Design of high lifetime suspension plasma sprayed thermal barrier coatings

Journal Pre-proof Design of High Lifetime Suspension Plasma Sprayed Thermal Barrier Coatings M. Gupta, X.-H. Li, N. Markocsan, B. Kjellman

PII:

S0955-2219(19)30743-5

DOI:

https://doi.org/10.1016/j.jeurceramsoc.2019.10.061

Reference:

JECS 12823

To appear in:

Journal of the European Ceramic Society

Received Date:

16 July 2019

Revised Date:

29 October 2019

Accepted Date:

30 October 2019

Please cite this article as: Gupta M, Li X-H, Markocsan N, Kjellman B, Design of High Lifetime Suspension Plasma Sprayed Thermal Barrier Coatings, Journal of the European Ceramic Society (2019), doi: https://doi.org/10.1016/j.jeurceramsoc.2019.10.061

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Design of High Lifetime Suspension Plasma Sprayed Thermal Barrier Coatings M. Gupta1, X.-H. Li2, N. Markocsan1, and B. Kjellman3 1

University West, Trollhättan, Sweden

2

Siemens Turbomachinery AB, Finspång, Sweden

3

GKN Aerospace, Trollhättan, Sweden

Abstract

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Thermal barrier coatings (TBCs) fabricated by suspension plasma spraying (SPS) have shown improved performance due to their low thermal conductivity and high durability along with relatively low production cost. Improvements in SPS TBCs that could further enhance their

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lifetime would lead to their widespread industrialisation. The objective of this study was to design a SPS TBC system with optimised topcoat microstructure and topcoat-bondcoat

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interface, combined with appropriate bondcoat microstructure and chemistry, which could

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exhibit high cyclic lifetime. Bondcoat deposition processes investigated in this study were high velocity air fuel (HVAF) spraying, high velocity oxy fuel spraying, vacuum plasma spraying, and diffusion process. Topcoat microstructure with high column density along with

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smooth topcoat-bondcoat interface and oxidation resistant bondcoat was shown as a favourable design for significant improvements in the lifetime of SPS TBCs. HVAF sprayed

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bondcoat treated by shot peening and grit blasting was shown to create this favourable design.

Keywords

Thermal barrier coatings; Suspension plasma spraying; Lifetime; Topcoat-bondcoat interface; Columnar microstructure.

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1. Introduction Thermal barrier coatings (TBCs) have been used in gas turbine engines since the 1970s to improve engine efficiency by allowing higher gas temperatures or reduced cooling airflow, and/or increased lifetime of turbine components by decreasing base metal temperatures [1-2]. Demands for reduction in engine emissions due to environmental restrictions as well as costs due to increased competitiveness in the gas turbine industry require significant improvements in the performance, durability and production costs of the TBC system.

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A TBC consists of an intermetallic bondcoat followed with by a ceramic topcoat deposited on the metallic gas turbine component as substrate. In a typical industrial gas turbine, topcoats deposited by atmospheric plasma sprayed (APS) are used mainly in stationary components

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(such as combustion chambers and guide vanes) while electron beam – physical vapour

deposition (EB-PVD) topcoats are used in rotating components with high mechanical loads

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(such as turbine blades). This is because columnar microstructures produced by EB-PVD

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yield better strain tolerance and show higher durability during operating conditions [3]. However, high equipment and process costs, and low deposition rates are major drawbacks of EB-PVD process why it is used only for restricted applications and limited suppliers are

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available worldwide [4]. An additional disadvantage with EB-PVD topcoats is the lower thermal insulation capability due to the dense columns with low porosity and wide column

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gaps providing hot gas conduction paths [5]. In APS, feedstock powder particles are injected

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into a plasma jet by a carrier gas and spread over the metal substrate as splats. While APS topcoats are much cheaper to produce and provide better thermal insulation than EB-PVD topcoats due to the lamellar microstructure resulting from the stacking of successive splats, they show lower strain tolerance and durability. TBCs produced by suspension plasma spraying (SPS) have been shown to produce a porous columnar microstructure that could combine the advantages of APS and EB-PVD coatings,

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that is low thermal conductivity and high durability along with relatively low production cost [5-8]. SPS is a considerably cheaper technique than EB-PVD to produce columnar TBCs, both in terms of initial investment in equipment cost as well as the running cost due to its capability to coat large parts without the need of a confined chamber with an inert environment [4]. Lower costs with SPS imply that SPS can not only be used as a replacement of EB-PVD but it could also be used for a variety of components in gas turbines including components where EB-PVD was not employed earlier due to too high costs. Additionally,

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most of the existing infrastructure for APS can be utilised making it a potentially widely available technology. However, further improvements in performance and lifetime of SPS TBCs are still necessary for its widespread industrialisation.

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Several factors affect TBC lifetime interactively such as (i) topcoat microstructure and

chemistry that govern topcoat strain tolerance and sintering resistance, (ii) topcoat-bondcoat

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interface roughness that governs affects topcoat adhesion, stresses created at the interface and

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oxide growth, and (iii) bondcoat microstructure and chemistry that influence govern oxide growth [9-13].

A favourable SPS topcoat microstructure could be a columnar microstructure as it has been

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shown to enhance TBC lifetime due to improved strain tolerance [8, 14, 15]. While SPS process is able to create columnar microstructures similar to EB-PVD TBCs, the column

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density and microstructure is generally significantly different. SPS TBCs generally show

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much lower column density as well as higher porosity in the columns as compared to EBPVD TBCs. Although higher porosity in the columns could be beneficial for lowering thermal conductivity, increase in column density could be beneficial in improving the strain tolerance of SPS TBCs, and thus lifetime. In earlier works, it has been shown that a SPS columnar microstructure with high column density can result in higher TBC lifetime than a columnar microstructure with low column density [15-17], suggesting that a columnar topcoat

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microstructure with high column density could be a favourable design to enhance SPS TBC lifetime. It should be noted here that in case of SPS TBCs, the columnar growth i.e. column density and porosity is affected by not only the topcoat deposition parameters, but also the bondcoat surface roughness which in turn affects TBC lifetime [16]. Therefore, optimisation of SPS topcoat and topcoat-bondcoat interface roughness needs to be done in parallel in order to achieve a favourable SPS topcoat design. In previous work done by the authors, topcoat-bondcoat interface roughness was altered

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varied by varying only the size distribution of the bondcoat feedstock powder, while keeping the same topcoat microstructure and chemistry as well as bondcoat chemistry [18]. It was shown in the study that in case of same bondcoat chemistry and similar bondcoat

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microstructures, smoother bondcoats enhanced the lifetime in case of SPS TBCs, while

rougher bondcoats led to lower lifetime possibly due to higher stresses in the topcoat and

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higher oxidation rate due to formation of oxide pegs [18]. This phenomenon was also

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observed by Bernard et al. who showed that smooth interface topography of diffusion bondcoats could improve result in better topcoat adhesion and high lifetime of in SPS TBCs resulting in lifetimes comparable to EB-PVD TBCs [14]. Based on the work done on APS

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TBCs, a smoother bondcoat is understood to reduce the induced stresses in the topcoat as compared to a rougher topcoat-bondcoat interface [11, 12]. These results suggest that a

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smooth topcoat-bondcoat interface could be a favourable design to enhance SPS TBC

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lifetime. While the as-sprayed interface topography is dependent on bondcoat fabrication process and powder size distribution, it can be modified through surface modification processes such as grit blasting and shot peening which will be investigated in this work. It is not fully understood yet what would be the preferred bondcoat microstructure and deposition process required for enhancing the SPS TBC lifetime. High velocity air fuel (HVAF) spraying has been shown to create dense microstructures with low internal oxidation

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due to the low spray temperature and high spray velocity of this process [7, 19, 20]. In previous work done by the authors, influence of varying bondcoat deposition process on lifetime of SPS TBCs was investigated [17]. Since the topcoat-bondcoat interface roughness also varied among the studied different bondcoats apart from the different bondcoat microstructure, the conclusions could not be directly correlated to the bondcoat deposition process as the resulting lifetime was influenced by a combination of differences in topcoat microstructure, interface roughness as well as bondcoat microstructure and chemistry [17].

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Regardless, HVAF spraying was shown to be a promising candidate for bondcoat deposition, which needs to be investigated further.

Vacuum heat treatment of bondcoat prior to topcoat deposition can promote protective

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alumina growth at the interface and improve TBC lifetime. Keller et al. studied the effect of

heat treatment in vacuum and surface roughness on low pressure plasma spraying (LPPS) and

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high velocity oxy fuel (HVOF) sprayed bondcoats in APS TBCs and concluded that lower

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pressure during heat treatment, low amount of yttrium in bondcoat, and lower number of surface protrusions are favourable for lifetime of TBCs [21]. The effect of vacuum heat treatment of HVAF bondcoats in SPS TBCs will be also investigated in this study.

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The goal of this study was to design a TBC system with appropriate bondcoat microstructure and chemistry, and optimised interface characteristics and topcoat microstructural features. In

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order to achieve this goal, the study was comprised of two parts. The first part was to

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investigate the influence of topcoat-bondcoat interface roughness, varied artificially by shot peening and grit blasting in combination with vacuum heat treatment of bondcoat, on topcoat microstructure and lifetime of SPS TBCs. This part was focussed on HVAF bondcoats, as they have been shown to produce favourable microstructures for high TBC lifetime [7, 17, 18]. The second part was to investigate the influence of bondcoat microstructure, deposited created by different spray processes but with similar topcoat-bondcoat interface roughness

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created artificially by shot peening and grit blasting, on lifetime of SPS TBCs. Bondcoat deposition processes investigated in this part were HVAF spraying, HVOF spraying, vacuum plasma spraying (VPS), and diffusion process. The purpose of performing the study in two parts was to examine study the effects of topcoat-bondcoat interface characteristics and bondcoat microstructure on topcoat microstructure and lifetime of SPS TBCs individually.

2. Experimental

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2.1. Sample preparation In this study, eight sets of samples were produced by varying bondcoat deposition process as well as bondcoat treatment. The samples analysed in this study are summarised in Table 1.

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Shot peening was performed with the main intention of creating mainly to create a smoother bondcoat surface topography. Light grit blasting was performed with the intention of creating

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to produce a favourable roughness for better adhesion of the topcoat as well as for cleaning

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the shot peening residues from the bondcoat layer. Vacuum heat treatment was performed with the purpose intention of promoting interdiffusion as well as controlled oxidation of the bondcoat surface.

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2.1.1. Bondcoat deposition

HVAF spraying was performed with M3 supersonic spray gun (Uniquecoat Technologies,

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USA) using identical spray parameters for all HVAF samples. HVOF spraying was performed

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with the Diamond Jet 2600 gas fuel gun (Oerlikon Metco, Switzerland). The VPS and diffusion bondcoats were ordered from commercial suppliers. Thickness of thermal sprayed bondcoats was around 180-200 µm while diffusion bondcoat was around 80 µm used as standard by the supplier. The diffusion bondcoats analysed in this study are typically used as a standard for EB-PVD TBCs.

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The feedstock powder used for HVAF spraying and HVOF spraying was AMDRY 386 (Oerlikon Metco), while AMDRY 386-2 (Oerlikon Metco) was used for VPS bondcoats as its powder size distribution is more suitable for VPS. It should be noted here that AMDRY 386 and AMDRY 386-2 have the same powder chemistry. The substrate material for HVAF and HVOF samples was Hastelloy-X, while for VPS and PtAl bondcoats, the substrate material was Inconel 792 used as standard by the suppliers. Button shaped substrates in dimensions 25.4 mm diameter x 5 mm thickness were used for all

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tests. 2.1.2. Bondcoat treatment

After bondcoat spraying, the samples were exposed to various treatments as summarised in

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Table 1. Shot peening was performed using cast steel shots while light grit blasting was

performed using alumina grit. The parameters for shot peening and grit blasting were selected

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based on prior experience with these processes. Vacuum heat treatment was performed in a

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furnace at 1120 ºC for two hours used as a standard for VPS bondcoats in industrial applications. It should be mentioned here that VPS and PtAl were already exposed to same vacuum heat treatment in as-received condition as per the standards by the supplier.

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2.1.3. Topcoat deposition

The topcoats in this study were sprayed by SPS using the Mettech Axial III gun with

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NanoFeed 350 suspension feeder (Nortwest Mettech Corp., Canada) using identical spray

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parameters for all samples. 8 wt.% yttria stabilised zirconia (YSZ) suspension in alcohol with 25 wt.% solid load and D50 ≈ 500 nm was used as feedstock material for topcoat spraying. The samples were preheated just before topcoat spraying by performing five strokes of the plasma gun without feeding the suspension. The preheating step was performed as it is understood to improve the adhesion of the topcoat. The samples were continuously cooled

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during topcoat deposition by forced air cooling to maintain the substrate temperature around 200-300 ºC. Thickness of the sprayed topcoats was around 300 µm.

2.2. Bondcoat surface topography Surface topography of samples with only bondcoat in as-sprayed condition as well as after treatments before spraying topcoat were analysed by a stylus-based surface profilometer, Surftest SJ-301 (Mitutoyo Europe GmbH, Germany) following the ISO 4288 standard. Ten

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measurements were carried out by the profilometer on each sample to measure the Ra roughness value.

Three-dimensional (3D) topography of the samples was analysed by white light

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interferometry with the Profilm3D equipment (Filmetrics Europe GmbH, Germany) following the ISO 25178 standard. The lateral resolution obtained with this equipment was around 0.5

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µm while the vertical resolution was 1 nm. The scan area for each measurement was 400 µm

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x 340 µm. The data was captured using the integrated Profilm software, and then exported to the software MountainsMAP (DigitalSurf, France) to post-process the data. Three measurements were performed on each sample in as-sprayed condition as well as after

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treatments before spraying topcoat.

The surface topography was also analysed qualitatively by taking top-view images of the

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samples in backscattered electron mode with the TM3000 tabletop scanning electron

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microscopy (SEM) equipment (Hitachi High-Technologies Europe GmbH, Germany).

2.3. Microstructure characterisation

After spraying the topcoats, the samples were first cold mounted in a low viscosity resin to infiltrate all pores in the coating, then sectioned using a cutting disc, and mounted again in a high viscosity resin to prepare for grinding and polishing. A Buehler AutoMet 300 Pro

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equipment was utilised for grinding and polishing the samples using a semi-automatic routine. The samples after lifetime testing and heat treatment were prepared in the same fashion. Microstructure cross-section was analysed by SEM. Porosity of the topcoats was measured by image analysis technique using the open source software Fiji (ImageJ). Standard procedure for measuring porosity of SPS TBCs developed by the research group in previous works was followed [6]. Images at low magnification (x600) were taken to capture the complete topcoat layer with large microstructural features in the

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coatings such as intercolumnar gaps, large cracks and micrometric pores, while images at high magnification (x4000) were taken to capture the fine intracolumnar porosity. Ten images were taken at each magnification at random locations across the coating cross-section. It should be

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noted here that the images at higher magnification were taken to cover a random area inside the columns and avoid the intracolumnar porosity, as the large porosity features in

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intracolumnar regions are supposed to be covered already in the low magnification images.

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The captured images were first converted into binary images using the image analysis software. Subsequently, the low magnification binary images were processed to remove porous features smaller than 2 µm2 and high magnification binary images were processed to

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remove porous features larger than 2 µm2. This step was performed to ensure that the same porous features were not counted twice. The average coarse porosity measured from low

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magnification images was added to the average fine porosity measured from high

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magnification images to calculate the total average porosity. The density of columns in the topcoat layer was measured by manual counting using the SEM images taken at 200x magnification. A line was drawn over the SEM image using Microsoft PowerPoint software and the number of intercolumnar gaps intercepting this line were calculated for each images. Five such images were analysed for each sample. The column

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density for each image was calculated by dividing the number of times the intercolumnar gaps intercepted the drawn line by the length of the line.

2.4. Thermal cyclic fatigue testing Lifetime of the TBC samples was evaluated by thermal cyclic fatigue (TCF) testing. TCF testing is designed to evaluate TBC performance under long exposures at high temperatures similar to operating conditions in an industrial gas turbine. This test is used as a screening

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method by the industrial partners in this study. As a consequence of the long exposures at high temperatures, the samples after TCF testing show significant bondcoat oxidation and thermally grown oxide (TGO), apart from sintering of topcoat. Failure in TCF testing is

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typically governed by swelling of the TGO and stresses generated due to different coefficients of thermal expansion of different layers in the TBC system during cyclic conditions. Sintering

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also play a significant role in failure.

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and thermal fatigue of the topcoat layer caused due to thermal cycling at high temperatures

During TCF testing, the samples were heated in a furnace (Entech Energiteknik AB, Sweden) at 1100 ºC for one hour, and then cooled down to around 100 ºC in ten minutes by blowing

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compressed air on the surface of the samples. The heating and cooling of samples was repeated in this manner until 20% spallation of the topcoat was observed, indicating failure. A

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photograph of the samples was captured after each cycle by a camera mounted in the cooling

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chamber to determine the exact number of cycles to failure. At least three samples from each set of TBCs were tested in this study.

2.5. Heat treatment While oxidation of the bondcoat can be observed after TCF testing, the samples can be analysed only at the end of testing which means that typically the samples would have been

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exposed to different exposure times. Therefore, in order to compare and analyse TGO growth versus a certain time among different samples, heat treatment was performed using the same equipment as used for TCF testing. The samples were cycled in the same manner as during TCF testing, but were taken out from the furnace after 25, 100, and 500 cycles. One sample from each of set of TBCs was tested for each of the heat treatment conditions.

2.6. Bondcoat hardness

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Hardness of the bondcoat in as-sprayed condition as well as after heat treatment was measured with HMV-2T Vickers micro hardness tester (Shimadzu Corporation, Japan)

following the ASTM E384 standard. During hardness testing, a load of 200 g was applied for

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15 s to calculate HV0.2 values. The samples mounted in epoxy for microstructure cross-

section analysis were used for hardness measurements. At least ten indents distributed across

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the length of the coating cross-section were made for each sample. In case of as-sprayed samples, the indents were made close to the middle of the bondcoat layer along its thickness,

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while in case of heat treated samples, the indents were made close to the middle of the remaining beta phase layer in the bondcoat along its thickness. The indents in the PtAl GB

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3. Results

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sample were made close to the middle of the bondcoat layer along its thickness in all cases.

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3.1. Bondcoat surface topography In this work, surface roughness measurements with profilometer and white light interferometry were combined with SEM top-view images taken in backscattered mode to characterise the bondcoat surface topography. In this combined surface characterisation approach, the 2D profilometer measurements provide mean quantitative data, the SEM topview images indicate quality and composition of the surface, and white light interferometry

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provides visualisation of 3D surface topography in high resolution. Apart from the roughness, 3D surface topography data also provides an indication of the particle deposition conditions (melting and flattening) during coating deposition, which is not possible to determine when measuring surface roughness only by profilometer. This approach has been shown to be highly beneficial in previous work in order to characterise the bondcoat surface topography in detail, both quantitatively and qualitatively, rather than relying only on traditionally employed arithmetic mean height (Ra) measurements [13, 18].

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3.1.1. Profilometer measurements Surface roughness of the bondcoat samples measured by contact profilometer in as-sprayed

condition as well as after different treatments is shown in Figure 1. The HVAF sample with

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bondcoat without any treatment showed a Ra roughness of around 5.5 µm. After shot peening process, the Ra roughness was significantly reduced to around 3.5 µm. Vacuum heat

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treatment of the HVAF bondcoat did not result in a significant difference in surface

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roughness. Light grit blasting after shot peening of the HVAF bondcoat did not result in a significant change in Ra roughness, resulting in Ra roughness of around 3.5 µm. Similarly, for HVOF and VPS bondcoat samples that showed as-sprayed Ra roughness of

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around 6.5 µm and 7 µm respectively, the Ra roughness was significantly reduced after shot peening to around 3.5 and 4 µm respectively, with little effect of light grit blasting on

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resulting Ra roughness. The PtAl bondcoat showed a low as-deposited Ra roughness of

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around 2 µm, again showing little effect of light grit blasting on resulting Ra roughness. These results show that shot peening process was successful in compaction of the surface features resulting in a smooth profile as intended. The prominent result from Figure 1 is that even though the HVAF, HVOF and VPS bondcoats showed different Ra roughness in assprayed condition, all of them showed similar Ra roughness after shot peening as well as after shot peening followed by grit blasting.

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3.1.2. Top-view images Figure 2 shows the SEM top-view images of the bondcoat surface of the samples analysed in this study, while Figure 3 shows the SEM top-view images of the bondcoat surface of HVOF and VPS samples in as-sprayed condition. It should be noted here that all images shown in Figure 2 and Figure 3 were captured at the same magnification. As seen in Figure 2, the HVAF sample showed presence of several large semi-molten particles on the surface as indicated by the arrows. These features are understood to form due

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to the low process temperature during HVAF spraying, resulting in insufficient melting of the coarser larger particles present in the bondcoat powder feedstock thus creating poorly

deformed large particles in the bondcoat. The surface of the HVAF SP bondcoat seemed to be

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quite smooth without any surface protrusions, showing no indications of the semi-molten

particles in the surface as seen in as-sprayed condition. This result implies that shot peening

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process was successful in compaction of the surface features resulting in a smooth profile as

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intended. The surface of the HVAF VHT bondcoat showed a similar topography as the HVAF bondcoat with semi-molten particles, but with a dark grey layer of oxides forming at several areas of the surface. Even though the samples were treated in vacuum, the probable presence

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of low amounts of oxygen during the heat treatment seemed to have resulted in the formation of oxides at the surface, which was confirmed by energy dispersive spectroscopy (EDS).

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These observations correlate well with the Ra roughness measurements shown in Figure 1,

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where the shot peened sample showed much lower Ra roughness than the as-sprayed and vacuum heat treated samples. The images of the HVAF SP+GB sample shown in Figure 2 looked very distinct as compared to the HVAF SP sample, indicating a surface with small-scale roughness exhibiting several small grooves on the surface, as indicated by the solid arrows, created certainly due to the grit blasting process. Several grit residues in the form of black particles can be observed on the

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surface as indicated by the dashed arrows. It seems that while grit blasting was successful in cleaning the shot residues from the surface, it resulted in formation of alumina grit residues. The HVOF as-sprayed bondcoat shown in Figure 3 indicated a similar profile as HVAF bondcoat but with apparently fewer semi-molten particles as indicated by the arrows, which is reasonable considering the higher process temperature during HVOF spraying as compared to HVAF spraying. The VPS bondcoat shown in Figure 3 showed no semi-molten particles on the surface, certainly due to much better melting of the particles in the high temperature VPS

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process. However, after shot peening and grit blasting, both HVOF and VPS bondcoats showed similar characteristics as the HVAF bondcoat, as shown in Figure 2. These

observations correlate well with the measurements shown in Figure 1, where the HVAF,

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HVOF, and VPS samples after shot peening and grit blasting showed similar Ra roughness. The small-scale roughness features created in the PtAl GB sample were non-trivial to

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differentiate due to the different phases present in the bondcoat as can be observed in Figure

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2. However, grit residues in the form of black particles can be observed on the surface similar to HVAF, HVOF and VPS samples after shot peening and grit blasting. 3.1.3. White light interferometry

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Figure 4 shows the 3D surface profiles of the bondcoat surface of the samples captured by white light interferometry. All profiles illustrated shown in Figure 4 are shown at the same

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scale as indicated in the figure.

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The large semi-molten particles in the HVAF sample appear as hemi-spherical hills on the bondcoat surface, as indicated by the arrows. This observation is deduced as the lateral size of the hemi-spherical hills, around 50-60 µm, corresponds well with the size of the semi-molten particles seen in the SEM top-view image in Figure 2, as well as the size of the larger particles in the bondcoat feedstock powder which has a powder size distribution of -63+5 µm. This type of topography has been also observed in previous work done by the authors [18]. After

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shot peening, the HVAF bondcoat resulted in a much smoother surface topography, correlating well with observations in Figure 1 and Figure 2. The formation of oxides observed in Figure 2 in HVAF VHT sample after vacuum heat treatment did not seem to affect the surface topography significantly, showing similar topography as HVAF sample. This result correlates well with the Ra roughness measurements. The HVAF SP+GB sample showed formation of small-scale roughness features in the form of small grooves on the surface created certainly due to the grit blasting process, correlating well

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with the observations in Figure 2. These small-scale roughness features resulted in a substantially very distinct topography as compared to the HVAF SP sample even though both samples showed similar Ra roughness, implying that the small-scale features were not gauged

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in the profilometer measurements perhaps due to being too fine as compared to the stylus diameter. The HVOF SP+GB and VPS SP+GB samples also showed a similar profile as

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HVAF SP+GB sample, correlating well with the results shown in Figure 1 and Figure 2.

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These results infer that after shot peening and grit blasting, HVAF, HVOF and VPS bondcoats resulted in a similar bondcoat surface with similar Ra roughness, surface quality and surface topography, despite exhibiting a clearly distinct as-sprayed surface. The fact that

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samples after shot peening and after shot peening followed by grit blasting showed similar Ra roughness despite showing very distinct topography emphasises the importance of using the

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combined approach for characterising the bondcoat surface instead of trusting only the

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traditionally employed profilometer measurements. The PtAl GB sample, although exhibiting a lower Ra roughness in as-deposited state, showed topographical characteristics similar to HVAF, HVOF and VPS samples after shot peening and grit blasting with similar small-scale roughness features in the form of small grooves on the surface created certainly due to the grit blasting process.

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3.2. Microstructure 3.2.1. Topcoat microstructure Cross-sectional images of the microstructure of all TBC samples analysed in this study are shown in Figure 5, while the top-view images of the topcoats are shown in Figure 6. The topcoat layer in all samples showed a porous columnar structure typically exhibited by SPS coatings. The density of columns in the topcoats measured by manual counting using the SEM images is shown in Figure 7. It can be noted from Figure 5-7 that the shape and density of

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columns varied significantly among the samples despite the fact that identical spray parameters were used for spraying all topcoats.

Similar column shape with irregular column widths and column density of around 18-20

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intercolumnar gaps/mm could be observed in the HVAF, HVAF SP, and HVAF VHT

samples as noted from Figure 5-7. In case of HVAF SP sample, a slightly brighter layer of

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about 20 µm thickness can be observed near the topcoat-bondcoat interface as indicated by

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the solid arrow in Figure 5, suggesting an initial layer with lower porosity and less columnar structure. This effect could be attributed to the flatter bondcoat surface in this sample. Due to the flatter bondcoat surface in this sample, the number of peaks initially available for column

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initiation would be lower than a rougher bondcoat surface, resulting in deposition of predominantly larger particles in the initial stage forming a denser layer [22, 23].

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Subsequently, the initiation peaks required for column formation would have been created by

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the large particles that adhered on the substrate’s surface during the first few layers of deposition [23, 24]. Some of the columns can in fact be noted to be originating from the brighter layer as indicated by the dashed arrows in Figure 5. However, the general shape and density of the columns was similar to HVAF and HVAF VHT samples, probably due to insufficient variation in bondcoat roughness.

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Topcoat layers with narrower columns, much more regular column widths and higher column density of around 25-30 intercolumnar gaps/mm could be observed in the HVAF SP+GB, HVOF SP+GB, and VPS SP+GB samples as noted from Figure 5-7. Despite similar Ra roughness as the shot peened samples, the reason for narrower columns with higher column density in the shot peened and grit blasted samples was deemed to be the small-scale roughness features on the bondcoat surface created due to the grit blasting process. Each of these small-scale roughness features probably acted as initiation peaks required for column

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formation resulting in higher density of columns in these samples. Since these features were quite homogeneously distributed on the bondcoat surface as shown in Figure 4, the width of

the formed columns was much more regular in these samples. Similar topcoat microstructure

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could be observed in the PtAl GB sample as noted from Figure 5-7. This can be attributed to the smoother roughness profile of the PtAl bondcoat [14, 17], as well as the small-scale

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roughness features on the bondcoat surface created due to the grit blasting process. Due to the

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flatter bondcoat surface in this sample, an initial dense layer of topcoat of about 10 µm thickness can be observed similar to HVAF SP sample as indicated by the solid arrow in Figure 5. Almost all columns in the PtAl GB sample seem to be originating from this initial

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dense layer, unlike HVAF SP sample where only some of the columns originated from the initial dense layer. The discontinuity around the middle of the HVOF SP+GB and VPS

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SP+GB topcoats across their thickness that can be observed in Figure 5 was formed due to

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undesired disruption during topcoat spraying of these samples. However, no effect of this discontinuity could be observed on column formation as seen in Figure 5 and Figure 6. The results shown in Figure 5-7 imply that for given spray parameters, column density in the TBC samples is not directly correlated to Ra roughness as may have been deduced based on earlier studies [16, 18]. For example, HVAF and HVAF SP samples showed similar column density despite showing different Ra roughness values, while HVAF SP and HVAF SP+GB

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showed significantly different column density despite showing similar Ra roughness values. The initiation of columns seems to be dependent on not only the large peaks related to the mean surface roughness generally represented by Ra, but also the small peaks related to the small-scale roughness on the surface (created by grit blasting in this case). Porosity of the topcoats measured by image analysis is shown in Figure 8. The coarse porosity values correspond to the porosity evaluated on low magnification images while the fine porosity values correspond to the porosity evaluated on high magnification images. No

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considerable differences among the porosity levels in all topcoats can be observed in Figure 8 when considering the standard deviations. The HVOF SP+GB and VPS SP+GB samples showed slightly higher total porosity values, which could have been a result of the

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discontinuity in the topcoat layer slightly exaggerating the porosity values. Changes in

significantly.

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3.2.2. Bondcoat microstructure

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column density among the samples did not seem to affect the total porosity values

Cross-sectional images of the bondcoat microstructure in all samples analysed in this study after topcoat spraying are shown in Figure 9. The HVAF sample showed a dense bondcoat

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microstructure typically created by HVAF spraying with a rough interface. Flattening of bondcoat surface due to the shot peening process can be clearly seen in the HVAF SP sample.

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Apart from the flattening of bondcoat surface, densification of the upper part of the bondcoat

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layers can be also observed. The densified layer created due to shot peening had a thickness of around 40 µm into the bondcoat layer from the surface. Vacuum heat treatment of the HVAF bondcoat seem to have resulted in even lower porosity as well as homogenisation of the beta phase present in the bondcoat as can be observed in HVAF VHT sample. The HVAF SP+GB bondcoat looked similar to HVAF SP bondcoat, which is reasonable as light grit blasting affected only the bondcoat surface. The HVOF SP+GB bondcoat showed

18

presence of oxides possibly created due to the higher process temperature in HVOF spraying along with some porosity. The denser layer close to the topcoat-bondcoat interface created due to shot peening can be also observed in this sample with a similar thickness of around 40 µm as in case of HVAF shot peened samples, although not as dense as the HVAF shot peened samples. The VPS SP+GB bondcoat showed a dense microstructure with homogeneous beta phase, possibly formed due to the vacuum heat treatment performed by the bondcoat supplier. The denser layer is difficult to distinguish in case of VPS samples probably due to the already

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high density of the bondcoat prior to shot peening. However, regions with incomplete bonding could be observed close to the interface as indicated by the arrows, suggesting incomplete

closure of the bondcoat surface asperities by the shot peening process. Both HVOF SP+GB

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and VPS SP+GB samples showed similar topcoat-bondcoat interface as HVAF SP+GB

samples, which correlates well with the results shown in Figure 1, 2 and 4. The PtAl GB

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bondcoat showed a typical microstructure created by the diffusion process. Minor cracks in

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the bondcoat near the bondcoat surface can be observed, created possibly due to the grit blasting process.

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3.3. Thermal cyclic fatigue testing

The results from TCF testing are shown in Figure 10. Significant variations in TCF lifetime

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among the samples can be clearly observed despite the same bondcoat chemistry (except for

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PtAl bondcoat). All samples showed good lifetime fulfilling the industrial requirements. Shot peening of the HVAF bondcoat resulted in 1.5 times higher lifetime, while vacuum heat treatment did not seem to improve the lifetime. Shot peening of the HVAF bondcoat followed by grit blasting resulted in 1.75 times higher lifetime as compared to the HVAF sample, showing the highest lifetime among all samples analysed in this study. HVOF SP+GB, VPS SP+GB and PtAl GB samples showed 1.25-1.5 times higher lifetime than the HVAF sample.

19

These results show that shot peening of the bondcoat followed by grit blasting can significantly improve the TCF lifetime of the TBC samples. Cross-sectional images of the microstructure of the failed TBC samples are shown in Figure 11. It can be observed that in all cases, failure occurred due to spallation of the topcoat layer due to cracking near the topcoat-bondcoat interface. The spallation occurred as a result of due to a combination of cracking at the topcoat-TGO interface, TGO-bondcoat interface, as well as cracking within the TGO layer. This type of failure is typically observed in TCF testing and

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is understood to occur due to be caused by the stresses created at the topcoat-bondcoat interface owing due to formation and swelling of the TGO layer, as well as mismatch in

coefficients of thermal expansion coefficients among different layers in the TBC system

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during thermal cycling [7, 20].

Opening up of several columnar gaps (indicated by solid arrows) in combination with

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merging of several columns resulting in reduction of columnar gap (indicated by dashed

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arrows) can be observed in all samples in Figure 11. Merging of columns seemed to be more dominant in the samples with narrower columns. These observations can be linked to sintering of the topcoats at high temperatures during TCF testing. It has been observed in

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earlier works that due to sintering of fine particles in the SPS topcoat during TCF testing, intercolumnar porosity is increased while intracolumnar porosity is reduced [17, 25-27].

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Based on earlier works done on APS TBCs, it is understood that for a given exposure

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temperature, sintering of topcoat is mainly dependent on topcoat material and porosity. It has been shown in earlier work on SPS TBCs that when different columnar densities were created due to different bondcoat surface topography using same topcoat material and spray parameters, the sintering rates were similar in all topcoats [17]. In this study, the topcoat material and spray parameters were the same for all samples, resulting in similar topcoat porosities but different column densities created due to different bondcoat surface topography.

20

Therefore, sintering rates are expected to be similar in all samples and thus not considered to play a major role in variation in lifetimes among the samples. When comparing the different bondcoats in the samples after TCF testing shown in Figure 11, all microstructures appeared fairly dense with low porosity content and little internal oxidation, except in the case of HVOF SP+GB sample that showed presence of several oxides as well as porosity within the bondcoat layer. The oxides observed in the HVOF bondcoat were probably a combination of oxides formed during bondcoat deposition and oxides formed

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during TCF testing promoted by the porosity present in the bondcoat before testing. MCrAlY coatings typically consist of two phases: Al-rich beta phase and Al-poor gamma

phase. When the TBC is exposed to service conditions, bondcoat oxidation occurs resulting in

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consumption of the beta phase content and formation of mainly alumina as the TGO layer.

The thickness of remaining beta phase layer in the bondcoat varied substantially among the

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analysed samples as marked in Figure 11. The remaining beta phase content in the MCrAlY

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bondcoat indicates the degree of bondcoat oxidation and the available aluminium reservoir that promotes the formation of the slow growing alumina layer [9, 12]. Complete consumption of beta phase indicates that the protective alumina layer growth is no longer

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favoured, leading to rapid oxidation of other constituents in the bondcoat such as chromium and nickel, which create very high stresses resulting in failure of the TBC [12]. The thickness

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of remaining dark-grey coloured beta phase layer in the bondcoat varied substantially among

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the analysed samples as marked in Figure 11. In case of HVAF and HVAF VHT samples that showed similar lifetime, thickness of the remaining beta phase layer was also similar (around 100-110 µm). In case of HVAF SP and HVAF SP+GB samples that showed higher lifetime, thickness of the remaining beta phase layer was lower (around 80 µm), which is reasonable as these samples were exposed to high temperatures for much longer duration. In case of HVOF SP+GB sample, thickness of the remaining beta phase layer was around 70 µm which was

21

similar to HVAF SP+GB sample. However, the remaining beta phase layer was trickier to distinguish in case of HVOF SP+GB sample probably due to higher local consumption of beta phase that formed oxides within the bondcoat layer, resulting in lower concentration of the remaining beta phase layer. In case of VPS SP+GB sample, the beta phase layer was not visible implying that it was fully depleted. This observation seems reasonable considering the very thick and non-uniform TGO layer that can be observed at the interface in this sample,

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which could have led to more severe beta phase consumption.

4. Discussions

As discussed in section 1, the purpose of this work was to study the effects of topcoat-

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bondcoat interface characteristics varied by different bondcoat treatments and bondcoat

microstructure varied by changing bondcoat deposition process on lifetime of SPS TBCs

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individually.

4.1. Influence of bondcoat treatment

The various treatments performed in this study on HVAF bondcoats resulted in variation of

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topcoat-bondcoat interface topography as well as bondcoat microstructure, even though the starting bondcoat microstructures were identical. These variations can affect not only the

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stresses in the topcoat near the topcoat-bondcoat interface, but also the TGO growth

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behaviour [18, 28-30].

Figure 12 and Figure 13 show the cross-sectional microstructure images of the samples after heat treatment in TCF furnace for 25 and 100 cycles respectively. The microstructure of all HVAF bondcoats after heat treatment looked similar with low porosity and little internal oxidation, with differences appearing when comparing the TGO layers in different samples. It can be observed in Figure 12 and 13 that in the HVAF sample, a dense and dark grey layer of

22

mainly alumina, with quite homogeneous thickness following the interface profile, was formed as TGO. A dense and thin layer of alumina is the most desired form of TGO as it forms a protective barrier for penetration of oxygen into the bondcoat restraining further bondcoat oxidation. The HVAF VHT sample showed a discontinuous TGO layer after 25 cycles as indicated by the arrows. The reason for discontinuity in the TGO layer formation after 25 cycles could have been preferential oxide growth over the oxides formed at the bondcoat surface after the vacuum heat treatment as discussed in section 3.1.2. These pre-

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formed oxides could have acted as nucleation sites for further oxide growth. However, after 100 cycles, this effect was not observed and the TGO layer appeared similar to the HVAF sample.

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It can be observed in Figure 12 and 13 that the HVAF SP sample showed a thin and uniform layer of alumina following the flatter interface profile. Zhang et al. showed for HVAF

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coatings that treatment by shot peening resulted in a TGO layer with lower thickness than assprayed coatings due to suppression of the fast growth of transient alumina by quick

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formation of protective inner α-alumina scale resulting in a much thinner outer α-alumina layer [30]. The reason for this behaviour was that the high surface defect density in the shot

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peened sample assisted the fast formation of a continuous protective α-alumina oxide layer at the initial oxidation stage, unlike the as-sprayed sample where transient alumina oxides were

ur

formed first [30]. Consequently, lesser depletion of beta phase was observed in the shot

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peened samples due to the slower growth of the alumina scale as compared to the as-sprayed samples [30]. This effect was found to be more pronounced in shot peened samples as compared to polished samples [30]. Polishing of bondcoat surface has already been shown to produce a denser and thinner layer of alumina with lesser spinel oxides, due to removal of seed oxides generated during coating deposition required for nucleation of spinel oxides, as compared to a rougher bondcoat surface [28, 31]. These observations correlate well with this

23

study, where the thickness of the TGO layer in the HVAF SP sample appeared to be lower than the HVAF sample, indicating slower TGO growth rate. Similarly, the TGO layer in the HVAF SP+GB sample also appeared to be slightly thinner than the HVAF sample, following the interface profile. Vickers hardness values of the bondcoats in the initial condition and after heat treatment in TCF furnace are shown in Fig. 14. When comparing the HVAF bondcoats in the initial condition, it can be observed that shot peening of the HVAF bondcoat increased the hardness

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of the bondcoat while vacuum heat treatment significantly reduced the hardness of the bondcoat. After 25 cycles of heat treatment in the TCF furnace, the differences in hardness values of the HVAF bondcoats was marginal as the hardness of the as-sprayed and shot

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peened samples was considerably reduced. The hardness values decreased further with increasing heat treatment time, although remaining in the same range, similar to the

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observations in previous work done by the authors [18]. The higher hardness of the shot

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peened samples in the initial condition could be attributed to several factors such as: (i) formation of high compressive residual stresses due to shot peening process, (ii) densification of the bondcoat due to shot peening process, and (iii) possible phase changes occurred during

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shot peening process [32]. The reasons for lower hardness of the HVAF VHT sample in initial condition could have been relaxation of the high compressive residual stresses, formed

ur

initially due to the high particle velocities in HVAF process, and phase changes occurred due

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to the vacuum heat treatment [32]. Apart from these two reasons, further decrease in the hardness values with time could be attributed to the gradual disappearance of the brittle and hard intermetallic beta phase [33]. Based on the results shown in section 3 and discussions in this section, it can be noted that the HVAF VHT sample showed a similar bondcoat microstructure and similar TGO growth as the HVAF sample, especially after longer heat treatment, as well as similar interface topography

24

and topcoat microstructure. These observations correlate well with the similar TCF lifetimes observed in these samples, showing that vacuum heat treatment of HVAF bondcoat is not beneficial for improving TCF lifetime. When comparing the bondcoat surface roughness values shown in Figure 1 with TCF lifetimes shown in Figure 10, it can be clearly observed that for HVAF bondcoats, a smoother bondcoat surface resulted in higher lifetime than a rough bondcoat surface. This observation correlates well with previous works on SPS TBCs [14, 16, 18]. Apart from the smoother

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bondcoat surface created artificially by shot peening process in HVAF SP sample, the reason for its significantly higher TCF lifetime than HVAF sample could have been the lower TGO growth rate as discussed earlier in this section, since the topcoat microstructures in both

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samples were similar. Since the initial effects induced by shot peening on mechanical

properties of the bondcoat, especially residual stresses, seemed to be eliminated after short

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cyclic treatment as shown in Figure 14, they are not deemed to affect the lifetime

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significantly.

The reason for higher TCF lifetime of the HVAF SP+GB sample as compared to HVAF SP sample is believed to be the improved strain tolerance of the topcoat microstructure due to

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higher column density in this sample created as result of the grit blasting process since both samples showed similar interface roughness values, bondcoat microstructure and TGO growth

ur

behaviour. As discussed in section 1, a columnar microstructure with higher column density

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has been shown to result in a higher SPS TBC lifetime [15-17]. It can be concluded from these results that shot peening followed by grit blasting is an effective process to design a favourable topcoat-bondcoat interface.

4.2. Influence of bondcoat deposition process

25

Differences in TGO growth behaviour among samples with bondcoats deposited by different processes can be observed in Figure 12 and 13. It should be emphasised here that HVAF, HVOF and VPS bondcoats had identical starting bondcoat chemistry. While the HVAF SP+GB sample showed a thin and uniform layer of alumina following the interface profile, the TGO in HVOF SP+GB sample appeared to be a bit thicker along with presence of light grey coloured mixed oxides as indicated in Figure 13. The faster oxide growth rate in HVOF SP+GB sample could be attributed to the presence of oxides initially formed close to the

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interface during spraying, and incomplete densification of the bondcoat surface due to shot peening process as can be observed in Figure 9. Additionally, formation of oxides within the HVOF bondcoat can be clearly observed, which would lead to faster consumption of the beta

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phase.

The VPS SP+GB sample showed a non-uniform oxide layer growth in Figure 12 and 13,

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along with formation of oxide pegs forming into the bondcoat as indicated by the arrows. The

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non-uniformity in the TGO layer thickness was exaggerated with longer heat treatment period. This type of oxidation behaviour has been also observed in earlier works and has been attributed to the surface roughness facilitating local depletion of aluminium in the hills at the

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bondcoat surface [17, 27, 34]. Since the bondcoat surface was relatively smoother in this case, the formation of oxide pegs can be attributed to the incomplete closure of the surface

ur

asperities in the initially rough VPS bondcoat, as indicated in Figure 9. These regions close to

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the bondcoat surface could have accelerated oxide growth resulting in a non-uniform TGO layer. The higher oxidation rate in the VPS SP+GB resulted in a wider gap between the interface and remaining beta phase layer than other samples, which could have easily facilitated formation of mixed oxides due to longer diffusion path for aluminium to reach the interface. The PtAl GB sample showed formation of a thin and uniform layer of alumina after heat treatment, as also observed in previous work done on SPS TBCs [17].

26

When comparing the Vickers hardness values of bondcoats deposited by different processes before and after heat treatment in Figure 14, it can be observed that HVAF bondcoat showed the highest hardness in the initial condition, which was reduced after heat treatment as discussed in section 4.1. HVOF SP+GB and VPS SP+GB samples showed a lower hardness in the initial condition despite being exposed to shot peening process, indicating that the initial high hardness in HVAF sample could be mainly attributed to the HVAF process rather than the shot peening process. The lower initial hardness in the HVOF SP+GB sample could

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be attributed to the higher porosity and lower compressive stresses possibly created in this sample by HVOF spraying as compared to HVAF spraying. The lower initial hardness in the VPS SP+GB sample could be attributed to the vacuum heat treatment performed by the

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bondcoat supplier as well as low compressive stresses possibly created in this sample by the

VPS process. Differences in proportions of phase constituents in as-sprayed condition created

re

due to different particle temperatures and velocities during spraying could have also been a

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reason for differences in hardness values of different bondcoats. The PtAl bondcoat showed similar hardness as HVOF and VPS bondcoats in the initial condition, but much higher hardness after 25 cycles, which could be attributed to redistribution of phases in the bondcoat

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due to the heat treatment. The hardness was reduced to similar values as the initial level after longer heat treatment.

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Based on the results shown in section 3 and discussions in this section, it can be noted that

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HVAF SP+GB, HVOF SP+GB and VPS SP+GB samples showed similar topcoat-bondcoat interfaces and topcoat microstructures. Therefore, the differences in lifetime in these samples can be mainly attributed to differences in the bondcoat microstructure created by differences in spray processes and parameters, as the starting bondcoat chemistry was identical in these samples. The reason for lower lifetime of HVOF SP+GB sample than HVAF SP+GB sample could have been the formation of oxides within the bondcoat both during spraying as well as

27

during cycling conditions, with the latter effect enhanced due to the porosity present in the bondcoat in the initial condition. The reason for lower lifetime of VPS SP+GB sample than HVAF SP+GB sample could have been the faster growth of oxide layer leading to faster depletion of the beta phase in the bondcoat. This difference can be also observed in Figure 11 where complete depletion of beta phase was observed in the VPS SP+GB sample at failure. The PtAl GB bondcoat showed similar topographical characteristics as HVAF SP+GB, HVOF SP+GB and VPS SP+GB samples, resulting in a topcoat microstructure similar to

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these samples. The high lifetime of the PtAl GB sample could be attributed to the superior chemical composition of the bondcoat in terms of oxidation resistance, resulting in a thin and homogeneous layer of slowly growing pure alumina. The results from this study show that

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shot peened and grit blasted HVAF NiCoCrAlY bondcoat resulted in higher average TCF

lifetime than PtAl bondcoat. Apart from the significantly lower material cost, NiCoCrAlY

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bondcoat deposited by HVAF results in lower process cost too as compared to the PtAl

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bondcoat. Therefore, it can be concluded from these results that HVAF spraying, combined with shot peening and grit blasting, could be a suitable process to design a favourable

5. Conclusions

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bondcoat microstructure.

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In this study, the influence of bondcoat treatment by vacuum heat treatment, shot peening and

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grit blasting, and influence of bondcoat deposition process on resulting topcoat microstructure and lifetime of SPS TBCs was investigated and discussed. A columnar topcoat microstructure with high column density was shown to improve the lifetime of SPS TBCs. Vacuum heat treatment of the HVAF bondcoat did not result in any improvement in the TBC lifetime, while shot peening and grit blasting processes resulted in

28

significant improvements in the TBC lifetime due to differences in interface topography, TGO growth behaviour and topcoat microstructure. Shot peened and grit blasted HVAF bondcoat demonstrated better oxidation resistance and lifetime than shot peened and grit blasted HVOF and VPS bondcoats investigated in this study, suggesting that HVAF could be a suitable process for bondcoat deposition in SPS TBCs. Shot peening followed by grit blasting was found to be an effective process to design a

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favourable interface due to the following benefits: (i)

smoothening of the bondcoat surface resulting in lower induced stresses in the topcoat,

(ii)

improved strain tolerance of topcoat microstructure with higher column density

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created due to the small-scale roughness features on the bondcoat surface acting as column initiation peaks, and

lower TGO growth rate, especially in HVAF bondcoats, due to quick formation of

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(iii)

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protective α-alumina scale at an early stage.

It can be concluded from this study that the favourable design of SPS TBCs with high lifetime, which is a strain tolerant columnar microstructure with a relatively smooth topcoat-

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bondcoat interface and a bondcoat microstructure showing high oxidation resistance, can be obtained by using a HVAF bondcoat treated by shot peening and grit blasting before spraying

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the SPS topcoat.

Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

29

Acknowledgements The authors would like to acknowledge The Knowledge Foundation (Grant number: 20160022) for the funding of this research work, Nicholas Curry at Treibacher AG for providing the suspension for topcoat spraying, and Olle Widman at Curtiss-Wright Surface Technologies for shot peening of bondcoat samples. Thanks to Stefan Björklund and Rohini Venkatesan at University West for the help with spraying and microstructure analysis

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respectively.

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[31] N. Czech, M. Juez-Lorenzo, V. Kolarik, and W. Stamm, Influence of the surface roughness on the oxide scale formation on MCrAlY coatings studied in situ by high temperature X-ray diffraction, Surf. Coat. Technol. 108-109 (1998) 36-42. [32] J.A. Haynes, M.K. Ferber, W.D. Porter, E.D. Rigney, Mechanical properties and fracture behavior of interfacial alumina scales on plasma-sprayed thermal barrier coatings, Mater. High Temp. 16 (1999) 49-69.

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[33] H. Brodin, and M. Eskner, The influence of oxidation on mechanical and fracture behaviour of an air plasma-sprayed NiCoCrAlY bondcoat, Surf. Coat. Technol. 187 (2004) 113-121.

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[34] P. Song, D. Naumenko, R. Vassen, L. Singheiser, and W.J. Quadakkers, Effect of Oxygen Content in NiCoCrAlY Bondcoat on the Lifetimes of EB-PVD and APS Thermal Barrier Coatings, Surf. Coat. Technol., 2013, 221, p 207-213

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Figure 1. Surface roughness of the bondcoat samples measured by contact profilometer in as-

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sprayed condition as well as after different treatments.

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Figure 2. SEM top-view images of the bondcoat surface of the samples analysed in this study.

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Figure 3. SEM top-view images of the bondcoat surface of HVOF and VPS samples in as-

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deposited condition.

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interferometry.

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Figure 4. 3D surface profiles of the bondcoat surface of the samples captured by white light

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Figure 5. Microstructure cross-section images showing an overview of all TBC samples.

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Figure 6. SEM top-view images of the topcoat surface of the samples.

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Figure 7. Density of columns in the topcoats measured by manual counting using SEM

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images.

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Figure 8. Porosity in the topcoats measured by image analysis.

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ro of -p re

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Figure 9. Microstructure cross-section images showing the bondcoat in all samples.

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Figure 10. Lifetime evaluated by TCF testing.

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Figure 11. Microstructure cross-section images after failure in TCF testing.

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Figure 12. Microstructure images of the samples after heat treatment in TCF furnace for 25

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cycles.

Figure 13. Microstructure images of the samples after heat treatment in TCF furnace for 100

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cycles.

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Figure 14. Vickers hardness values of the bondcoats in the initial condition and after heat

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treatment in TCF furnace.

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Table 1. Samples analysed in this study Bondcoat process

Treatment process

HVAF

HVAF

None

HVAF SP

HVAF

Shot peening

HVAF VHT

HVAF

Vacuum heat treatment

HVAF SP+GB

HVAF

Shot peening + Grit blasting

HVOF SP+GB

HVOF

Shot peening + Grit blasting

VPS SP+GB

VPS

Shot peening + Grit blasting

PtAl GB

Diffusion

Grit blasting

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Sample nomenclature

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