Figure 9 as a function of thermal cycles. The obvious gain in the amplitude of the contour indicates that the bond coat surface rumpled accumulatively with increasing thermal cycles. The amplitude increases faster for the side with bond coat only. Moreover, the bond9
coated-only surface develops both upward and downward contour change, while only grooving is generally observed for the side with YSZ topcoat. This is consistent with bond coat instability for the YSZ-coated side since it is appropriately termed ratcheting that does not require the upward displacement of the bond coat as in the rumpling behavior observed for the bond-coated-only side [8,9,21,22]. The presence of YSZ topcoat would place a rigid restriction on the instability of the bond coat surface, and can cause a preferred formation of grooves only. A small decrease in L/L0 for the YSZ-coated side was observed at the beginning cycles as shown in Figure 9. Although not substantiated, it could be related to the partial recovery of the severe plastic strains induced by grit blasting. The rumpling and/or ratcheting at the TGO/bond coat interface was observed as a significant damage-factor causing the spallation failure, since the growth of the TGO scale on a wrinkled surface can develop out-of-plane stress. Based on curvature, the stress is tensile at the ridge region and compressive at the groove region, which causes separation between TGO/bond coat interface [23]. Figures 4(a) and 4(b) present the HAADF micrographs of the interfaces of the specimens after 10 and 200 thermal cycles, respectively. Figure 4(c) is the BF TEM micrograph of the interface for the specimen with 10 thermal cycles, highlighting the nucleation of cracks at the TGO/bond coat interface. It is recalled from Figure 1 that no cracks were observed in the as-coated specimen. These cracks would coalesce and grow with further thermal cycling, as observed in Figure 4(b) for the specimen with 200 thermal cycles. The separation between the TGO and the bond coat eventually leads to the failure of the TBC. The damage at the TGO/bond coat interface is the primary reason for final failure, while the degradation via decohesion at the TGO/YSZ interface reduces the suppression ability of the YSZ against buckling of the TGO. In addition, impurities introduced by grit blasting may degrade the fracture resistance of the TGO/bond coat interface [16]. 10
4. Compressive residual stress in the thermally grown oxide The average compressive residual stress in the TGO scale was measured by PSLS as a function of thermal cycles as presented in Figure 10. The open circles represent the magnitude of average compressive residual stress in the TGO underneath the YSZ topcoat. The squares and cross represent TGO stress of the flat area and the ridge regions, respectively, for the bond-coated-only side. In general, all three average compressive residual stress increase slightly initially, followed by continues decrease with thermal cycling. A peak magnitude of approximately 3 GPa is observed for all curves, which corresponds to that determined by thermal expansion mismatch. The initial increase in TGO stress can be caused by the TGO growth or the transformation from metastable Al2O3 to stable Į-Al2O3 [24,25]. A gradual release of the TGO stress can be attributed to the surface rumpling or ratcheting of the bond coat [6,8] or creep of the bond coat [26]. The magnitude of average compressive residual stress in the TGO underneath the YSZ topcoat remains slightly higher than that of the TGO developed on the overlay (Ni,Pt)Al bond coat. Faster TGO growth (and thus the higher growth stress [27]), and lower rate of undulation for the TGO/bond coat interface potentially due to the suppression by the YSZ topcoat, may help retain the higher magnitude of compressive residual stress within the TGO underneath the YSZ topcoat. On the other hand, the magnitude of compressive residual stress within the TGO on top of the ridges observed on the bond-coat-only side decreased rapidly after 50 cycles, consistent with formation of cracks along the ridges as seen in Figure 5.
11
5. Discussion Experimental findings can be correlated quantitatively and qualitatively to provide some insight into the failure mechanisms of TBCs examined in this study. The elastic energy accumulated in the TGO scale on cooling is considered as the main driving force for the final TBC failure [28,29]. In a simplified model, the TGO elastic energy per unit area can be defined [28,29] to contain two parts: strain energy within the TGO and TGO/bond coat interfacial energy, as expressed by:
TGO energy(J / m 2 ) =
σ 2 (1− υ ) h L E
(
Lo
)2 + γ M /O (
L 2 ) Lo
[3]
where ı is the magnitude of compressive residual stress in the TGO, Ȟ is the Poisson’s ratio (~0.25), h is the TGO thickness, E is the TGO elastic modulus (~370 GPa), L/Lo is the tortuosity and ȖM/O is the interfacial energy at the TGO/bond coat interface (~7 J/m2) [22]. By substituting the measured values for ı, h and L/L0 from this study, the TGO energy was calculated as a function of thermal cycles as plotted in Figure 11. The first term on the right-hand-side of the Eq. [4], strain energy within the TGO, is much larger in magnitude than the second term in Eq. [3], the TGO/bond coat interfacial energy. With an increase in thermal cycles, the TGO energy continues to increase to about 90 J/m2. After 400 thermal cycles, a sudden decrease in the TGO energy is observed, corresponding to a partial, yet a significant damage in the TGO scale or at the TGO/bond coat interface. With failure, the decrease in the TGO energy would correspond to the relief of the energy stored in the scale. Previous studies indicate that phase transformations in the bond coat can contribute to the rumpling of the bond coat. The ȕ-to-Ȗ’ phase transformation is associated with a significant molar volume increase of approximately 8% to 38% [7], and the martensitic transformation can induce a transformation strain of 0.7% due to a 2% molar volume difference between the martensite and the ȕ 12
phase [30,31]. In this study, the transformation to Ȗ’ and/or martensite was observed earlier for the YSZ-coated side along with the faster depletion of Al due to faster TGO growth. The undulation of the bond coat surface, however, occurred faster for the bond-coated-only side than the YSZ-coated side as seen in Figure 9. Therefore, the undulation of the bond coat surface is strongly influenced by the adherence of the YSZ coating to the TGO scale, and the formation of Ȗ’ and martensites does not appear to be the only factor on the overall rumpling behavior of the bond coat underneath the YSZ topcoat. 5. Summary Microstructural evolution and failure characteristics of EB-PVD YSZ TBCs with Pt-modified NiAl bond coat were examined as a function of furnace thermal cycling at 1100°C with dwell time of 1 hour using PSLS, XRD, SEC/XEDS, and TEM. To examine the influence of EB-PVD topcoat during thermal cycling, the other side of the button specimen was only coated with (Ni,Pt)Al. Findings from this investigation includes: •
Parabolic oxide growth constants were determined, based on the assumption of diffusion-controlled kinetics, to be 0.241 ȝm/h1/2 and 0.175 ȝm/h1/2 for the TGO developed on the YSZ-coated and bond-coated-only sides, respectively. The TGO scale grew faster underneath the YSZ topcoat.
•
The β-to-γ’ and/or martensite transformation in the (Ni,Pt)Al bond coat occurred faster on the YSZcoated (Ni,Pt)Al bond coat, corresponding to the faster Al-depletion associated with the faster TGO growth observed on the YSZ-coated (Ni,Pt)Al bond coat.
•
Rumpling occurred faster and with larger amplitude on the (Ni,Pt)Al coating without the YSZ topcoat. Rumpling caused a significant decohesion at the YSZ/TGO interface with thermal cycling.
13
•
Crack initiation at the TGO/bond coat interface was observed as early as 10 thermal cycles. Despite the decohesion at the YSZ/TGO interface, and eventual spallation failure primarily through the TGO/bond coat interface.
•
The magnitude of compressive residual stress in the TGO showed an initial increase up to 3~4 GPa followed by a gradual decrease. The rate of stress relaxation was much quicker for the TGO scale without the YSZ topcoat with distinctive relief corresponding to the cracking at the top of geometrical ridges associated with the (Ni,Pt)Al bond coat.
•
The maximum elastic energy for the TGO scale was estimated at 90 J/m2 at 50% of its thermal cycling lifetime (Nf = 545 cycles).
Acknowledgment Authors would like to express sincerely appreciation of specimens provided by Mr. Ken Murphy of Howmet Research Corporation, Whitehall, MI, USA, for this study.
References [1] [2]
[3]
[4]
N.P. Padture, M. Gell, E.H. Jordan, “Thermal barrier coatings for gas-turbine engine applications,” Science, 296 (2002) 280-284. B.A. Pint, I.G. Wright, W.Y. Lee, Y. Zhang, K. Pruessner, K.B. Alexander, “Substrate and bond coat compositions: factors affecting alumina scale adhesion,” Materials Science and Engineering A., 245 (1998) 201-211. Y. Zhang, W.Y. Lee, J.A. Haynes, I.G. Wright, B.A. Pint, K.M. Cooley, P.K. Liaw, “Synthesis and cyclic oxidation behavior of a (Ni,Pt) Al coating on a desulfurized Ni-base superalloy,” Metallurgical and Materials Transactions A, 30 (1999) 2679-2687. Y. Zhang, J.A. Haynes, W.Y. Lee, I.G. Wright, B.A. Pint, K.M. Cooley, P.K. Liaw, “Effects of Pt incorporation on the isothermal oxidation behavior of chemical vapor deposition aluminide coatings,” Metallurgical and Materials Transactions A, 32 (2001) 1727-1741. 14
[5]
[6] [7] [8] [9] [10]
[11]
[12]
[13] [14] [15]
[16]
[17] [18] [19]
[20]
[21] [22]
B. Gleeson, N. Mu, S. Hayashi, “Compositional factors affecting the establishment and maintenance of Al2O3 scales on Ni–Al–Pt systems,” Journal of Materials Science, 44 (2009) 1704-1710. V.K. Tolpygo, D.R. Clarke, “Surface rumpling of a (Ni, Pt) Al bond coat induced by cyclic oxidation,” Acta Materialia, 48 (2000) 3283-3293. V.K. Tolpygo, D.R. Clarke, “On the rumpling mechanism in nickel-aluminide coatings: Part I: an experimental assessment,” Acta Materialia, 52 (2004) 5115-5127. M.Y. He, A.G. Evans, J.W. Hutchinson, “The ratcheting of compressed thermally grown thin films on ductile substrates,” Acta Materialia, 48 (2000) 2593-2601. H.E. Evans, “Oxidation failure of TBC systems: An assessment of mechanisms,” Surface and Coatings Technology, 206 (2011) 1512-1521. S. Laxman, B. Franke, B.W. Kempshall, Y.H. Sohn, L.A. Giannuzzi, K.S. Murphy, “Phase Transformation of Thermally Grown Oxide on (Ni,Pt)Al Bond Coat During Electron Beam Physical Vapor Deposition and Subsequent Oxidation,” Surface and Coatings Technology, 177178 (2004) 121-130. A.J. Burns, R. Subramanian, B.W. Kempshall, Y.H. Sohn, “Microstructure of As-Coated Thermal Barrier Coatings with Varying Lifetimes,” Surface and Coatings Technology, 177-178 (2004) 89-96. J. Liu, J.W. Byeon, Y.H. Sohn, “Effects of Phase Constituents and Microstructure in Thermally Grown Oxide on The Thermal Cycling Lifetime and Failure of Thermal Barrier Coatings,” Surface and Coatings Technology, 200 (2006) 5869-5876. N. Mu, J. Liu, Y.H. Sohn, Y.L. Nava, “Long-term Oxidation and Phase Transformations in Aluminized CMSX-4 Superalloys,” Surface and Coatings Technology, 188-189 (2005) 27-34. M.J. Stiger, N.M. Yanar, M.G. Topping, F.S. Pettit, G.H. Meier, “Thermal barrier coatings for the 21st century,” Zeitschrift fur Metallkunde. 1999. 90(12): p. 1069-1078. M.J. Stiger, N.M. Yanar, F.S. Pettit, G.H. Meier, “Mechanisms for the Failure of Electron Beam Physical Vapor Deposited Thermal Barrier Coatings Induced by High Temperature Oxidation,” Elevated Temperature Coatings: Science and Technology III, TMS, Warrendale, OH, 1999, 5165. V.K. Tolpygo, D.R. Clarke, K.S. Murphy, “The effect of grit blasting on the oxidation behavior of a platinum-modified nickel-aluminide coating,” Metallurgical and Materials Transactions A. 2001. 32(6): p. 1467-1478. Y. Zhang, J.A. Haynes, B.A. Pint, I.G. Wright, W.Y. Lee, “Martensitic transformation in CVD NiAl and (Ni, Pt) Al bond coatings,” Surface and Coatings Technology, 163 (2003) 19-24. V.K. Tolpygo, D.R. Clarke, “Microstructural study of the theta-alpha transformation in alumina scales formed on nickel-aluminides,” Materials at High Temperatures, 17 (2000) 59-70. B.W. Kempshall, Y.H. Sohn, S.K. Jha, S. Laxman, R.R. Vanfleet, J. Kimmel, “A microstructural observation of near-failure thermal barrier coating: a study by photostimulated luminescence spectroscopy and transmission electron microscopy,” Thin Solid Films, 466 (2004) 128-136. Y.H. Sohn, B. Jayaraj, S. Laxman, B. Franke, J.W. Byeon, A.M. Karlsson, “The nondestructive and nano-microstructural characterization of thermal-barrier coatings,” Journal of Metals, 56 (2004) 53-56. D.R. Mumm, A.G. Evans, I.T. Spitsberg, “Characterization of a cyclic displacement instability for a thermally grown oxide in a thermal barrier system,” Acta Materialia, 49 (2001) 2329-2340. A.G. Evans, D.R. Mumm, J.W. Hutchinson, G.H.Meier, F.S. Pettit “Mechanisms controlling 15
[23] [24]
[25]
[26] [27] [28] [29]
[30]
[31]
the durability of thermal barrier coatings,” Progress in Materials Science, 46 (2001) 505-553. X.Y. Gong, D.R. Clarke, “On the measurement of strain in coatings formed on a wrinkled elastic substrate,” Oxidation of Metals, 50 (1998) 355-376. Y.H. Sohn, K. Vaidyanathan, M. Ronski, E.H. Jordan, M. Gell, “Thermal cycling of EBPVD/MCrAlY thermal barrier coatings: II. Evolution of photo-stimulated luminescence,” Surface and Coatings Technology, 146 (2001) 102-109. M. Wen, E.H. Jordan, M. Gell, “Evolution of photo-stimulated luminescence of EBPVD/(Ni,Pt)Al thermal barrier coatings,” Materials Science and Engineering A, 398 (2005) 99107. J. Rösler, M. Bäker, K. Aufzug, “A parametric study of the stress state of thermal barrier coatings: part I: creep relaxation,” Acta Materialia, 52 (2004) 4809-4817. E. Schumann, C. Sarioglu, J.R. Blachere, F.S. Pettit, G.H. Meier, “High-temperature stress measurements during the oxidation of NiAl,” Oxidation of Metals, 53 (2000) 259-272. V.K. Tolpygo, D.R. Clarke, K.S. Murphy, “Oxidation-induced failure of EB-PVD thermal barrier coatings,” Surface and Coatings Technology, 146 (2001) 124-131. N.M. Yanar, F.S. Pettit, G.H. Meier, “Failure characteristics during cyclic oxidation of yttria stabilized zirconia thermal barrier coatings deposited via electron beam physical vapor deposition on platinum aluminide and on NiCoCrAlY bond coats with processing modifications for improved performances,” Metallurgical and Materials Transactions A, 37 (2006) 15631580. M.W. Chen, R.T. Ott, T.C. Hufnagel, P.K. Wright, K.J. Hemker, “Microstructural evolution of platinum modified nickel aluminide bond coat during thermal cycling,” Surface and Coatings Technology, 163 (2003) 25-30. M.W. Chen, M.L. Glynn, R.T. Ott, T.C. Hufnagel, K.J. Hemker “Characterization and modeling of a martensitic transformation in a platinum modified diffusion aluminide bond coat for thermal barrier coatings,” Acta Materialia, 51 (2003) 4279-4294. RMS =
1 n 2 ( yi − yo ) ¦ n i =1
[1]
Y = Yo + K p t
[2]
2
TGO energy ( J / m ) =
σ 2 (1 − υ ) h L (
E
16
Lo
)2 + γ M /O (
L 2 ) Lo
[3]
List of Figures Figure 1. (a) High angle annular dark field micrograph and (b) bright field TEM micrograph of the YSZ/bond coat interface including TGO for the as received specimen. Selected area electron diffraction patterns from the (c) α-Al2O3 TGO, (d) tetragonal YSZ and (e) B2 (Ni,Pt)Al bond coat. Figure 2. Thickness of the TGO scale for both the YSZ-coated and bond-coated-only sides as a function of square root of time for oxidation at 1100°C. Figure 3. Optical micrographs highlighting the phase transformations within the (Ni,Pt)Al bond coat as a function of thermal cycling at 1100°C: (a) as-received, (b) 10 cycles, (c) 50 cycles, and (d) 200 cycles underneath the YSZ topcoat; (e) as-received, (f) 10 cycles, (g) 50 cycles, and (h) 200 cycles on the bond-coated-only side. Figure 4. High angle annular dark field micrograph of the interface for specimens with (a) 10 and (b) 200 thermal cycles at 1100°C. Bright field micrograph of (c) the TGO and (d) martensite phase for specimen with 10 thermal cycles. Both (e) Ȗ’-Ni3Al and (f) martensite phases were observed in the bond coat after 10 thermal cycles at 1100°C. Figure 5. High magnification secondary electron micrographs highlighting the TGO scale development on the bond-coated-only surface: (a) in the as-received specimen, and specimens after thermal cycles of (b) 10, (c) 50, (d) 100, (e) 200 and (f) 400 at 1100°C. Figure 6. Cross-sectional backscatter electron micrographs from the YSZ-coated TBC specimens after thermal cycles of (a) 10, (b) 50, (c) 100, (d) 200, (e) 300 and (f) 400 at 1100°C.
17
Figure 7. Secondary electron micrographs of the (a) exposed bare bond coat surface and (b) bottom of the spalled YSZ after 545 thermal cycling at 1100°C. Cross-sectional backscatter electron micrographs of the (c) bond coat and (d) spalled YSZ. Figure 8. Evolution of the TGO/bond coat interfacial contours for the (a) bond-coated-only and (b) the YSZ-coated sides as a function of thermal cycling. Figure 9. Root mean square (RMS) roughness and tortuosity (L/L0) of the TGO/bond coat interface as a function of thermal cycles. Figure 10. Average compressive residual stress in the TGO scale as a function of thermal cycles measured by photostimulated luminescence spectroscopy. Inset shows the first few thermal cycles. Figure 11. TGO energy per unit area determined for the YSZ-coated TBCs as a function of thermal cycling.
18
Figure 1
Figure 2
Figure 3
Figure 4
Figure 5
Figure 6
Figure 7
Figure 8
Figure 9
Figure 10
Figure 11