Thermal Cycling Behavior of Thermal Barrier Coatings with MCrAlY

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Thermal Cycling Behavior of Thermal Barrier Coatings with MCrAlY Bond Coat Irradiated by High-Current Pulsed Electron Beam Jie Cai,†,‡ Peng Lv,§ Qingfeng Guan,*,§ Xiaojing Xu,†,‡ Jinzhong Lu,*,‡ Zhiping Wang,∥ and Zhiyong Han∥ †

Institute of Advanced Manufacturing and Modern Equipment Technology, Jiangsu University, Zhenjiang, 212013, China School of Mechanical Engineering, Jiangsu University, Zhenjiang, 212013, China § School of Materials Science and Engineering, Jiangsu University, Zhenjiang, 212013, China ∥ College of Science, Civil Aviation University of China, Tianjin, 300300, China ‡

ABSTRACT: Microstructural modifications of a thermally sprayed MCrAlY bond coat subjected to high-current pulsed electron beam (HCPEB) and their relationships with thermal cycling behavior of thermal barrier coatings (TBCs) were investigated. Microstructural observations revealed that the rough surface of air plasma spraying (APS) samples was significantly remelted and replaced by many interconnected bulged nodules after HCPEB irradiation. Meanwhile, the parallel columnar grains with growth direction perpendicular to the coating surface were observed inside these bulged nodules. Substantial Y-rich Al2O3 bubbles and varieties of nanocrystallines were distributed evenly on the top of the modified layer. A physical model was proposed to describe the evaporation−condensation mechanism taking place at the irradiated surface for generating such surface morphologies. The results of thermal cycling test showed that HCPEB-TBCs presented higher thermal cycling resistance, the spalling area of which after 200 cycles accounted for only 1% of its total area, while it was about 34% for APS-TBCs. The resulting failure mode, i.e., in particular, a mixed delamination crack path, was shown and discussed. The irradiated effects including compact remelted surface, abundant nanoparticles, refined columnar grains, Y-rich alumina bubbles, and deformation structures contributed to the formation of a stable, continuous, slow-growing, and uniform thermally grown oxide with strong adherent ability. It appeared to be responsible for releasing stress and changing the cracking paths, and ultimately greatly improving the thermal cycling behavior of HCPEB-TBCs. KEYWORDS: thermal barrier coatings (TBCs), high-current pulsed electron beam (HCPEB), thermally grown oxide (TGO), microstructural modifications, thermal cycling behavior

1. INTRODUCTION

thermal expansion match between the TC and superalloy substrate and prevent the substrate from oxidation.6,7 In service, however, TBCs have a tendency to spall or debond of the ceramic topcoat, which originates from the formation and growth of microcracks at the TC/BC interface. The main reasons for this damage are oxidation of the BC layer8 as well as thermally- and mechanically induced stress in the TBC system9 during the thermal cycling process. Previous

Thermal barrier coatings (TBCs), serving as a heat-insulating layer on the surfaces of metallic components, have been extensively applied to combustion chambers and turbine blades of gas turbines.1−3 A typical TBC system consists of a thermally insulating ceramic top coat (TC) and an oxidation resistant metallic bond coat (BC). A 6−8% wt % yttria partially stabilized zirconia (YSZ) was commonly used as TC layer due to its high coefficient of thermal expansion and a significantly low thermal conductivity.4,5 MCrAlY (M = Ni, and/or Co) is an appropriate choice for the BC layer to provide a good © XXXX American Chemical Society

Received: September 3, 2016 Accepted: November 7, 2016 Published: November 7, 2016 A

DOI: 10.1021/acsami.6b11129 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Figure 1. Morphologies of powders: (a) CoCrAlY powders; (b) 8YSZ powders.

studies1,6−8,10 have demonstrated that the cracking is believed to be related to the thickness and nature of thermally grown oxide (TGO), typically α-Al2O3, which was formed at the interface between the TC and the BC due to the reaction of the penetrating oxygen and the metallic elements from the MCrAlY bond coat during high temperature operation. Rabiei and Evans8 found that the propagation of cracks in the YSZ coating moves toward the TC/TGO interface as TGO thickness increases. It is suggested11 that a higher TGO-growth rate should result in shorter TBCs-lifetime, which corresponds to the driving force (growth stress and thermal stress) accumulated in the TGO during thermal cycling due to the TGO growth and thermal expansion mismatch. Therefore, it is preferred to obtain a stable, continuous, slow-growing, and uniform TGO layer with strong adherent ability at the TC/BC interface. The TGO evolution is susceptible to the surface conditions of the MCrAlY coating, including the surface feature, microstructure, coating composition, and residual stress, which are dependent on the preparation process.7,12 The air plasma spray (APS) technique has been widely employed to deposit TBCs, especially the MCrAlY bond coat, because of the advantages associated with low cost, good practicality, and high efficiency.13,14 However, the as-sprayed coatings exhibit typical layered and porous structures, microcracks, a rough surface, and inhomogeneities, which dramatically affect the service life of TBCs, particularly on the growth of the TGO layer. Therefore, an effective method to conquer these thermally sprayed defects of the MCrAlY coating will be beneficial for the lifetime prediction of the TBC system. High-current pulsed electron beam (HCPEB) has been successfully developed as an efficient surface modification technique of different materials in recent years.15−20 During HCPEB irradiation, the high-density electron pulses of short duration can introduce different physical processes in the surface layer, such as rapid melting−solidification, evaporation−condensation as well as surface smoothing and annealing. These nonequilibrium processes can easily change the surface morphologies, microstructures, chemical compositions, phase structures, and stress states at the surface. Many of the previous studies have suggested that the substantial modifications obtained by HCPEB irradiation can effectively improve the mechanical properties of various metals, in particular, their fatigue, corrosion, and wear resistance,21−25 while the studies on HCPEB irradiation of coatings are comparatively few. Hao et al.26 investigated an arc-sprayed FeCrAl coating by HCPEB irradiation. They found that the high temperature corrosion resistance of the irradiated coating was improved effectively due to the glossy and compact remelted coating surface. In our studies,7,27 we found that HCPEB irradiation can make the

coating surface remelted, thermal defects annealed, grain crystals refined, and elemental distributions changed. These irradiated effects, especially the refined microstructures, can accelerate the preferential nucleation and instantaneous formation of α-Al2O3 at the very beginning stage of oxidation and, consequently, lead to a thinner, continuous, slow-growing, and compact TGO layer during the high temperature oxidation process. We have proved that the HCPEB technique is a promising method for enhancing the oxidation behavior of TBCs. It is generally known that thermal cycling with dwell time plays a vital role in controlling the durability, efficiency, and reliability of TBCs. However, to our knowledge, the microstructural modifications of the MCrAlY coating after HCPEB treatment and their relationships with thermal cycling behavior of TBCs are seldom paid full attention. The emphasis of the present work is put on the relationship between the microstructure evolution of the APS-MCrAlY coating and thermal cycling behavior of TBCs before and after HCPEB treatment. A detailed characterization of surface modification and failure mode associated with the crack evolution at the TC/BC interface under thermal cycling condition before and after HCPEB irradiation was analyzed in detail.

2. EXPERIMENTAL PROCEDURE 2.1. TBCs Preparation. GH4169 (Inconel 718), a nickel based superalloy, was used as the substrate (cut as ϕ 25.4 × 6 mm). Before deposition, a grit blasting machine (using corundum of 60-mesh) was employed on the substrate surface for achieving the required roughness and the mechanical interlocking between coating and substrate. The coatings were sprayed using an SG100 APS plasma torch, Praxair 3710 type, using Ar plasma. The coating powders were injected externally to the spraying gun, connected to a Six-axis Robot of type 2400 M (ABB), and the injection was optimized prior to each coating’s deposition to ensure process robustness and spraying parameters. The bond coat was deposited of approximately 160 μm in thickness, using commercially available CoCrAlY alloy powders with a nominal composition of Cr-23, Al-13, Y-0.5, and Co-balance (wt %). The CoCrAlY powers (PWA1348-2) exhibited a spherical morphology, and the particle sizes were about 5−45 μm, as shown in Figure 1a. Subsequently, the coated samples were irradiated by using a HCPEB technique of Nadezhda-2 type with 30 pulses. More details about the HCPEB irradiated conditions and irradiated processes have been reported in our previous studies.27,28 At last, a ceramic topcoat was deposited to a thickness of 240 μm also by the APS method, with powders of 8% Y2O3-ZrO2 (8YSZ). The hollow spherical 8YSZ powers (PWA 1375) had diameters ranging from 30 to 106 μm, as shown in Figure 1b. The corresponding spraying parameters employed in this work are summarized in Table 1. 2.2. Thermal Cycling Tests. Thermal cycling tests were carried out by using a muffle furnace. All the samples were subjected to thermal cycles in air, each cycle consisting of 20 min holding at 1050 B


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ACS Applied Materials & Interfaces Table 1. Plasma Spraying Parameters parameters

bond coat

top coat

voltage (V) current (A) powder feed rate (rpm) primary gas pressure (Ar, psi) secondary gas pressure (H2, psi) powder carrier pressure (Ar, psi) spray distance (mm) spray rate (mm/s)

38 750 2.5 60 110 40 85 450

39 860 3.5 60 30 30 72 250

°C and 10 min cooling to ambient temperature under the compressed air (pressure 0.8 MPa). After testing, the oxidized samples after 50, 100, 150, and 200 cycles were taken out to analyze in detail. More than 20% of the spalled region of the coating surface was adopted as a judging criterion for the failure of the TBC system.29,30 2.3. Characterization. In order to perform more detailed investigations of the morphology characteristics and thermal cycling behavior before and after HCPEB irradiation, four kinds of coating systems named A, B, C, and D were employed, as shown in Table 2.

Figure 2. XRD results of coatings A and B.

Table 2. Summary of Experimental Coating System coatings

preparations

single-layer

A B

BC BC + HCPEB

double-layer

C D

BC + TC BC + HCPEB + TC

surrounded by oxides were obtained. Protrusions could be observed scattering everywhere on the surface and caused a drastically undulating surface with a ravine feature or spiky margin, as seen in Figure 3a. Figure 3b,c shows many thermally spraying defects, like microcracks, large cavities, and unmelted and semimelted particles that were visible on the highly rough surface, which have proved to be detrimental for the growth of TGO and, ultimately, result in a spalling damage or even failure of TBCs as a potential defect origin. Figure 4 gives the surface micrograph of coating B performed by 30 HCPEB pulses. It is clear that the coating surface was modified ultimately and replaced by a series of hills (i.e., interconnected bulged nodules) and valleys with pores and large cavities sealing, as seen in Figure 4a. Seen visibly in the magnified LSM image in Figure 4b, substantial bubbles were attached to the bulged nodule surface. The magnified SEM image (Figure 4c) reveals that the top surface of these bulged nodules possessed a completely compact appearance, and a mass of small protrusions were deposited on it, which were consistent with the bubbles in Figure 4b. As seen in the magnified Figure 4d, black oxides were detected at the concave part between the surface bulged nodules, which were testified as Al2O3 by the EDS analysis in Figure 4i. Moreover, a detailed examination of Figure 4e reveals that refined grains were homogeneously distributed throughout the entire surface. At higher magnification (Figure 4f), these very fine grains with a size distribution of 50−800 nm can be predicatively characterized as distinguishable spherical grains. The same situation could be noticed from the AFM observation within the scanning range of 2 × 2 μm2, as shown in Figure 4g. These spherical grains possessed the average height of about 37 nm and an average size of about 285 nm, statistically analyzed by NanoScope Analysis Software. The possible compositions of these small protrusions (Figure 4c) and very refined grains (Figure 4f) were determined by EDS, which indicates that these small protrusions were mainly Y-rich Al2O3 (by further analysis combining the results of EDS and XRD, it could be confirmed that these Y-rich Al2O3 were YAlO3-Al2O3), while these refined grains were primarily an Al-rich Co matrix. It can be expected that the Co matrix was covered with a very thin Al2O3 film or even Y-rich Al2O3 bubbles were on the top of surface.

experimental analysis XRD analysis microstructural characterization thermal cycling test microstructural characterization

The microstructural characterizations of coatings A and B were comprehensively performed by using a scanning electron microscope (SEM) of type JSM-7100F on a microscope equipped with an energydispersive spectrometer (EDS), a three-dimensional laser scanning microscope (LSM) of type VK-X100/X200, an atomic force microscope (AFM) of type AFP-3D, and a transmission electron microscope (TEM) of type JEM-2100. The macroscopic characteristics of coatings C and D after thermal cycling were observed by a high megapixel camera. Thereafter, the oxidized coatings of C and D were sectioned by a low-speed diamond saw, mounted using epoxy, mechanically polished to 0.5 μm, and finally characterized by SEM.

3. RESULTS 3.1. XRD Analysis. Figure 2 gives the XRD results of coatings A and B. The initial coating A was mainly composed of three phases: γ-Co solid solution, β-CoAl intermetallic phase, and a minority σ-CoCr phase. Moreover, a small content of Cobased mixed oxides can be also detected. After HCPEB irradiation, the diffraction intensity of γ-Co was increased while that of β-CoAl was decreased. The peaks of the σ-CoCr phase and Co-based mixed oxides disappeared, and simultaneously, the Al2O3 and YAlO3 phases were clearly detectable. It was noted that the increment of the diffraction intensity, located at 2θ = 46.17°, was a consequence of the superposition of peaks of Al2O3 and YAlO3 phases. In addition, the peaks of coating B became broadened owing to the very fine substructures obtained in the top layer. Certainly, the emergence of chemical gradients that existed on the irradiated coating surface can also result in the peak broadening.31,32 3.2. Surface and Cross-Sectional Morphologies. The basic microstructures of coating A are provided in Figure 3, which shows that the coating possessed a typical APS characteristic structure. It is obvious that lamellar structures with many layered “composites” made out of alloy splats C


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Figure 3. Cross-sectional morphology (a), surface SEM image (b), and LSM image (c) of coating A.

obtained in the SAED pattern indicated the presence of the CoAl phase. (4) Cellular nanostructures. As shown in Figure 7d, these nanostructures had grain sizes of 20−70 nm. The picture depicts that the boundaries of the cellular grains had a darker color than that inside, which is suggested that the content of alloy elements in the cell boundaries was higher. Moreover, very fine precipitates of less than 10 nm were located at the grain boundaries and in particular most often at triple junctions. Similar results were previously reported in 3Cr13 steel33 and D2 steel17 irradiated by HCPEB irradiation. The existence of these granular precipitates can suppress the expansion of grain boundaries, which has a positive effect on the high temperature oxidation resistance of TBCs by a grain size constraining effect. Dense dislocation walls are visible under TEM observations, as seen in Figure 8a. A detailed examination in Figure 8b shows that the boundaries of dislocation walls were very clear and parallel to each other, with an occasional bifurcation case. The space between adjacent dislocation walls was estimated as 15− 40 nm. Interestingly, the existence of parallel sub-dislocation walls inside dislocation walls can be apparently found in Figure 8b. The space between adjacent sub-dislocation walls, which were approximately perpendicular to dislocation walls, can be estimated as 5−10 nm. The microstructure clearly shows that a pronounced plastic deformation occurred, and consequently, the effect of grain refinement can be expected by the intersection between dislocation walls and sub-dislocation walls. The diffraction rings obtained in the SAED pattern can be also regarded as a witness mark of grain refinement induced by HCPEB irradiation. 3.4. Temperature Field Simulation. The simulation of the temperature field was used to estimate thermal-dynamic changes of the CoCrAlY coating associated with 1-pulsed HCPEB treatment under an accelerating voltage of 27 keV and an energy density of 4 J/cm2, as shown in Figure 9. A procedural detail about the calculation for the present case was based on the previous studies.19,34 Figure 9a shows the evolution of the temperature profile vs sample depth and time. On heating, two plateaus were noticeable. One was the melting plateau at 1728 K corresponding to the melting point, where the solid−liquid biphase region was present. Another was the evaporation plateau at 3000 K corresponding to the

Figure 5 shows the cross-sectional SEM images of coating B. As can be seen in Figure 5a, the cross section of the irradiated coating was featured by two distinct layers: the “discontinuously” modified layer with an average depth of 35 μm (the upper layer) and the initial as-sprayed layer with lamellar structures (the lower layer). For further magnification (Figure 5b), these bulged nodules were completely dense and free from pores and oxides. After chemically corroding in the aqua regia (HCl:HNO3 = 3:1), the parallel columnar grains with a growth direction perpendicular to the coating surface were observed inside these bulged nodules, as shown in Figure 5c,d. The size of the columnar crystals was 1−2 μm, approximately. Besides, dendrites with the same orientations were distributed in the interior of columnar region. To confirm the element distributions on the cross section of coating B, a detailed map scanning analysis was employed, as shown in Figure 6. The results clearly reflect that depletions of Al, Y, and O elements and the relatively uniform distribution of Co and Cr elements were found within the remelted layer. Simultaneously, the aluminum concentration that occurred closer to the interface was much higher than that inside the bulged nodules. It is suggested that the partial Al element was released from the irradiated layer and then enriched particularly at the top surface. 3.3. TEM Analysis. The TEM observations were further carried out to investigate the microstructural modifications in the remelted layer, as shown in Figure 7. It is clear that, after HCPEB irradiation, varieties of nanocrystallines were formed and distributed evenly on the top of modified layer, which can be divided into four kinds: (1) Circular bubbles. As seen in Figure 7a, these bubbles exhibited a near-perfect circular morphology and their sizes were within 100 nm. The corresponding selected area electron diffraction (SAED) pattern (inset in Figure 7a) demonstrates that these bubbles can be indexed by the CoAl phase, which was consistent with the SEM result in Figure 4f. (2) Nanoscale particles. As displayed in Figure 7b, the average size of these particles was only a few nanometers. The corresponding SAED pattern exhibited the representative diffraction rings, which can be also indexed by CoAl phase. (3) Equiaxed nanograins. As shown in Figure 7c, these domains (∼20 nm) had a regular shape and clear grain boundary, where the similar diffraction rings D


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Figure 4. (a) SEM image showing a succession of hills and valleys of coating B; (b) substantial bubbles; (c) completely compact appearance; (d) black oxides at concave part; (e) very fine grains; (f) magnified SEM image showing nanograins of (e); (g) AFM image of (f); (h)−(j) EDS analysis detected form (c), (d), and (f), respectively.

calculation also revealed a very fast heating (108 K/s) and cooling (107 K/s) thermal cycle occurring in the near surface layer. 3.5. Thermal Shock Resistance. The cross-sectional SEM observations of coatings C and D are present in Figure 10. As seen in Figure 10a, the initial TBC possessed two typical coatings: the porous ceramic coating with many cracklike discontinuities, and the lamellar MCrAlY coating. The highly undulated nature of the TC/BC interface was featured by local convexities and concavities. By contrast, the cross-sectional morphology of the irradiated TBC displayed three distinct layers: the porous ceramic layer, the refined melted layer, and the initial lamellar bonding layer, as can be seen in Figure 10b.

boiling point, within which the material kept a constant temperature due to the evaporation. On cooling, it is clearly that the solidification plateau can be also observed, which lasted much longer than that of the melting one. The temperature evolving curve over a depth range of 0−10 μm and for a duration of 0−5 μs was emphasized in Figure 9b. In comparison with the curves, the total vaporizing duration at the very top surface with a maximum temperature of 3000 K was estimated as 1.2 μs, while it was about 0.6 μs at a depth of 1 μm. Accompanied by the heat constantly transferring to the matrix, the temperature beneath the surface increased first and decreased afterward. It is suggested that a very high temperature gradient along the depth occurred. Besides, the E


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Figure 5. (a) Cross-sectional morphology of coating B; (b) further magnification of (a); (c) corroded morphology of coating B; (d) further magnification of (c).

Figure 6. Map scanning analysis showing the element distributions within the irradiated layer: (a) SEM image, (b) Co, (c) Cr, (d) Al, (e) O, and (f) Y element.

apparent spot spallations were observed at the edge of top layer after 100 cycles. With the thermal cycling going on, the spalled area continuously enlarged. There were about 5% of the cracked and peeled-off regions at the edge of the coating surface after 150 cycles. As the number of cycles increased up to

Notably, the ceramic layer and the remelted layer had a satisfactory adhesion, which showed a regular and wavelike interface. Surface morphologies of coating C after thermal cycling test are present in Figure 11. As seen in the macro photographs, F


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Figure 7. TEM observations of coating B: (a) circular bubbles; (b) nanoscale particles; (c) equiaxed nanograins; and (d) cellular nanostructures.

Figure 8. (a) TEM image showing dislocation walls of coating B; (b) magnified TEM image of (a).

Figure 9. Calculated temperature evolution (a) and corresponding temperature evolving curve (b) in coating B irradiated by HCPEB with high energy of 4 J/cm2.

200, the spallation area accounted for 34%, which far exceeded 20% of the total surface area and was identified that this coating

has failed. In order to thoroughly examine the failure behavior after 200 cycles, a deep study concerning the microstructure G


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Figure 10. Cross-sectional SEM images of coating C (a) and coating D (b).

Figure 11. Macro images of coating C after different thermal cycles; (a) SEM image of spalling area after 200 cycles; (b, c) magnified images of region A and region B.

and composition of the spalled region (red outline) was conducted. Observed from Figure 11a, the spalling surface was uneven and characterized as a delaminated fracture. High magnifications of regions A and B are illustrated in Figure 11b,c. It is clear that region A consisted of many small holes and oxidized multilayer flats with finger-like granules. The corresponding EDS result in Table 3 shows that a large amount of Cr-Co-O was detected, which can be confirmed as spinel inclusions. Region B was mainly composed of large ZrO2 particles, surrounded by small spinel particles (Figure 11c and Table 3). It is suggested that the spallation of coating C mainly occurred in the TGO/TC interface and, simultaneously, was coupled with a slight ceramic spalling.

Table 3. EDS Results of the Points in Figure 11 compositions (at. %) EDS points

O

1 2

39.13 74.82

Al

Cr

Co

21.13

39.74

Zr 25.18

Figure 12 gives the surface morphologies of coating D after different thermal cycles. As is apparent from the figures, in all cases, nominally the samples did not fail. After 150 cycles, some spots were found on the fringe of the surface. When sustained up to 200 cycles, only 1% of the evidently visible area on the edge of the surface was spalled. The magnified SEM image of H


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Figure 12. Macro images of coating D after different thermal cycles; (a) SEM image of spalling area after 200 cycles; (b, c) magnified images of region A and region B.

from higher repeated stress. Besides, the pre-existing cracklike discontinuities in the ceramic developed into fully open cracks. Most of them remained in the ceramic, but some have directly extended to the TGO layer and coalesced with the pores within the TGO layer. This, eventually, led to the formation of large cracks that distributed along the TC/TGO interface. By contrast, pores and cracks were absent at the TGO/BC interface. As thermal cycling proceeded to 150, as shown in Figure 13c, planar coarse cracks appeared at and near the interface of TGO and TC. After 200 thermal cycles, as shown in Figure 13d, the top coat was delaminated from the TC/TGO interface, indicating that the TBC has failed. Figure 14 shows the cross-sectional SEM images of coating D after different thermal cycles. It is clear that, after 50 cycles, a nearly continuous and compact TGO of predominantly Al2O3 was expected at the TC/melted layer interface, which was uniform in thickness (Figure 14a). Some cracklike discontinuities and porosities were also present in the ceramic coating. After 100 cycles, the TGO layer with dense morphology was still fully attached with no evidence of cracks and pores at the interface, whereas the pores within the ceramic coating were enlarged and merged (Figure 14b). After 150 cycles, a clear cracking was found initiating from and propagating along the interface between the TGO layer and the top coat, as shown in Figure 14c. As the thermal cycles increased to 200, the longest cracks in the ceramic topcoat penetrated into the TGO layer, as shown exemplarily in Figure 14d, or started in the TGO and subsequently shifted the cracking path toward the ceramic topcoat.

the spalling area (red outline), as shown in Figure 12a, reveals that the residual surface contained the bottom side (region A) and splash-type flat (region B). Further magnified SEM images in Figure 12b,c exhibit that the feature and constituent were very distinctive in comparison to the initial one. Given from the morphologies and EDS results (Table 4), region A was mainly Table 4. EDS Results of the Points in Figure 12 compositions (at. %) EDS points

O

Al

Cr

Co

Zr

1 2

61.02 66.18

21.43

7.53

10.03

4.65 33.82

composed of Al2O3 granules and doped with small amounts of Co-oxides, while region B contained regular lamellar ZrO2. It is clear that, compared with the coating C, the spallation of coating D mainly occurred in the ceramic lamellar structure and TGO/TC interface. Figure 13 shows the delamination crack path within the TBC and the TGO after different thermal cycles of coating C. When subjected to 50 cycles (Figure 13a), the TBC was composed of three distinct layers: the porous ceramic layer (the upper), containing plenty of crisscrossing cracklike discontinuities and porosities; the TGO layer with a highly undulated nature (the middle), characterized by two regions, predominantly mixed Co/Cr-O spinels adjoining to the top coat and a relatively small portion of Al2O3 closed to the bond coat; and the typical lamellar bond coat layer (the bottom), containing segmented Al2O3 veins. After 50 cycles, both the oxygen permeating inward and the metallic elements diffusing outward contributed to the formation of TGO, which has been discussed in our previous studies in detail.7,27 It is clear that the thickness and compositions of the TGO layer were rather less uniform, and a large amount of pores were visible within the TGO layer. After thermal cycles added up to 100 (Figure 13b), the TGO thickness increased obviously and local differences in scale thickness appeared. The pores were enlarged and merged, which happened at locations where the mixed oxides suffered

4. DISCUSSION 4.1. Evaporation Mode and Its Effects on the Surface Modifications. Referring to the formation of “hills and valleys”, deposited with abundant micron- and/or nanoscale particles demonstrated in Figures 4−7, it is clear that an evaporation−condensation mode played a dominant role in generating such surface morphologies, which is very different from the ones observed under the more conventional melting− solidification mode observed in HCPEB treated metals. During I


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Figure 13. Cross-sectional microstructures of coating C after different thermal cycles: (a) 50, (b) 150, (c) 150, and (d) 200 cycles.

has a significant “selective purification” effect during the melting or evaporated process, corresponding to the eruptive mechanism of subsurface hot spots at impurities (Co-based oxides). The absence of the Co-based oxides in coating B revealed in the XRD pattern may be an implication. Viewed from Figures 4, 7, and 8, nanostructures on the coating surface induced by HCPEB treatment can be divided into three main types: (1) Nanoparticles induced by vapor deposition. As discussed, abundant small bubbles were splashed out from the liquid melt due to the boiling impact under the evaporation mode (Figure 15a,b). In the subsequent irradiated process, these sputtering particles were redissolved in liquid melt and then re-evaporated, redeposited, and ultimately formed more stable condensation nucleuses. In the following cooling process, these nucleuses experienced growth and amalgamation stages (Figure 15c−e) and, consequently, formed nanoscale deposited particles massively, as shown in Figures 4f and 7a. It is worth noting that abundant Y-rich Al2O3 particles were also detected on the remelted surface, observed in Figure 4b,c,h. This aspect witnessed a selective evaporation of the specific chemical species that had a lower boiling point and higher activity. Under the evaporation mode by HCPEB irradiation, part of Al element was released from the CoAl phase and converted to the Al vapor. As analyzed previously,19

the HCPEB irradiated process, when the boiling point of the target material was reached, evaporation occurred within a certain depth, and simultaneously a large amount of microbubbles were generated in the liquid melt (Figure 15a). When approaching the top surface, these microbubbles showed a significant burst and splashed out abundant small droplets, which had a certain impact on the surrounding material (Figure 15b). It is similar to the water boiling, which can make the liquid surface accidented. Subsequently, the boiling and the recoil force from the evaporation, accompanied by the following rapid solidification, led to the formation and freezing of such a wavy morphology. Besides, Al2O3 veins that existed within the initial lamellar structures were mostly agglomerated at the concave part between the surface bulged nodules. Some residual evaporated−condensed droplets (i.e., Al2O3 droplets) also existed. Meanwhile, the parallel columnar grains perpendicular to the coating surface, observed inside the bulged nodules, were the consequence of directed solidification, which can be interpreted by the thermal conduction flowing mainly toward the base matrix during HCPEB irradiation.5,26,35 It was noted that, during this process, a spot of the CoCr phase that existed in the initial coating was melted and solid dissolved in the main phases. It was the reason for the disappearance of σ-CoCr peaks shown in Figure 2. Besides, previous studies20,21,24,25 have proved that HCPEB irradiation J


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Figure 14. Cross-sectional microstructures of coating D after different thermal cycles: (a) 50, (b) 150, (c) 150, and (d) 200 cycles.

the pressure was at least 1 order of magnitude higher than that of the normal value of 2−5 × 10−3 Pa, due to the removal of gas from the coating or possible evaporation process. In the present case, the residual amount of oxygen was inevitable. Consequently, a large number of Al2O3 protrusions or even thin film could be redeposited from the vapor and then condensed directly on the top surface. (2) Nanoparticles induced by rapid solidification. After HCPEB irradiation, there was a very high temperature gradient along the depth. The melt was stabilized and remained a highly undercooling condition, which led to an increase in the nucleation rate within the melt.17,36 The solidification and following cooling process occurred very rapidly, giving no time for the occurrence of grain growth, and therefore, nanoscale grains were formed, as shown in Figure 16. Nanograins observed in Figure 7c,d witnessed this effect. In addition, the continuity and integrity of dendrite growth could be broken by new heterogeneous nucleuses, and ultimately finer scale particles can be obtained, as shown in Figure 7b. (3) Nanoparticles induced by plastic deformation.

Figure 15. Formation mechanism of MCrAlY bond coat under the evaporation mode after HCPEB irradiation (the schematic illustration of nanoparticles induced by vapor deposition after HCPEB irradiation).

K


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Figure 16. Schematic illustration of nanoparticles induced by rapid solidification after HCPEB irradiation.

Figure 17. Schematic illustration of nanoparticles induced by plastic deformation after HCPEB irradiation.

Figure 18. Crack types of coating C formed during thermal cycling process.

therefore, refined the crystalline structures. Besides, the formation of sub-dislocation walls was considered to be associated with the repeating plastic deformation and high strain rate after multiple HCPEB pulses. In this case, dislocation walls and sub-dislocation walls intersected each other in order to achieve the desired grain refinement effect, as seen in Figure 17. 4.2. Failure Mode during Thermal Cycling. It is assumed that thermal cycling with dwell time consists of a superposition of two load components “thermal cycling” and “isothermal oxidation”.39 TBC spallation is, therefore, not a sudden fracture event at the end of life. It is a continuous damage evolution with the formation of microcracks and crack growth, and meanwhile, TGO growth plays a dominant role. Figure 18 illustrates the schematic illustration of various types of cracks in the TGO layer and its vicinity, summarized from Figure 13. There are four types of cracks in the case of coating C:

As mentioned, during HCPEB irradiation, each pulse of the electron beam can lead to rapid thermal cycle on the irradiated surface to which is also related to the dynamic stress. The coupling of the temperature field and stress field is the critical factor for the substantial modifications. Zou et al.34 proposed that the nonstationary thermal stress fields induced by a temperature gradient consisted mainly of quasistatic stress and thermal stress waves. The thermal stress propagates in the longitudinal direction with small amplitudes of about 0.1 MPa, while the quasistatic stress acts in the transverse direction with several hundreds of MPa, which is sufficiently high for the irradiated surface to deform. The observation of dislocation walls from Figure 8 was a signature of severe stress and plastic deformation caused by rapid heating and solidification, which was the result of reducing its total energy during the process of dislocation accumulation and rearrangement.37,38 Dislocation structures were transformed from the original line defects to dislocation walls. Continuing transformation can change the crystal orientations on each side of dislocation walls and, L


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Figure 19. Crack types of coating D formed during thermal cycling process.

aggressive fracture and eventual spallation of the TGO could happen. Figure 19 illustrates the schematic illustration of various types of cracks in the TGO layer and its vicinity in the case of coating D, summarized from Figure 14. In comparison with the initial ones, there are three types of cracks: Type A: cracklike discontinuities isolated remained within the TC. Type B: cracks along the TC/TGO interface. Type C: cracks penetrated into both the TC and TGO. As seen in the schematic illustration of the TGO layer, it is clear that the morphology of the TGO layer was similar to a sine curve. For simplicity, some researchers1 assumed the primary form of TGO to be accorded with a sine curve. They pointed out that the maximum tension stress appeared at the undulation crests while the maximum compression appeared at the troughs, and the stress distribution varied with the locations along the interface of the sine curve. However, this assumptive model possessed small wave width, large amplitude, and close waveform. Moreover, geometrical imperfections at the interface resulted in many stress concentrative points, which can make damage initiation and progression of microcracks occur. By contrast, the sine wave shape of TGO formed after HCPEB irradiation was completely different from the initial one, which had large wave width, small amplitude, and sparse waveform. As seen in Figure 19, the shape of TGO was highly regular and dense whether in the hills or valleys, which was related to the HCPEB irradiated effects. As discussed before, the thermal cycling process included TGO growth during “isothermal oxidation” and TGO shrinkage during “thermal cycling”. During the very beginning stage of oxidation of the first thermal cycle, a stable, continuous, slow-growing, and uniform TGO with strong adherent ability was accelerated to be formed. Several mechanisms can be called upon: (i) An entirely compact remelted surface with higher Al concentration provided a favorable precondition for forming such a protective oxide layer. As shown in Figure 4, the wavy aspect of a series of hills (interconnected bulged nodules) and valleys (black oxides Al2O3) created a completely compact surface “sealing layer”. This layer with an almost defect-free morphology eliminated the hindering effect of other oxides for Al2O3 grains, which was beneficial to the formation of such protective monolayer TGO. (ii) Nanoparticles, refined columnar grains, and plentiful deformation structures contributed to the selective oxidation of Al element, promoting the fast formation of the Al2O3 layer. These refined structures obtained within the remelted layer

Type A: cracklike discontinuities isolated remained within the TC. Type B: cracks located within the TGO. Type C: cracks penetrated into both the TC and the TGO. Type D: cracks along the TC/TGO interface. During the beginning stage of thermal cycling, the cracks at the interface mainly referred to type A. The opening of the preexisting discontinuities within the ceramic, associated with thermal shrinkage during the thermal spraying process, appeared to propagate and extend. After 50 cycles, bilayer TGO with porous structures was obtained, as shown in Figure 13a, the formation of which was related to the thermal sprayed defects like rough surface, spherical particles, and the underlying splats.27 As reported,40 the growth rate of these Co/Cr-O mixed oxides can reach at least 3 orders of magnitude compared to that of the monolayer Al2O3. They were loose and fragile, and their rapid growth was accompanied by a constrained volume expansion that led to compressive “growth” stresses. Besides, upon cycling, the thermal-expansion mismatch between the TGO and the BC layer led to very high “thermal” compressive residual stresses in the TGO. In this case, void formation occurred within the TGO, which was more susceptible to crack nucleation. As thermal cycling continued, the TGO layer grew thicker. In the combination of accumulated exposure time for isothermal oxidation and thermal cycling alternating stress, the thickness of Co/Cr-O mixed oxides further increased, and simultaneously, the voids grew and connected, forming transverse microcracks. As the cycles increased to 100, type B cracks within the TGO layer were formed by microcracking extension, followed by linking-up and then large crack propagation. Once the type B cracks got started, they extended to the TC layer driven by the stress fields with the assistance of the pores and type A cracks within the ceramic layer. Ultimately, type C cracks that penetrated into both the TC and the TGO layers occurred. With the increment of thermal cycles, type D cracks along the TC/TGO interface started to form. Because of the highly undulating nature of the TC/BC interface, out-of-plane stresses resulted from the vicinity of the TC/TGO interface: tensile at the undulation crests and compression at the troughs.1,41 Therefore, in-plane shear stress near the point of inflection showed favorable driving force for cracking.6 At some point, after reaching a critical shear stress, fracture along the TC/TGO interface occurred, forming the type D cracks. Various types of cracks were propagated, coalesced, and finally grew throughout the TGO layer and partly the TC layer. On this occasion, M


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(2) A physical model has been proposed to describe the evaporation−condensation effect, which played a dominant role in generating typical surface morphologies. Nanostructures on the coating surface induced by HCPEB treatment can be divided into three main types: nanoparticles induced by vapor deposition, nanoparticles induced by rapid solidification, and nanoparticles induced by plastic deformation. (3) The substantial differences concerning the final failure, TGO morphologies, crack growth behavior, and, finally, lifetime of APS-TBCs and HCPEB-TBCs were analyzed in detail. After 200 thermal cycles, the spalling area of HCPEB-TBCs accounted for only 1% of its total area while it was about 34% for APS-TBCs. The irradiated TBCs by HCPEB treatment exhibited superior thermal cycling performance. The modified microstructures, fine microcracks, and lower growth rate of the TGO layer after HCPEB irradiation played a dominant role in improving thermal cycling behavior of TBCs.

enhanced short-path diffusion of Al element, increased its diffusion coefficient, and accelerated its selective oxidation in a short time. (iii) Abundant Y-rich Al2O3 bubbles improved the self-repairing capability and adhesion strength of the TGO layer. As shown in Figure 4b,c,h, microscale Y-rich Al2O3 bubbles can be acted as nucleus and/or repairing points of the oxide layer, conducive to its transverse connection and growth. Besides, the enrichment of Y element had an important impact on suppression of the outward growth of TGO, prevented vacancy condensation at the TGO/metal interface, decreased the growth rate of the TGO layer, and increased its adhesive ability.42 The presence of the continuous and uniform α-Al 2 O 3 layer in Figure 14a witnessed these specific mechanisms. After 100 cycles, no cracks were formed except the type A. They failed to penetrate into the TGO, even when subjected to loop tensile stress of the wave crest. First, the nonuniform stress distributions from the initial undulated bond coat were eliminated by HCPEB irradiation, which was replaced with an entirely compact surface. Additionally, the refined remelted layer was expected by HCPEB treatment, and therefore, the grains of Al2O3 scale formed under this circumstance must be also nanosized, which was proved in our previous study.7,27 According to ref 43, the grain size of the oxide layer is inversely proportional to its oxide-creep strain. During the thermal cycling process, the nanosized oxide layer was favorable to the stress relaxation by the fast creep strain effect, thereby reducing the stress level. In addition, the stable Al2O3 layer slowed down the Al consumption and ensured the high Al concentration within the composition on the top surface. On this occasion, even if the destruction to the TGO layer occurred in the melted layer because of the porosity, cracks, or other imperfections, new protective oxide films could be produced immediately, which guaranteed the protective effect of the TGO. When the cycles added to 150, type B cracks were formed at the TC/TGO interface, illustrating that the tensile stress normal to the TGO/BC interface (due to the thermal-expansion mismatch) reached the fracture toughness, and finally led to the interface separation. Obviously, the irradiated samples had good interfacial adhesion because the interfacial imperfections were significantly less than the initial one, which was primarily responsible for TBC failure. The tensile stress between the peaks of the TC/TGO interface profile promoted the further growth of cracks that were initiated at the TGO with the increase of cycles and finally formed type C cracks. These three types of cracks were propagated, coalesced, and ultimately resulted in the TBC failure by some spallation of the top coat. However, it is worth noting that the microcracks existed in the ceramic coating almost parallel to the interface, keeping parts of the ceramic and TGO remained in the low valley avoided by the effects of them, which was a reason that ceramic lamellar structures can be observed in Figure 12d. It is clear that the irradiated coating D exhibited superior thermal cycling performance.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (Q.G.). *E-mail: [email protected]. Tel: +86 511 88790769. Fax: +86 511 88790083 (J.L.). ORCID

Qingfeng Guan: 0000-0002-1946-2040 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support from the National Natural Science Foundation of China (Nos. 51601072, U1233111, and 51575242), the Jiangsu Province Natural Science Foundation for Youths (No. BK20160530), the Jiangsu Provincial Natural Science Foundation for Distinguished Young Scientists (No. BK20140012), the China Postdoctoral Science Foundation (No. 2016M601730), the Postdoctoral Foundation of Jiangsu Province (No. 1601007C), and the Senior Talent Foundation of Jiangsu University (No. 15JDG149) is acknowledged.

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REFERENCES

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Thermal Cycling Behavior of Thermal Barrier Coatings with MCrAlY

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