Comparative study of the failure mechanism of atmospheric and suspension plasma sprayed thermal barrier coatings

Surface & Coatings Technology 370 (2019) 163–176

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Comparative study of the failure mechanism of atmospheric and suspension plasma sprayed thermal barrier coatings

T

M. Gizynskia,b, , X. Chena, N. Dusautoya,c, H. Arakia, S. Kurodaa, M. Watanabea, Z. Pakielab ⁎

a

National Institute for Materials Science (NIMS), 1-2-1 Sengen, Tsukuba, 305-0047, Ibaraki, Japan Faculty of Materials Science and Engineering, Warsaw University of Technology (WUT), Wołoska 141, 02-507 Warsaw, Poland c Science of Ceramic Processes and Surface Treatments (SPCTS), University of Limoges, 12, rue Atlantis, 87068 Limoges, France b

ARTICLE INFO

ABSTRACT

Keywords: Plasma spraying Thermal barrier coatings (TBC) Thermal shock/thermal shock resistance Mechanical properties Suspension plasma spraying

Among numerous works on suspension plasma sprayed thermal barrier coatings, only a few focused on the relationship between the thermal shock resistance of SPS coatings and their mechanical properties. In this study, several SPS yttria-stabilized zirconia coatings were deposited with various spraying parameters including plasma gas composition, stand-off distance, and suspension formulation, then they were subjected to thermal cyclic test at 1150 °C. The results showed that despite segmented/columnar microstructure most of them failed after fewer cycles than the benchmark APS coating with conventional lamellar microstructure. Changes in the mechanical properties of the coatings such as Young's modulus and toughness after exposure to elevated temperature were evaluated. Most of the SPS coatings showed unusual change of mechanical properties, i.e., significant decrease of either Young's modulus or toughness or both. It was shown that such peculiar evolution of mechanical properties cannot be explained by the tetragonal to monoclinic transformation of YSZ, but it is rather related to the formation of very fine pores smaller 1 μm after heat treatment. The evolution of mechanical properties was linked with the results obtained in the thermal cyclic test, showing that thermal cycling resistance of SPS coatings was predominantly controlled by the mechanical properties after heat exposure, for which such phenomena as sintering-induced embrittlement may give crucial influence.

1. Introduction Ceramic Thermal Barrier Coatings (TBCs) are widely used in hot sections of land-based gas turbines and aero-engines to increase the efficiency and extend the life of metallic components [1–3]. The primary function of the TBC system is to provide a low thermal conductivity barrier, allowing the temperature of metallic components to be kept moderate while increasing the operating gas temperature. Nowadays, YSZ-based TBCs are most widely used and fabricated either by Plasma Spraying (PS) or Electron Beam Physical Vapor Deposition (EB-PVD). Application of each method results in dissimilar and distinctive coatings microstructures. The PS coatings consist of lamellae of rapidly solidified YSZ particles called splats, and have low thermal conductivity due to a myriad of inter-splat voids and microcracks, while EB-PVD coatings typically show columnar microstructure, which exhibits low stiffness and high strain tolerance [4–7]. Recently, a new method which might combine the benefits of both PS and EB-PVD has emerged. In Suspension Plasma Spraying (SPS) [8–10] particles in the size range of sub-micrometer to nanometer are

⁎

suspended in water or organic solvent are used as a feedstock to realize fine microstructure with sub-micrometer splats and nano-porosity. These attributes of microstructure may greatly reduce the coating's thermal conductivity [11,12]. Previous studies have also demonstrated that various coating structures can be realized by the SPS process, including vertically cracked and columnar microstructure [13,14], which may result in significant improvement of compliance, durability, and lifespan of YSZ TBC. Despite these promising features and close attention paid to various aspects of SPS process, only a limited number of publications have focused on performance-related properties for TBC, such as thermal shock resistance or thermal stability [12,15–17]. There are even no unequivocal results showing the advantage of SPS coatings in thermal cycle lifetime over APS coatings. Results of several works dealing with thermal shock resistance of SPS coatings are summarized in Table 1. Results presented in Table 1 shows very big discrepancy between results for SPS coatings even if they were reported in the same paper. Such a considerable variation may indicate that the performance of SPS coatings are highly sensitive to their microstructure [21] and

Corresponding author at: National Institute for Materials Science (NIMS), 1-2-1 Sengen, Tsukuba, 305-0047, Ibaraki, Japan. E-mail address: [email protected] (M. Gizynski).

https://doi.org/10.1016/j.surfcoat.2019.03.013 Received 27 November 2018; Received in revised form 16 February 2019; Accepted 6 March 2019 Available online 07 March 2019 0257-8972/ © 2019 Published by Elsevier B.V.

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coatings and underlined that failure modes of the SPS and APS could be different. The SPS coatings failed by partial spallation of segments/ columns, while the APS coatings failed by the bridging of interface delamination close to the bond coat/top coat interface and the intrinsic segmentation cracks. Other investigators have reported the opposite trend of APS coatings outperforming SPS ones. Such results were presented by Seshadri et al. [21] and Ganvir et al. [22], the latter of which suggested that the reason could be the high magnitude of deposition stresses, randomly oriented columnar features or presence of interpass porosity bands. To the best of the authors' knowledge, there is a very limited number of published study aimed at understanding of SPS coating's failure in a comprehensive manner. Because the failure mechanism of SPS TBC has not been clarified yet, the current study aims to closely investigate this issue. Several SPS TBCs were sprayed with various spraying conditions and then tested in thermal cycle test. Microstructure, mechanical properties and residual stresses of the coatings were investigated in order to understand the thermal cycle performance of the SPS coatings.

Table 1 Comparison of thermal cycle resistance of SPS and APS coatings in the literature. Number of cycles before failure APS coatings

SPS coatings

The relative difference (APS = 0%)

290

110 135 270 375 450

−62% −53% −7% +29% +55%

290

330

+14%

280

−4%

355

+22%

310

+7%

360

+24%

165

−43%

150

−48%

85

165

+94%

50

40

−20%

25 3 320 260

−50% −94% −21% −36%

405

Notes

HVOF BC, nanosuspension PS BC, Nanosuspension HVAF BC, submicrometer suspension Solid content 25%, d50 ~ 45 nm Solid content 33% d50 ~ 200 nm Solid content 25% d50 ~ 200 nm Solid content 17% d50 ~ 200 nm Solid content 25% d50 ~ 500 nm Solid content 25% d50 ~ 400 nm Solid content 25% d50 ~ 45 nm

Porous Feathery Columnar Wide Columnar Segmented Lamellar Particles < 10 μm Particles < 3 μm

Cycle profile

Source

1 h at 1100 °C air jet cooling

[12]

1 h at 1100 °C air jet cooling

[18]

2. Experiments 2.1. Feedstock materials and spraying conditions

0.5h at 1100 °C water cooling 24 h at 1100 °C

[19,20]

1 h at 1100 °C air jet cooling

[22]

Two types of suspension of YSZ (8 wt% Yttria-Stabilized Zirconia) were used in this study. One of which was ethanol-based suspension with the powder loading of 25 wt%. For the suspension preparation commercially available fused and crushed powder (YSZ, Imerys Fused Minerals, Laufenburg, Germany) with a particle size < 2 μm (Fig. 1) and reagent grade ethanol (99.5%) as the solvent was used. The suspension was prepared by mixing the YSZ powder with ethanol, followed by ball milling for 24 h by using ZrO2 balls and polyethyleneimine as a dispersant. The ball milling allowed to break up the particle agglomerates before spraying. The second suspension was water-based suspension prepared by H.C. Stark GmbH, of which solid content was 30 wt% and D50 of YSZ powder was 0.9 μm. Plasma spraying of these suspensions was carried out by an Oerlikon Metco Multicoat plasma spray unit with a Triplex Pro 210 gun. The plasma torch nozzle diameter was fixed at 9 mm. Suspension in the tank was pressurized with argon and radially injected into the plasma jet through an orifice with a diameter of 200 μm. Detailed parameters of the SPS process are listed in Table 2. Two plasma gas compositions, two stand-off distances were chosen in order to find the effect of the spraying parameters on the microstructure and thermomechanical properties of the resultant coatings. A plasma jet generated by an argon‑helium (C1) mixture operated at a power of 44 kW and another plasma jet formed by an argon‑hydrogen mixture were used for spraying, the latter of which should provide a higher thermal conductivity of the plasma and

[21]

mechanical properties, which in turn can significantly vary depended on deposition parameters such as substrate [12], size of ceramic powder [18,22], and concentration of solid phase in suspension [18]. Abundance of factors affecting SPS coating's performance makes difficult to determine differences between failure modes of SPS and APS or EBPVD. However, some researchers made an effort to explain observed differences Zhai et al. [20], who have reported a superior lifetime of SPS coatings due to a columnar or segmented microstructure, which reduces thermal stress in a manner similar to the EB-PVD coatings. Also, results presented by Zhao et al. [19] showed a longer lifetime of SPS

Fig. 1. YSZ powder used for the ethanol-based suspension preparation (a), the particle size distribution (b). 164

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Table 2 APS and SPS parameters for YSZ coating deposition. Coating

Method

Feedstock

Stand-off

Plasma composition

Plasma gun power

APS C1-50E C1-70E C2-50E C2-70E C2-70W

APS SPS SPS SPS SPS SPS

204NS powder Ethanol-based suspension Ethanol-based suspension Ethanol-based suspension Ethanol-based suspension Water-based suspension

100 mm 50 mm 70 mm 50 mm 70 mm 70 mm

Ar/He Ar/He Ar/He Ar/H2 Ar/H2 Ar/H2

26 kW 44.2 kW 44.2 kW 53.4 kW 53.4 kW 53.4 kW

enhance heat transfer from the plasma to the suspension. When argon‑hydrogen plasma (C2) was used the power level rose to 54 kW. As a benchmark coating, we used conventional YSZ coating deposited by APS using commercially available Sulzer Metco 204NS Hollow Sphere Oxide Powder (HOSP) with particle size D50 of 40 μm as the feedstock powder. For the APS process a Praxair SG-100 plasma gun was used. The target thickness of both APS and SPS coatings was around 350 μm. The specimens tested in the thermal cyclic test (TCT) were prepared in the following manner. Disc-shaped substrates with a diameter of 30 mm and thickness of 3 mm, made of Inconel®-718 alloy were roughened by grit blasting with alumina particle, final toughness was around Ra = 3 μm, then cleaned and degreased with acetone before spraying. Later the superalloy substrate was coated with CoNiCrAlY alloy (Oerlikon Metco, AMDRY 9954) by High-Velocity Oxy-Fuel (HVOF) method. The roughness of the as-sprayed bond coating was in the order of Ra = 5–7 μm. The bond coating was vacuum heat-treated at 1000 °C for 4 h in order to improve adhesion to the substrate and residual stress relief. Prior to YSZ deposition, the samples were degreased and cleaned in acetone using ultrasonic for about 5 min. Before spraying substrates were heated by several passes of the plasma gun over the substrates to a temperature around 200 °C in order to improve adhesion. The torch transverse speed during SPS process was 1000 mm/ s, while during APS it was 200 mm/s. The microstructural investigations, residual stress, and toughness test were carried on coatings deposited in the same manner as the TCT samples. The bending test was carried on coatings deposited onto machined strips of Inconel®-718 substrate with dimensions of 1.9 × 10 × 100 mm3. In order to reduce the complexity of the measurement and facilitate calculations of Young's modulus, the bond coating was not applied to bending test specimens. In order to investigate the effects of high-temperature exposure on the microstructure and mechanical properties of the coatings, some specimens were heat-treated at 1150 °C for 5, 10 or 50 h in the air atmosphere.

bending (4 PB) setup using a Shimadzu Epigraph testing machine, equipped with a 10 kN load cell. The support and load span were 80 and 40 mm, respectively and the center point deflection was monitored with a deflection sensor with 1 μm resolution. The coating was on the compression side of the specimen. This testing procedure is similar to those employed by other researchers [23,24]. Loading ranged between 0 N and −240 N in four consecutive cycles at a strain rate of 3 × 10−6 s−1. The recorded stress and strain were employed to find mean stress and strain in the TBCs and further to calculate the mean Young's modulus. Coatings' moduli were evaluated at a strain of 0.3%, within the elastic limit of the substrate. Using the deflection u measured at the center the apparent Young's modulus Eapp of the composite beam is given by [25]:

Eapp =

d (3L2 48I

4d 2)

P u

(1)

where d is the distance between the inner and the outer loading point, L is the distance between the two outer spans, P is the applied force and I is the second moment of area. The coating's modulus Ec could be estimated by using Eq. (2):

Eapp =

(

3 12Es yNA + 1

Ec Es

) (t

s

yNA )3 +

t3

Ec (t Es

yNA )3 (2)

and the neutral axis yNA is given by:

yNA =

1 2

Es ts2 + Ec (t 2 ts2 ) Es ts + Ec tc

(3)

where: Es is the substrate Young's modulus, which for Inconel®-718 is 205 GPa (own study-4-point bending of the substrate without coating), t, ts, and tc are the thickness of the whole specimen, the substrate and coating, respectively. The cohesion of a top coat can be experimentally evaluated by measuring the resistance of the ceramic layer against initiation and propagation of cracks when locally loaded. In this study, Vickers indents were generated with gradually increased forces. Cracks resulting

2.2. Microstructure evaluation Cross-sections and fractured surfaces of the samples were examined using a scanning electron microscope (JEOL, JSM-6010LA). For the cross-section analysis, as-sprayed samples were cut with a SiC cutting wheel, mounted in epoxy resin and subjected to standard metallographic preparation with grinding using emery papers with gradation from #120 to #1200, polishing using diamond suspensions with grain diameter of 3 μm and 1 μm, and final polishing step using 0.2 μm alumina suspension. Prior to observation, the samples were coated with a thin Pt layer. Coating porosity was evaluated by image analysis of at least 5 cross-sections images taken with a magnification of 5000×. The Image J 1.51j8 (National Institute of Health, USA) software was used to analyze the image, for each image the grayscale threshold was set manually in order to obtain data covering the widest possible range of pores. 2.3. Mechanical tests

Fig. 2. Method of the critical load and crack length determination in the indentation method.

The coatings' Young's moduli were measured by a four-point 165

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from the application of local force were observed on the indents' corners. The method relies on the determination of the critical load Pc required to initiate a crack with a critical size cc. If the load is lower than Pc, any cracks are not observed in the indent vicinity. If the applied load is higher than Pc, cracks starting from the indentation are observed and the crack length increases with the applied force. In order to find the critical values Pc and cc, the crack length – load and indent halfdiagonal – load data were plotted on a single log-log graph (Fig. 2). The intersection of a “crack” line and “hardness” line determines the critical values. A similar approach was employed by e.g. Sniezewski et al. [26] and Lesage and Chicot [27] to measure adhesion in bilayer systems. Once the critical values Pc and cc are determined, the “Vickers” toughness KVIF can be calculated with the following formula [28]:

KVIF =

E1/2Pc H1/2cc3/2

APS coatings typical lamellar microstructure without vertical cracks is shown in Fig. 3a. At higher magnification pores and intersplat voids are visible. Microstructures of C1-50E and C2-70W coatings are presented in Fig. 3b and f, respectively, both of which exhibit relatively dense microstructure with deep vertical cracks (labeled “A”). Fig. 3b inset reveals very fine porosity in coating C1-50E, the largest pores are about 1 μm in diameter, whereas the enlarged image of C2-70W coating shows coarser splats and pores. During suspension spraying it is known that a jet of water-based suspension breaks up into bigger droplets than in the case of ethanol-based suspension [35,36], which could be an origin of the generally coarser microstructure of C2-70W coating. The microstructure of C1-70E (Fig. 3c) coating is entirely different from the ones described so far. The coating is made of columns (“B”) separated by porous boundaries (“C”), and porosity bands (“D”) are clearly visible inside individual columns. The zoom-in image reveals a wide size distribution of pores as well as loosely bonded spherical particles (“E”). Fig. 3d reveals dense columnar microstructure of the C2-50E coating. Unlike the highly porous C1-70E coating, columns of the C2-50E coating are dense and only some porosity bands were found. High magnification image confirms dense microstructure and reveals horizontal voids between splats. Fig. 3 e shows the microstructure of C270E coating, whose appearance is similar to C1-70E, i.e., the coating possesses columnar microstructure which can be attributed to the long stand-off distance and low particle velocity at impact. Many defects are visible in the microstructure, such as poorly bonded spherical particles and significant porosity. As compared to C1-70E coating, however, C270E exhibits narrower columns and the porosity bands and intra-columnar gaps are not as well developed as in C1-70E. The microstructure developed in the examined SPS coatings well agree with the structures proposed by VanEvery et al. [37]. In case of the C1-50E and C2-50E most likely the build-up mechanism resulting in the formation of the flat, continuous coating is most prominent in this case as stand-off distance is short and impacting droplets have significant momentum in a direction perpendicular to the substrate's surface. Thus, particles impact the substrate almost perpendicularly with no or very small velocity parallel to the surface. In such case, the vertical cracks may be formed during cooling while thermal misfit stress is arising and eventually leads to through-thickness cracking. A similar mechanism could be applied to the formation of C2-70 W- longer spraying distance and therefore particles' lower velocity is compensated by higher mass of particles formed from water-based suspension, thus the ceramic particles remain high momentum at the impact. Longer spraying distance applied during spraying of C1-70E and C2-70E coating resulted in lower velocity and momentum of the particles, therefore near the surface due to drag the particles gain component of the velocity parallel to the substrate. As a result deposited coating has columnar microstructure or structure with vertical porosity bands. Image analysis of these coating cross-sections revealed various levels of porosity as shown in Table 3. The reference APS coating shows around 11% porosity, which is within a typical range for APS YSZ coatings. Since the SPS coatings possess a significant portion of submicron porosity, which is very hard to closely analyze by image analysis, due to the closing of the fine pores caused by mechanical polishing. Thus, it should be noted that the measured porosity can be somewhat lower than the true values. Coatings C1-50E, C2-50E, and C2-70 W were dense and their porosity was around 5 to 7%. Long spraying distance resulted in the significantly higher porosity of C1-70E and C2-70E coatings around 20%. In C1-70E coating, pores are located mostly in the well-developed porosity bands inside columns and intracolumnar highly porous spaces. 20%. Comparing to the C1-70E coating, the C5-70E coating exhibits narrower columns, the porosity bands, and intra-columnar gaps are not as well established as in C1-70E. The microstructure of C2-70E does not reveal such regular arrangement of pores and the large fraction of the porosity is located randomly inside columns. The difference may be caused by various temperature and velocity reached by ceramic particles under C1 and C2 conditions [14].

(3)

where α is a constant (α = 0.016 ± 0.004 [29]), E and H refer to the Young's modulus and hardness of the tested material, respectively. Hardness was measured on the coatings' cross-sections using a Vickers type indenter at a loading of 0.98 N. 2.4. Residual stress The residual stresses in both YSZ top coat and TGO are usually considered as the driving force for TBC failure. The stress in the top coat was measured using a method based on Raman spectroscopy as employed by Limarga et al. [30] and Zou et al. [31]. Since the coating is much thinner than the substrate, the coating stress can be assumed as a biaxial stress σin-plain: in plain

=

1

C

2

EC Edense

where Δν refers to the frequency shift from the reference stress state, υC is the Poisson's ratio of the YSZ top coat (in this study υC = 0.22 [32,33]), EC is the Young's modulus of the tested coating and Edense is the Young modulus of the bulk YSZ (Edense = 210 GPa [30]). Π refers to the piezospectroscopic coefficient. Another study of Limarga et al. [34] shows that a peak located around the wavenumber of 465 cm−1 is the most sensitive one to stress. The piezospectroscopic coefficient of this peak is 2.01 cm−1/GPa. Raman measurements were made using a Micro Raman Spectrometer (Horiba-Jovin Yvon T64000) with an excitation of 514.532 nm from an argon‑krypton laser. The intensity of the laser was 100 mW. All spectra were fitted by a Lorentzian-Gaussian function to obtain the peak position. 2.5. Thermal cycle test Thermal cyclic test is one of the standard lifespan assessment tests for TBCs. In this study, TCT samples were heated in a furnace to 1150 °C for a period of 10 min followed by rapid cooling using compressed air to approximately 60–100 °C within 5 min. After each cycle, an image of the specimen surface was taken by a camera allowing delamination progress to be observed and cycles to failure to be counted. A thermocouple was used to record the specimen temperature profile during the test. Once every 100–200 cycles thermovision image was taken by an IR camera (Avio InfReC R300), by which the surface temperature of each individual specimen can be evaluated. Samples were repeatedly cycled until failure. In this study spallation of 10% of the coating area was defined as the failure criterion. 3. Results and discussion 3.1. Microstructure Cross-sectional microstructures of the as-sprayed coatings are presented in Fig. 3 with high-magnification images in the insets. For the 166

Fig. 3. Cross-section of as-sprayed APS (a), C1-50E (b), C1-70E (c), C2-50E (d), C2-70E (e) and C2-70 W (f) coatings.

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Sintering of YSZ usually results in healing of microcrack, closure of pores and enhanced intersplat bonding, which should lead to increase of both stiffness and toughness [44]. In fact, such change was observed in the case of the APS coating (Fig. 5a) - sintering led to a notable increase in toughness from 2.48 in the as-sprayed condition to 2.93 MPa m1/2 after 50 h at 1150 °C. The increase of Young's modulus was moderate, the reason could be a relatively high density of the APS coating. Surprisingly, most of the SPS coatings showed an opposite trend as can be seen Fig. 5b for C1-50E – resulted in a decrease of both toughness and stiffness. The initial value of the SPS coatings' toughness was quite high- around the value of 2.7–2.9 MPa m1/2. It is well known that SPS YSZ coatings are built of continuously stacked splats, which are much finer than APS YSZ splats. These fine SPS YSZ splats could realize better bonding among each other in the as-spray condition because in SPS coatings there are no long intersplat boundaries, which may act as an easy path for crack propagation [42]. Similar changes were previously reported by Chen et al. [45], in that study 100 h long aging of the SPS made YSZ coatings at 1200 and 1400 °C led to drop of Young's modulus and hardness. Also, similar heat-treatment inducted decrease of the mechanical properties (hardness, Young's modulus) was previously reported for EB-PVD coatings by Renteria et al. [46] and Zotov et al. [47]. In contrast to the C1-50E coating, the C1-70E SPS coating showed typical sintering-induced change of both stiffness and toughness (Fig. 5c). The mechanical properties of the SPS coatings were degraded except for C1-70E and C2-70E as summarized in Fig. 5d. Phase transformation can be ruled out as a possible mechanism for the observed degradation of the mechanical properties as any transformation in APS and SPS coatings (e.g. t’-YSZ → m-YSZ) was not observed in the XRD spectra. The change in the mechanical properties should, therefore, be related to microstructure evolution during heat treatment. The SPS coatings in the as-sprayed condition show a variety of small pores and inter-splat voids with size in the range of tens or hundreds nm (Fig. 6a). After aging at 1150 °C for 50 h, cross-sectional examination revealed pore rearrangement into evenly distributed globular closed pores with the size of few-hundred nm (Fig. 6b). A similar change in the porosity of SPS YSZ coatings was reported by Chen et al. [45] and Bacciochini et al. [48]. The only exception within the SPS coatings was a C1-70E coating, which showed “typical” enhancement of the mechanical properties. Heat treatment of the C1-70E coating led to the significant densification of the intra-columnar spaces and microcracks (Fig. 6c), reduction of the small pores density – mainly large pores with a diameter over 1 μm emerged after 50 h exposure at 1150 °C (Fig. 6d).

Table 3 Coating porosity values as measured by crosssection image analysis. Coating

Porosity [%]

APS C1-50E C1-70E C2-50E C2-70E C2-70W

11.3 ± 2.8 4.8 ± 0.8 21.7 ± 5.3 6.3 ± 1.7 18.2 ± 4.9 7.0 ± 1.9

Fig. 4. Force-displacement curves recorded during bending of C1-50E coating in the as-sprayed and heat-treated conditions at 1150 °C for 5, 10 and 50 h. Notice that the curves for 5 and 10 h almost completely overlap with each other.

3.2. Mechanical properties Fig. 4 shows representative results of the bending test of coatings in the 4-point bending setup with the coatings in compression. The plot shows force-displacement curves for C1-50E coatings in the as-sprayed condition as well as after heat treatment at 1150 °C for 5, 10 and 50 h. All the coatings showed some anelastic behavior (hysteresis and nonlinear elastic behavior). Unexpectedly, heat treatment led to decrease of elastic modulus as clearly seen in Fig. 4, i.e. the load-unload curves became less steep with prolonged heat treatment time. Table 4 presents the average thickness of the coatings and their mechanical properties – Young's modulus and fracture toughness- measured for the as-sprayed condition as well as for coatings heat-treated at 1150 °C for 5, 10 and 50 h. One can expect that coating with low stiffness [37,38,40] and high toughness [41–43] would perform well at cyclic tests, thus changes in the TBC mechanical properties are presented in the E-KVIF plots. The plots on Fig. 5 show tendencies of Young's modulus and coating toughness evolution after heat treatment at 1150 °C. It is well established that coatings undergo sintering at the high temperature.

3.3. The results of thermal cycle test The results of thermal cycle test are presented in Fig. 7, in terms of cycles to failure for each coating. Comparing the APS and SPS coatings, it can be seen that the lifetime of SPS coatings except C1-70E was significantly shorter than the APS coatings. Despite of apparently favorable initial microstructures and mechanical properties as TBC, lifespans of the C1-50E, C2-70E, and C2-70W were in the same range of values- between 400 and 800 cycles. Based on the reported results of

Table 4 Mechanical properties of the tested coatings in the as-sprayed and heat-treated conditions. Coating

Thickness t [μm]

APS C1-50E C1-70E C2-50E C2-70E C2-70 W

358 379 421 335 316 406

± ± ± ± ± ±

8 20 25 2 7 12

As-sprayed

5h 1/2

E [GPa]

KVIF [MPa m

65 70 45 70 42 65

2.5 2.91 2.05 2.79 2.12 2.95

]

10 h

E [GPa]

KVIF [MPa m

73 60 52 66 45 67

2.60 2.68 2.22 2.83 2.18 2.24

168

1/2

]

50 h

E [GPa]

KVIF [MPa m

76 59 52 65 48 68

2.7 2.40 2.32 2.15 2.15 2.66

1/2

]

E [GPa]

KVIF [MPa m1/2]

77 57 54 70 53 63

2.93 2.03 2.56 2.12 2.21 2.36

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Fig. 5. Evolution of Young's modulus and toughness by heat treatment at 1150 °C: APS (a), C1-50E (b), C1-70E (c) and summary plot showing mechanical properties of all tested coatings before heat treatment and after 50 h at 1150 °C (d).

Fig. 6. High-magnification secondary-electron images of the intracolumnar microstructure of the C2-50E coating before (a) and after 50 h long heat treatment at 1150 °C (b), the microstructure of the C1-70E before (c) and after (d) the heat treatment.

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about 50% shorter time. 3.4. Fractography The TGO layer and its vicinity are believed to be crucial for the performance of TBC's as these are sites of crack nucleation and growth [51]. In Fig. 8a, the thickness of TGO in the TBC samples after failure in TCT is plotted against the square root of the cumulated time at high temperature. The error bars correspond to the standard deviation of measurements. The plot shows that the oxidation kinetics of TGO growth was likely to follow the same oxidation law for all tested samples. It is expected that the oxidation kinetics of TGO growth followed the parabolic oxidation law: Fig. 7. Lifetime measured by the thermal cyclic test.

= kp where δ is TGO thickness, τ is oxidation time, and kp is the parabolic rate constant. This observation is reasonable since the APS and SPS YSZ coatings are permeable to oxygen, the bond coat oxidation rates were very similar. Many authors developed models of TBC's lifetime where TGO growth and thickness are key factors [52–54], though the TGO seems not to be a crucial factor affecting TBC durability in this study. The images in Fig. 8b and c show cross-sections of the TGO layer in C250E and APS specimens, respectively. At the initial stage of oxidation, the main constituent of the TGO was α-Al2O3, what is shown in Fig. 8b and was confirmed by additional EDS and XRD analysis (not presented here). After longer oxidation other oxides such as Cr2O3, NiO or spinels were formed on the top of alumina layer. Fig. 8b reveals that a delamination crack propagated predominantly in the SPS YSZ coatings, so the mechanical properties of the ceramic layer itself should be the major factor in the failure process. In contrast to SPS, the delamination crack in the APS sample ran along the TGO or TGO/TC interface as shown in Fig. 8c. Fig. 9 presents specimens' surfaces after delamination. Images shown in Fig. 9a–c were taken using back-scatter electron mode, which makes use of the various chemical composition of the specimen.

DVC coating lifetime [49] one can expect that coatings with discontinuous segmented/columnar structure would have a superior lifespan. Also, high porosity level is generally known to be a factor promoting good resistance to thermal cycle fatigue [50]. None of these factors observed in these SPS coating seem to have improved the performance of the SPS coatings. The C2-50E coating deposited using “hot” argon‑hydrogen plasma and short stand-off distance failed after just 260–310 cycles. One can find this result surprising as C2-50E coating has similar microstructure to C1-50E or C2-70W coatings, i.e. deep vertical cracks and dense slabs with very fine porosity. Also, based on the initial mechanical properties the C2-50E had been expected to have superior resistance to thermal shocks. The C1-70E coating is the only SPS coating with similar or slightly higher performance than the APS coating. The long lifespan can be attributed to its columnar microstructure with high porosity and compliance. This coating was deposited with the “cooler” argon‑helium plasma and it is noteworthy that coating with similar microstructure and mechanical properties but sprayed with the “hotter” AreH2 plasma, namely C2-70E, failed after

Fig. 8. The linear relationship between TGO thickness and the square root of accumulated dwell time (a), TGO appearance of C2-70E (b) and APS (c) coatings. 170

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Fig. 9. Specimen surface after TBC delamination in TCT - APS (a), C1-70E (b) and C2-50E (c and d).

Generally, the light gray areas correspond to areas rich in heavy elements (not presented here EDS analysis showed a high concentration of Zr and Y), while darker regions are dominated by lighter elements (mainly Al, Co, Cr, Ni). Since the dark areas correspond to TGO, they allow estimating the area fraction of TGO in the fractured surface, i.e.,

so-called black failure. The surface after delamination of APS coating (Fig. 9a) shows a large area of TGO and islands of YSZ, which probably remained in the valleys of asperities at the TGO/TBC interface. In contrast, the fractured surfaces of the SPS coatings as shown in Fig. 9b and c were predominantly white regardless of lifespan indicating that

Fig. 10. Change of Raman peak position of free-standing and APS coat attached to the substrate after high-temperature exposure at 1150 °C (a). The residual stress of the YSZ top coats and its dependence on the isothermal heat treatment time for APS, C1-50E and C1-70E (b). Horizontal crack observed in C1-50E coating heat treated for 10 h (c). 171

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the crack propagated mostly in the ceramic top coat. Fig. 9d shows that cracks ran along the splat boundaries, uncovering very fine spherical YSZ crystals. White failure is usually observed in TBC systems with short thermal cycles, TBCs exposed to high temperature for a long time tend to fail usually by TGO/TC delamination [44] due to TC sintering and toughening.

residual stress should have limited driving force for delamination and contributed to the observed longer lifetime of C1-70E coating under thermal cycling. 4. Discussion There exists a large number of reports describing properties affecting the lifespan of TBCs under thermal cyclic or thermo-mechanical fatigue conditions. Many authors have emphasized the role of TGO growth and thermal expansion mismatch between the bond coat, TGO, and YSZ top coat, leading to high magnitude of residual stress in the TGO layer (1 GPa and more) and subsequent crack nucleation and growth [50]. Other researchers have found the cause of failure in sintering driven stiffening [64], leading to increase of delamination driving force. Many studies on TBC performance have reported a good correlation between mechanical properties in the as-sprayed condition (Young's modulus [37,38,40], toughness or interfacial toughness [41–43]) and coating resistance to thermal cycles. In this study, however, none of these approaches seem to be applicable. All tested TBCs were deposited onto the same HVOF-sprayed CoNiCrAlY alloy. At hightemperature exposure, the TGO growth kinetics obeyed the same parabolic relation regardless of the top coat microstructures (Fig. 8). Also, the fact that the predominant mode of failure in SPS coatings was white failure (crack propagating inside the ceramic coating) strongly suggests that the properties of the top coatings affected the lifetime more significantly than the toughness of the TGO/top coat interface. Especially, the peculiar evolution of the mechanical properties of the SPS coatings makes “stiffening approach” useless. In order to better understand the lifetime of the tested SPS coatings, we propose the following explanation. The driving force of TBC delamination is strain energy U stored during thermal cycles. The stored energy is a consequence of the CTE mismatch between the ceramic coating (αTBC = 11.5·10−6 K−1) and the metallic bond coat (αBC = 14·10−6 K−1). The energy per unit area in a thin planar film in one cycle can be given as:

3.5. Residual stresses Residual stresses accumulated in the TBC act as a component of the crack driving force and can facilitate delamination of the top coat. The residual stress is found by determining the difference in the peak position between free-standing YSZ layer and the YSZ top coat attached to the substrate. Fig. 10a is an example of such peak position change and shows a change of peak position after exposure of APS coat at 1150 °C. Generally speaking, with an increase of the heat treatment time, the Raman peak position shifted to lower wavenumber. For the freestanding TBC, the shift of the Raman peak can be attributed to decreasing O2−anion disorder [55–56]. After 20 h at 1150 °C, the peak reached a stable position around 464.5 cm−1. Due to the residual stress present in the supported TBC, the peak of the constrained coating takes a different position. Based on the difference in the peak positions between the supported and free-standing YSZ coatings, the stress in the YSZ top coat was evaluated for APS, C1-50E and C1-70E. Fig. 10b shows such stresses as a function of thermal exposure time. For the APS top coat in the as-sprayed condition, residual stress was highly tensile at around 150 MPa. Such a high value may be attributed to the predominant quenching stresses (tensile) and low level of thermal stress (compressive) [57]. Low thermal stresses could be a result of long spraying distance and low enthalpy of the plasma, so the temperature of ceramic particles and substrate during the spraying process was low. The particle temperature was not recorded in this study, but thermocouple was attached to the back-side of the substrate during spraying, during APS spraying substrate temperature reached merely from 200 to 250 °C. For comparison, the substrate temperature in SPS processes exceeded 600 °C. The residual stress sharply decreased to about -95 MPa (compressive) during initial 20 h of heat treatment and then remained at a similar level after 50 h of exposure. At 1150 °C any internal stress should be relaxed by the plastic deformation and creep of the BC. During cooling, however, the coating should be placed under an increasing compressive residual stress induced by the thermal expansion mismatch between the ceramic coating and the metallic bond coat/ substrate couple. Usually, increase of Young's modulus caused by sintering leads to further increase of the compressive residual stress [58]. Similar values and evolution paths were reported previously by Chen et al. [59], Hamacha et al. [60], Jordan and Faber [61] and Texeira et al. [62]. Both tested SPS coatings exhibited compressive residual stress around −50 MPa in the as-sprayed condition due to the high specimen temperature during spraying. For C1-50E the isothermal heat treatment up to 10 h led to a steep increase of compressive residual stress, but then stress relaxation took place and after 20 h residual stress reached a low level of −20 MPa. This relaxation could be caused by the formation of short cracks in the vicinity of TGO/top coat interface as it is presented in Fig. 10c. Longer exposure led to a moderate increase of compressive stress, which could be explained by competition between stress accumulation by the mismatch of thermal strain and crack growth. Finally, isothermal heat treatment can lead to delamination; in some specimens (some C1-50E and all C2-50E) failure occurred already after 50 h. The residual stress level for C1-70E, on the other hand, was compressive with a rather moderate value from −25 to −50 MPa independent of time. Many factors, such as low Young's modulus, high coating porosity [63], highly-developed columnar microstructure can be the cause of low-level residual stress. After 10 h, some stress relaxation occurred, but no obvious vertical or horizontal cracks were observed and no phase transformation took place either. Such moderate

U=t

( )d

1 tC EC ( )[( C 2 1

BC )

T ]2

where tC is the coating thickness, σ and ε are thermal mismatch stress acting on the coating and its thermal strain during the cycle, respectively; EC Young modulus, which depends on heat treatment time τ, αC coefficient of thermal expansion of the YSZ coating, ΔT temperature change in the thermal cycle, υ TC Poisson ratio (υ = 0.22 [32,33]). For simplicity, we assumed that there is linear elastic relationship between stress and strain. Several simplifications were made at this point but they do not have a strong effect on the general concept of the discussion. First of all, since the bond coat in all tested samples was the same HVOF sprayed CoNiCrAlY alloy, properties of the bond coat, such as CTE or creep behavior was not taken into account. Also, uniform stress distribution across coatings thickness and the linearly elastic stressstrain relationship of TBCs were assumed over the whole range of temperature ΔT. Finally, in the calculation measured values of Young's moduli were considered while the YSZ stiffness varies with temperature and changes during thermal cycling. On the other hand, propagation of the crack in the ceramic requires energy G dissipated due to the formation of new surfaces and other dissipative processes such as non-elastic deformation in the crack tip vicinity. The energy Gc required for crack propagation is given by:

Gc =

(1

KVIF ( ) 2 ) E( )

where KVIF is represented by the TBC toughness measured by indentation. In order to compare the two counteracting energies, i.e. the crack propagation driving force U and the TC toughness acting against crack growth Gc, a parameter ζ was introduced as the ratio: 172

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Fig. 11. The relationship between a number of cycles to failure and ζ parameter for the as-sprayed (a) and coatings heat treated for 50 h (b) at 1150 °C.

Fig. 12. Schematic of a design of two different TBCs with (a) non-recommended and (b) recommended morphology which can result in long lifespan in TCT. (c) Undesired high-temperature evolution of microstructure resulted in the formation of arrays of submicron pores. (d) Desired change leading to a dense structure with micrometer-sized pores.

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Table 5 The sequence of changes which took place in the APS and SPS coatings. Short-living TBC coatings (SPS coatings except for C1-70E)

Long-living TBC coatings (A1 and SPS C1-70E coatings)

During spraying ceramic particles is well heat treated due to the application of highly conductive plasma and short spraying distance

During spraying heating of ceramic particle is limited due to the application of plasma with limited heat conductivity and long spraying distance, thus particles temperature at impact could be considered as moderate The microstructure of coatings possess many defects such as big pores and relatively wide intersplat cracks Heat treatment leads to spherization of the wide intersplat cracks and formation of the dense microstructure with big spherical pores The change of microstructure causes stiffening and toughening of the coatings, as a result, the ζ parameter shifts toward lower values Since the coating is tough the failure starts after a significant number of thermal cycles when TGO layer is thick and delamination crack propagates along TGO

Inside columns/segments microstructure of the coatings contains fine defects – small pores and narrow intersplat cracks Heat treatment causes bridging of narrow intersplat cracks and formation of arrays of nanopores The change of microstructure leads to significant embrittlement of the coating and resulting in an increase of the ζ parameter Low toughness is a parameter affecting the lifetime and failure of the coating, thus failure occurs even if the TGO layer is thin and delamination crack propagates predominantly in the YSZ layer

( )=

Paper of Ganvir et al. [51] summarized studies focused on the effect of SPS coating's microstructure on its thermal shock resistance. The report demonstrates a key role of the TC microstructure consisting of narrow columns, empty intercolumnar spacings, nano-pores, and βphase rich bond coating in thermal-shock resistant SPS coatings. With our results in the present work, it is possible to complement the suggestions given in the mentioned paper. The schematics of undesired and desired of SPS deposited TBCs are shown in Fig. 12a and b, respectively. A segmented morphology with dense structure (low porosity) did not produce good results in the thermal cyclic test, despite numerous reports presenting enhanced reliability of segmented coatings deposited by APS. Dense interiors of columns or segments cause building-up of residual stress which facilitates cracking. Presence of very fine porosity and narrow intersplat cracks is not recommended since they seem to lead to the formation of arrays of nano-porosity (Fig. 12c) [52]. According to the results of this study, in order to provide good resistance to thermal cycles, it is necessary to control the evolution of the mechanical properties of the YSZ coatings. The results suggest that nano-porosity present in heat-treated the SPS coatings may cause a decrease of toughness. In such case failure of the TBC depends primary on embrittlement of the coating- cracks start and propagate inside the YSZ layer. The presented results support recommendations from [66]- a columnar morphology with high column density (number of column per length) and high porosity ensures good resistance to thermal cycles. Such microstructure provides high compliance of the coating and moderate level of thermal stresses. Porosity inside individual column should have size around micrometer because during high-temperature exposure such pores join together and become spherical (Fig. 12d). Results of mechanical tests show that such change of microstructure may be linked with stiffening and toughening of the coating. In such a case, the thick TGO layer formed after a considerable number of thermal cycles should become the primary path of crack propagation leading to failure.

U( ) Gc ( )

For each individual specimen this parameter was calculated using the general coatings' mechanical properties common to each type but with individual thickness and temperature variation for each sample in one cycle since there were small temperature differences between specimen holders, which did not exceeded 20 °C. One would expect that the higher the value of ζ is, the shorter the lifespan of the TBC will be. Fig. 11a presents the relation between the number of cycles to failure for all the tested coatings and ζ parameter in the as-sprayed condition. Despite of the narrow range of ζ values the coatings showed a wide range of lifespan. For instance, A1 and C2-50E coatings possessed similar ζ around 2.5 × 10−2, but the lifetime of A1 coating was 4 to5 times longer. As already discussed, high-temperature exposure significantly changed the top coatings' mechanical properties, which in turn change ζ parameter. Fig. 11b shows the ζ – lifespan relation for coatings heattreated for 50 h. Long-living coatings (A1 and C1-70E) moved slightly left on the plot due to counteracting increasing toughness. On the other hand, C1-50E, C2-50E, and C2-70 W moved right toward higher ζ values. This shift is mainly due to significant embrittlement of the coatings after heat treatment at 1150 °C. The plots after heat treatment show a broad tendency toward shorter lifespan with increasing ζ. The plots presented in Fig. 11 suggest that premature failure of the most SPS coatings was caused by the change of their mechanical properties, especially the decrease of the ceramic toughness. The mechanisms behind the aforementioned evolution are, at the moment, unknown and will be closely investigated together with the changes in coatings' microstructure. A similar approach was recently used by D.Zhou et al. [50] who looked at the relationship between mechanical properties of YSZ coatings deposited by Axial SPS. The authors highlighted a role of the top coat's fracture toughness and stress intensity factor arising during thermal cycling in the lifespan of the coating. Since constrained thermal strain Δε induces stress in the coating, stress intensity factor K arises at flaws (pores, microcracks) with the size of c:

K

Y

E

5. Conclusions In this study, several SPS coatings were sprayed with various spraying conditions (spraying distance, plasma, output power, and by using two types of suspensions). Most of the SPS coatings failed after a smaller numbers of cycles than the benchmark APS coating. The majority of SPS coatings show short lifespan in spite of apparently columnar or segmented microstructure. Only one SPS coating – C1-70E – performed comparatively with the APS coating. Link between thermal cycle performance and mechanical properties of SPS coatings was sought in order to understand the obtained results. The main conclusions are listed below:

c

where Y is a geometry factor. They considered a change of Young's modulus E during heat treatment by taking into account the average value measured during the test. Over thermal cycling, the crack consistently grows until the critical stress concentration factor KIC, when crack becomes unstable and lead to failure. Lifetime tf was found to be proportional to:

tf

C

Y

KIC E c

n

1 Investigated SPS coatings (except C1-70E) showed sintering-inducted evolution of mechanical properties. Most importantly, high

where C and n are constants characterizing the crack growth process. 174

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initial toughness decreased significantly at 1150 °C and led to the socalled white failure of the SPS coatings (crack propagation in the ceramic coating). These coatings failed after small numbers of thermal cycles. 2 The high magnitude of compressive residual stress in SPS coatings led to the formation of horizontal cracks starting from asperities of bond coat/top coat interface. 3 The reference APS and one SPS coating (C1-70E) underwent toughening and stiffening after high-temperature exposure. This change is well recognized and frequently reported in the TBC literature. These two coatings showed good reliability level in thermal cycle test. 4 Thermal cycle lifespan of SPS coatings does not depend on the initial mechanical properties because they tend to undergo more significant change at elevated temperature.

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