Daniel Scotson
Ahmet Hilmi Paksoy
Ping Xiao- University of Manchester, Department of Materials, Henry Royce Institute, Manchester, United Kingdom
Materials characterisation plays a crucial role in developing thermal barrier coatings and environmental barrier coatings for gas-turbine engines. The failure of thermal barrier coatings and environmental barrier coatings is influenced by a complex interdependence of microstructure, residual stress, and thermomechanical properties. Validating our mechanistic understanding of each of these factors that contribute to failure requires a selection of suitable characterisation techniques. Presented in this review are characterisation techniques, both ex situ and in situ, that have advanced the understanding of thermal barrier coating and environmental barrier coating failure. Targeted coating development that is both effective and efficient depends on these characterisation techniques to obtain superior coatings with improved performance and lifetime.
1 Introduction
Thermal barrier coatings (TBCs) and environmental barrier coatings (EBCs) facilitate using nickel based superalloys and silicon carbide composites respectively in demanding environments, such as the gas-turbine engine. Subjecting materials in a gas-turbine engine to ever more demanding environments is desirable to improve efficiency, accessed by increasing engine core temperatures. As a result, this reduces fuel burn, which simultaneously lowers operating costs and reduces environmental impact. However, limitations in the high-temperature performance (>1,200°C) and steam corrosion of the substrate materials require the introduction of TBCs and EBCs for protection.
The lifetime of TBCs and EBCs is dependent on their many failure mechanisms arising because of their demanding, operating environment. This environment typically involves thermal cycling, steam exposure and ingested contaminants, such as calcium magnesium alumino-silicates (CMAS), that can also result in solid particle erosion and foreign object damage (FOD) (Spitsberg and Steibel, 2004; Lee, 2018; Padture, 2019). Commonly, these conditions occur simultaneously, posing challenges to accurately model and predict the evolution of coating properties throughout lifetime.
Understanding TBC and EBC failure and consequently developing coatings with longer lifetimes is therefore only possible with an extensive use of characterisation techniques. Characterisation techniques, both ex situ and in situ, evidence the evolution of microstructure, residual stresses and thermomechanical properties throughout lifetime, Figure 1. This leads to a better understanding of the complex interplay between these factors contributing to coating failure. Understanding their failure then facilitates the targeted development of more effective coatings by tailoring compositions and obtaining better microstructures to extend lifetime.
FIGURE 1. Outlining the role of characterisation techniques to understand the failure of thermal barrier coatings (TBCs) and environmental barrier coatings (EBCs).
2 Purpose of TBCs and EBCs
TBCs are designed to thermally insulate nickel based superalloy components from the gas-turbine engine core's operating temperature. Typical TBC systems consist of two layers: a ceramic top coat (6-8 wt% Y2O3-ZrO2) and a metallic bond coat (NiCoCrAlY, β-(Ni,Pt)Al or Pt-diffused γ/γ'-based NiAl). Commonly resulting in a millimetre scale TBC, the coating must then have a low thermal conductivity to insulate the nickel based superalloy from a combustion environment that is now hotter than the superalloy's melting temperature. Publications highlighting the development of TBC systems can be found elsewhere (Miller, 1997; Evans et al., 2008; Clarke et al., 2012; Ghosh, 2015; Bakan and Vaßen, 2017; Gopal Thakare et al., 2021).
A step change to engine performance requires even higher engine core temperatures, which means replacing nickel based superalloys with silicon carbide (SiC) based composites. SiC composites' advantages include: higher temperature capability (above 1,200°C); higher creep resistance, and lower density (1/3rd of nickel based superalloys) (Lee, 2000; Spitsberg and Steibel, 2004). Consequently, the introduction of SiC composite components results in weight reduction and removes the need for additional cooling infrastructure, which increases the available air for combustion to generate engine thrust (Lee et al., 2021). Therefore, next-generation gas turbine engines can achieve currently unattainable thrust to weight ratios (Lee et al., 2021).
Despite these advantages, SiC in the combustion environment of a gas-turbine engine is susceptible to steam corrosion at an unacceptable rate of 1 μm/h at 1,200°C (Jacobson, 1993; Opila et al., 1999). Therefore, EBCs are required to protect SiC (Eaton et al., 2000; Eaton and Linsey, 2002; Kimmel et al., 2002) and enable even higher efficiencies in the gas-turbine engine (Spitsberg and Steibel, 2004).
A common EBC system, with a sub-millimetre thickness, is composed of a mixed rare-earth silicate (i.e Yb2Si2O7/Yb2SiO5) top coat and a silicon bond coat. Advantages of these silicate based top coats include their stability at high-temperatures and their lack of polymorphs and resulting phase transformations. Their thermal expansion coefficient is also similar to SiC (4.5 x 10−6 ⁰C−1 (Padture, 2019)), which prevented the established TBC system containing zirconia-based materials (10 x 10−6 ⁰C−1 (Hayashi et al., 2005)) protecting SiC. The silicon bond coat improves the adhesion between the coating and substrate (Kane et al., 2021), but the silicon imposes a temperature limit to an EBC, given its melting temperature of ∼1,416°C (Desai, 1986; Zhang et al., 2022). Therefore, this has started the development of alternative EBC systems with different bond coats (Anton et al., 2020; Harder, 2020; Lee et al., 2022; Zhang et al., 2022). Comprehensive reviews of other EBC systems and compositions can be found in separate publications (Tejero-Martin et al., 2021; Paksoy and Xiao, 2023).
3 Failure mechanisms of TBCs and EBCs
The extreme and complex in-service environmental conditions of TBCs and EBCs, which can simultaneously involve thermal cycling, steam oxidation and CMAS interaction, result in various failure mechanisms.
For TBCs, the thermally grown oxide (TGO), formed at temperatures exceeding 700°C at the bond coat/top coat interface (Stott and Wood, 1987; Takahashi et al., 2019), highlights the need for a holistic approach to understand its associated failure. The TBC TGO can contain a variety of phases dependent on the underlying bond coat; these include α-Al2O3 and mixed oxides, potentially consisting of metastable alumina, zirconia and NiAl2O4 spinel phase (Stiger et al., 1999; Chen et al., 2012). Consequently, residual stresses can then arise from TGO growth, and also from the thermal mismatch between the TGO and adjacent coating layers (Evans et al., 2001). These stresses can then result in buckling and spallation of the coating, initiated at imperfections in the TGO (Evans et al., 2001). Identifying the mechanisms that cause these residual stresses is important to understand their interplay with microstructural evolution.
Similarly, oxidation of EBCs forming a SiO2 TGO is an important failure site. Silicon bond coat oxidation occurs with a typical parabolic growth rate at micrometre thicknesses (Deal and Grove, 1965; Jacobson, 1993; Richards et al., 2016) and residual stresses arise from the difference in bulk thermal expansion coefficients between coating layers. These include: silicon (4.1 × 10−6 °C−1 (Desai, 1986; Richards et al., 2016)), one of the most used EBC top coat materials, Yb2Si2O7 (4 × 10−6 °C−1 (Fernandez-Carrion et al., 2013)), and an α-cristobalite, SiO2 TGO (30 × 10−6 °C−1 (Peacor, 1973; Perrotta et al., 1989)). Consequently, multi-layer models estimate thermal mismatch stresses of 4 GPa in the TGO (Richards et al., 2016; Begley and Hutchinson, 2017). However, these models assume planar interfaces, and that there is no stress relief by cracking or creep deformation (Richards et al., 2015; Richards et al., 2016; Begley and Hutchinson, 2017). An added complication for modelling is representing the complex, inhomogeneous microstructures achieved by thermal spray processes, such as air plasma spray (APS). Therefore, an evaluation of these coatings' thermomechanical properties and residual stresses is required to validate the suitability of these simplified models.
Furthermore, the EBCs' high temperature cubic, β-cristobalite SiO2 TGO also transforms to the tetragonal α-cristobalite phase upon cooling at approximately 220°C (Evans et al., 2001; Richards et al., 2016). The associated 4.5% volume shrinkage upon cooling can initiate cracks to cause coating spallation (Evans et al., 2001; Breneman and Halloran, 2014; Richards et al., 2016). In situ characterisation techniques to evaluate the resulting microstructure changes and residual stresses as a function of temperature are consequently important here.
The morphology of the TBC TGO also affects failure. Repeated thermal cycling can result in a roughening of the TBC bond coat, termed ratcheting, which propagates top coat cracks on each thermal cycle, before eventual crack coalescence and coating spallation (Evans et al., 2001). Therefore, models assuming planar interfaces are unsuitable given the concave and convex interfacial regions more closely resemble a misfitting sphere model (Evans et al., 2001). Furthermore, these models need to incorporate the microstructural evolution to track the change in residual stress distribution along an undulating interface, and the residual stress relief provided by cracking. Consequently, tracking the evolution of microstructure, determining the thermomechanical properties, and evaluating the local distribution of residual stress are all required to fully understand the TGO's role in TBC failure.
As well as undulations at the TGO interface, an increased TGO thickness also increases the likelihood for failure. Considering an energy-based approach for fracture, a thicker TGO increases the strain energy release rate, increasing the driving force for crack propagation and coating delamination (Hutchinson and Suo, 1991; Evans and Hutchinson, 1995; Begley and Hutchinson, 2017). Therefore, a critical TGO thickness is a simple failure criterion, commonly cited for both TBCs (Evans et al., 2001) and EBCs (Richards et al., 2016).
Reducing and controlling the TGO growth rate are therefore approaches pursued for both TBCs and EBCs. These include altering the bond coat morphology with TBCs (Gil et al., 2006; Yanar et al., 2006; Martins et al., 2020) and for EBCs, making top coat additions, such as Al2O3 or Al2O3-containing compounds including mullite and YAG (Y3Al5O12) (Lee, 2018; Lee et al., 2021; Paksoy et al., 2022) Validating these approaches requires quantifying the thermal oxide thickness by evaluating microstructure evolution throughout lifetime from deposition to failure.
However, since coating delamination is typically life-limiting, this involves mixed-mode loading propagating interfacial cracks between dissimilar materials (Piva and Viola, 1980; Liechti and Hanson, 1988; Hutchinson and Suo, 1991). Inducing this failure mode with representative loading boundary conditions is difficult. Hence, the evaluation of the fracture toughness is simplified to individual coating layers, meaning that nanoindentation or micromechanical approaches (Di Maio and Roberts, 2005) are necessary to satisfy the assumptions that the crack size is much smaller than the sample geometry (Irwin, 1948; Irwin, 1957; Lawn, 1993).
Another environmental challenge for TBCs and EBCs is the ingestion of siliceous contaminants into a gas-turbine engine. These can include sand, salt, and volcanic ash, also termed calcium magnesium alumino-silicates (CMAS) (Kim et al., 1993; Poerschke et al., 2017; Cao et al., 2021). The reaction of the coating and CMAS can cause dissolution-reprecipitation of new phases, and at very high temperatures (i.e. >1,100°C), the infiltration of molten CMAS through porosity and grain boundaries (Stott et al., 1994; Borom et al., 1996; Grant et al., 2007; Levi et al., 2012). The infiltration of CMAS can also lead to a loss of strain tolerance, provided by the columnar structure of TBC top coats deposited by electron-beam physical vapour deposition (EB-PVD) (Mercer et al., 2005; Evans and Hutchinson, 2007; Krämer et al., 2008). Understanding the mechanisms of infiltration, phase transformation and microstructural evolution inform the choice of coating compositions that are stable, and also microstructures that provide enhanced CMAS protection (Cao et al., 2021; Morelli et al., 2022).
Furthermore, the physical interaction of ingested contaminants can also result in solid particle erosion and foreign object damage (FOD) of coatings (Wellman and Nicholls, 2004; Bhatt et al., 2008; Gohardani, 2011; Presby et al., 2023). Erosion can also reduce the thermal insulation provided by the coating from a reduction in thickness and thus accelerate TGO growth for failure (Wellman and Nicholls, 2008). This highlights the need for testing in representative environments to establish the links between multiple environmental conditions, for example, between steam oxidation and foreign object damage (Hoffman et al., 2023) that together accelerate coating failure.
In conclusion, a wide variety of failure mechanisms occur in complex multi-layered TBC and EBC systems operating in extreme environments. All these failure mechanisms involve a complex interplay between microstructure, residual stresses, and thermomechanical properties. As a result, it is challenging to obtain reliable and representative models, meaning that characterisation techniques are essential to understand failure and then develop coatings with predictable performance throughout lifetime.
The following sections detail techniques suited to characterising microstructure (Section 4), evaluating stress (Section 5) and also inducing failure (Section 6) in TBCs and EBCs.
4 Characterising microstructure
4.1 Morphology evolution
Characterising a coating's morphology is necessary to study its evolution as a function of its extreme environmental conditions. There are several microstructural features of particular interest for TBCs and EBCs; these include the top coat grain size and grain orientation, and also the top coat's porosity. Additionally, the oxide layer thickness and bond coat roughness play a significant role in the interfacial failure mechanisms related to the TGO.
Furthermore, once failure mechanisms are established, the visualisation of the microstructure is an important feedback loop for optimising coating processing to avoid failure (Whitehouse, 1997). For example, achieving a grain size that reduces steam permeation (Wada et al., 2017; Kitaoka et al., 2018; Matsudaira et al., 2021); optimising the orientation of porosity for strain tolerance (Singh and Wolfe, 2005; Mauer and Vaßen, 2020), or obtaining a suitable interfacial roughness for adherence without increasing the likelihood for failure by ratcheting or rumpling (Evans and Hutchinson, 1984).
The following sections present the application of various characterisation techniques visualising and quantifying microstructural features in the top coat (4.1.1), thermally grown oxide (4.1.2) and bond coat (4.1.3) in TBCs and EBCs.
4.1.1 Top coat
Porosity in the top coats of TBCs and EBCs has both advantages and disadvantages. Porosity present between EB-PVD columns or between APS splats is essential for ensuring stress compliance during extreme thermal cycling conditions (Fleck et al., 2014). However, porous coatings can also increase the rates of CMAS corrosion by providing pathways for diffusion or infiltration (Naraparaju et al., 2017; Tejero-Martin et al., 2022). Furthermore, pores can accelerate the steam corrosion of the Yb2Si2O7 EBC top coat (Mccormack et al., 2023), affect the resistance to foreign object damage (Wellman and Nicholls, 2007; Kedir et al., 2020), and increase solid particle erosion (Presby et al., 2023). Thus, porosity is a parameter influencing TBC and EBC performance, and evaluating its evolution is important to obtain predictable performance (Nakamura et al., 2000; Tejero-Martin et al., 2021).
There are many methods to evaluate porosity within coatings. The Archimedes method is simple (ASTM International, 2013), but cannot quantify a distribution of pore sizes as achieved with mercury intrusion porosimetry (MIP) (Siebert et al., 1999; Medricky et al., 2015). However, both these techniques cannot indicate the spatial distribution of pores. This is important as densification gradients, due to sintering, can lead to elastic modulus variation as a function of depth into a coating (Kulkarni et al., 2006; Fleck et al., 2014). Other techniques include X-ray computed tomography (XCT) to provide this spatial distribution, although its resolution is limited (>1 µm), while SEM offers improved resolution, but the investigated region is limited to the sectioned area.
Characterisation techniques with improved resolution are also required since EB-PVD TBCs contain nanometre scale porosity. Capable techniques include small angle neutron scattering (SANS) and transmission electron microscopy (TEM) (Kulkarni et al., 2006). For example, in situ SANS at temperatures up to 1,100°C and laser flash analysis indicate pore coarsening increases the thermal conductivity after thermal ageing (Saruhan et al., 2011). The development of anisotropic intra-columnar porosity, observed with bright field TEM images after isothermal ageing (over 100 h at 1,100°C), also explains the plateauing of thermal conductivity with thermal ageing of a fully-stabilised EB-PVD YSZ top coat (Saruhan et al., 2011). The limited sample volume of a TEM lamella, means that significant effort is required to reliably investigate any evolution of this porosity.
Measuring the grain size and determining the grain boundary density is another common and straightforward microstructural feature to quantify. As EBC top coat grain boundaries are pathways for both steam permeation (Wada et al., 2017; Kitaoka et al., 2018; Matsudaira et al., 2021) and CMAS infiltration (Sternlicht et al., 2022b), determining the grain size can provide a measure of the EBC's corrosion resistance. Additionally, measuring grain size can evaluate grain coarsening during thermal exposure (Yilmaz et al., 2023), or understand the effect of secondary phases pinning grain boundaries (Paksoy and Xiao, 2021). For this measurement of grain size, electron backscattered diffraction (EBSD) (Humphreys, 2001) or image processing of SEM micrographs are standard methods (ASTM International, 2021b). More traditional methods using intercept methods (Heyn, 1903; Abrams, 1971) are unsuitable given the heterogeneity with thermal spray microstructures.
Grain orientation can also affect thermal conductivity (Zhu and Miller, 2000) and CMAS infiltration (Naraparaju et al., 2017), especially with the highly orientated columnar structure of EB-PVD TBCs. Identifying the intensity variation of X-ray diffraction (XRD) pattern peaks and SEM cross-section micrographs can determine suitable preheating temperatures and deposition times to achieve well-orientated and strain tolerant EB-PVD TBC columnar structures (Czech et al., 1999; Wada et al., 2004).
4.1.2 Thermally grown oxide
Examining a coating's cross-section is a common characterisation approach for assessing the evolution of oxide layers. Microstructures can be imaged with SEM or optical microscopy, with the greyscale contrast variation aiding the identification of oxide layers, Figure 2 (Ridley et al., 2022; Ridley et al., 2023). Capturing high-resolution SEM micrographs of oxide layers with automated sample stages and then stitching these images enables the measurement of thickness variation across significant regions of interest (Ridley et al., 2022, Ridley et al., 2023). For example, cross-sectioning identified that the EBC top coat thickness is not a limiting factor in TGO thickness, suggesting that the diffusion of oxidants across the TGO is the growth rate limiting mechanism (Lee, 2018; Lee et al., 2021; Lee et al., 2021).
FIGURE 2. Evolution of an EBC SiO2 TGO layer with steam cycling. (A) 120 (B) 300 and (C) 500 cycles with each cycle consisting of a 1-h exposure at 1,350°C with a 10-min air quench cool and calculated steam flow rate of 1.5 cm s-1. Blue and pink lines indicate evaluated contours by image thresholding performed by SOFIA software (Stack, 2022). Reprinted from Ridley et al. (2022) with permission from Elsevier.
Methods for segmenting microstructural features, such as oxide layers, are also advancing. Software tools, such as AutoTracer (Su et al., 2021), CMAT (Su et al., 2021) and SOFIA (Stack, 2022), provide a user-interface to segment images and identify the oxide according to greyscale contrast. This automation streamlines the measurement of oxide thickness and roughness, Figure 2 (Ridley et al., 2022; Ridley et al., 2023). Additionally, automatic segmentation with machine learning driven approaches is progressing. For instance, the successful segmentation of EBC oxides has been achieved through transfer learning and MicroNet, a training dataset of over 100,000 microscopy images (Stuckner et al., 2022). Providing open access to these methods leads to greater standardisation and ultimately offers more reliable comparison of TGOs formed in different environments with a variety of coating microstructures and compositions.
4.1.3 Bond coat
Measuring the bond coat's roughness evolution with thermal cycling can be performed with confocal microscopy (Evans et al., 2001). Confocal microscopy reveals the 3D morphology of an exposed surface, such as an oxidised TBC bond coat (Chen et al., 2017; Martins et al., 2020). Correlating between roughness datasets from confocal microscopy and TBC lifetime, enabled the definition of a Total Thresholded Surface Area parameter (Martins et al., 2020). Notably, this parameter reflected the bond coat morphology and showed that bond coats with broad, symmetric peaks resulted in longer lifetimes than bond coats with steep, narrow, asymmetric peaks (Martins et al., 2020).
However, confocal microscopy cannot characterise bond coats or TGOs when a top coat is present. Therefore, for real coating systems, other non-destructive techniques are necessary. An alternative approach is XCT, which has tracked the TBC bond coat roughness evolving with thermal exposure time (Zhao et al., 2012). XCT has also provided insight into the 3D crack geometry relative to the bond coat's peaks and troughs, Figure 3 (Ahmadian et al., 2015). Moreover, XCT identified cracks initiated at off-peak bond coat locations rather than on the peaks (Ahmadian et al., 2015), further highlighting the effect of bond coat morphology influencing failure.
FIGURE 3. X-ray computed tomography (XCT) of an air plasma spray (APS) TBC. (A) Segmentation of cracks showing clusters of cracks (colour) with bond coat/substrate (grey). (B) Contour map of bond coat surface roughness with centroid of cracked region overlaid (centre mark) for geometry shown in (A). Reprinted from Ahmadian et al. (2015) with permission from Elsevier.
4.1.4 Technique overview
Table 1 provides an overview of techniques for characterising the morphology of TBCs and EBCs.
TABLE 1. Overview of techniques characterising the morphology of TBCs and EBCs.
The morphology features of particular interest include the top coat grain size and the TGO interface roughness on representative regions of interest to reliably inform of their evolution and spatial variation. SEM and EBSD, additionally with stitching, provide suitable regions of interest (centimetre-scale) at high resolution (tens of nanometres). However, these techniques are only representative of the prepared slice of material. Confocal microscopy is advantageous for quantifying interfacial roughness of exposed surfaces. Although, this limits any study of bond coat roughness evolution with confocal microscopy as a top coat cannot be present or must be removed.
Alternatively, XCT provides a 3D perspective of the entire coating, with a useful interpretation of porosity. However, the limited resolution (>1 µm) and restricted penetration depth without specialist synchrotron facilities limit XCT's application in characterising ceramic coatings. A representative quantification of porosity at high resolution, while also mapping its spatial distribution is a limiting factor for characterisation techniques with a safer alternative to MIP not currently available.
4.2 Assessing chemistry and phase analysis
The high-temperature environment of TBCs and EBCs means that oxidation, phase transformations and amorphous/crystalline transitions all occur. Affecting these processes are diffusion mechanisms, composition effects and elemental segregation. The characterisation techniques informing their roles in TBC and EBC failure are outlined in this section.
4.2.1 Top coat
With EBCs, crystalline, crack-free top coats are preferred to reduce TGO formation. Typically this involves performing a heat-treatment (above 1,200°C) to crystallise the amorphous phases present after APS deposition, given the high cooling rates (Richards et al., 2015; Garcia et al., 2019; Xiao et al., 2020). Choosing suitable heat-treatment temperatures is important as metastable phases can induce volumetric changes from phase transformations that also affect the porosity and micro-crack density (Garcia et al., 2021; Bakan and Vaßen, 2023).
In situ variable temperature XRD is valuable in identifying such an amorphous to crystalline phase transformation and tracking the EBC top coat's evolution at temperatures of interest, Figure 4 (Garcia et al., 2021). In an investigation of a Yb2Si2O7-based EBC system, the top coat's amorphous to crystalline transformation started at around 1,050°C. Furthermore, the heat treatment temperature determined the crystallised phase; above 1,220°C resulting in the stable β-Yb2Si2O7 phase, rather than metastable phases (Garcia et al., 2021). These metastable phases were also only present with particular processing conditions, Figure 4B, so in situ variable temperature XRD can inform the choice of desirable EBC processing parameters.
FIGURE 4. X-ray diffraction (XRD) patterns as a function of temperature (RT–Room Temperature, 1,000°C, 1,100°C, 1,200°C, 1,300°C) in air of (A) YbDS-3H and (B) YbDS-9H. These are two free-standing Yb2Si2O7 (YbDS) EBC top coat samples, deposited with different APS processing parameters, with further details in Garcia et al. (2019). Reprinted from Garcia et al. (2021) with permission from Elsevier.
For TBCs, the yttria doped zirconia top coats can transform due to the limited composition variation of its phase fields (Scott, 1975). Furthermore, top coats are typically only partially stabilised with yttria to obtain better erosion resistance, which inherently means the phase stability needs understanding through operation (Stecura, 1978; Fang et al., 1999; Medvedovski, 2001; Reed, 2006). The yttria content and hence the ZrO2 phase has a large effect on the thermal cycling lifetime (Stecura, 1985; Miller, 1987). For example, the metastable tetragonal phase or cubic phase yttria-stablised zirconia (YSZ) is preferable to avoid the martensitic transformation and the associated unit cell volume change from tetragonal to monoclinic (Kobayashi et al., 1981; Chevalier et al., 2009). Although, rate-dependent with temperature and composition, a heat treatment around and above 1,300°C can also lead to the as-sprayed metastable tetragonal phase decomposing (Stecura, 1985; Miller, 1987; Brandon and Taylor, 1991).
Therefore, the phase stability of a variety of TBC YSZ compositions has widely been studied with XRD (Brandon and Taylor, 1991; Jones, 1997; Witz et al., 2007). However, XRD can misjudge the cubic phase fraction and also the yttria content in the tetragonal phase (Argyriou and Howard, 1995; Ilavsky and Stalick, 2000). Neutron diffraction can provide a more representative estimation with better peak separation between zirconia's tetragonal and cubic phases, although the limited availability of neutron sources restricts its use (Ilavsky and Stalick, 2000; Ilavsky et al., 2001; Kulkarni et al., 2006).
The phase stability of TBC top coat phases can also be changed by interaction with CMAS. The Y3+ can diffuse from the metastable tetragonal yttria-stabilised zirconia (t'-YSZ), given the higher solubility of Y3+ in CMAS (Nagelberg, 1985). Consequently, this yttrium depleted ZrO2 can transform upon cooling to the monoclinic phase at 600°C, evaluated by in situ variable temperature Raman spectroscopy (Garces et al., 2014). In situ techniques are again invaluable in defining the temperature dependence of phase transformations.
CMAS infiltration and reaction with the top coat on small length scales means that TEM is important. For example, the formation of wetting films at grain boundaries and penetration of CMAS glass into inclusions in contact with grain boundaries are seen in EBCs by TEM techniques, such as scanning transmission electron microscopy (STEM), high-angle annular dark-field (HAADF) and electron energy loss spectroscopy (EELS) (Sternlicht et al., 2022a; Sternlicht et al., 2022b). Complementing this understanding at larger length scales can be XCT, showing a time lapse of CMAS infiltration and the effect of pore morphology (Stolzenburg et al., 2016a; Zhang et al., 2020).
4.2.2 Thermally grown oxide
The formation of a thermally grown oxide (TGO) layer in TBCs and EBCs has seen much attention due to it being a critical failure mechanism. As a result, this has led to many segregation and phase stability studies.
For TBCs, the underlying bond coat affects the TGO phase present. The TGO can contain areas of coarse α-Al2O3 (above γ-phase bond coat) and also areas with mixed oxides, consisting of metastable alumina, zirconia and NiAl2O4 spinel phase (above γ' phase bond coat) (Stiger et al., 1999; Chen et al., 2012; Chen et al., 2022). A pre-oxidation treatment in a low-pressure oxygen environment can suppress the formation of NiO and the Ni(Cr,Al)2O4 spinel phase by forming α-Al2O3, shown with XRD and TEM (Chen et al., 2008a, Chen et al., 2008b; Zhao et al., 2008). However, the depletion of aluminium from α-Al2O3 can result in a mixed oxide and spinel TGO after thermal cycling, identified with XRD (Shillington and Clarke, 1999). The formation of these different phases should be avoided to avoid spallation failure of the TBC (Shillington and Clarke, 1999).
Rumpling of the TBC TGO is a roughening of the compressed oxide to relieve its strain energy (Tolpygo and Clarke, 2000). An increased TGO thickness increases this strain energy, and its growth rate can be controlled by reducing the outward diffusion of aluminium with the introduction of rare-earth elements in the bond coat, including rhenium (Bai et al., 2016), yttrium (Chen et al., 2017) and hafnium (Tolpygo and Clarke, 1998). STEM coupled with energy-dispersive X-ray spectroscopy (EDS) can characterise the TGO grain boundary segregation of rare-earth elements, such as yttrium (Chen et al., 2017). The accompanying smaller lateral growth strain of the TGO then resulted in less rumpling (Chen et al., 2017).
A more quantitative measurement of elemental segregation in the TGO requires atom probe tomography (APT) (Miller and Kenik, 2004). Although a sample is typically only representative of a 100 nm region, APT provides unmatched resolution for 3D mapping of elemental segregation (Miller and Kenik, 2004). Such a technique is therefore suitable for fundamental understanding, although sampling many regions of interest as a function of environmental exposure is unrealistic, given the time-consuming nature of sample preparation and experiments.
Figure 5 shows an example of elemental segregation analysed with APT at an α-Al2O3 grain boundary (Marquis et al., 2013). The concentration of zirconia and yttrium segregating on the α-Al2O3 grain boundary was also shown to increase with proximity to the top coat, suggesting these elements originated from the YSZ top coat (Chen et al., 2022). APT also showed there was preferential segregation towards dislocations (Marquis et al., 2013; Chen et al., 2022). Transmission Kikuchi diffraction (TKD) of needles prepared for atom probe tomography also highlighted whether segregation was related to grain boundary misorientation, although no dependence was noted (Chen et al., 2022).
FIGURE 5. (A) Atom probe tomography (APT) reconstruction from an Al2O3 TBC TGO layer containing a grain boundary and dislocation lines decorated with Y, Zr, Hf and Si. (B) Quantification of atomic concentrations of Zr, Y, Hf and Si at the grain boundary via a concentration profile along the red arrow indicated in (A). Reprinted from Marquis et al. (2013) with permission from Elsevier.
In TBCs, the segregation of sulphur and resulting sulphides reduces the adhesion between the TGO and bond coat due to the different bonding present (Evans et al., 2008). Secondary ion mass spectroscopy (SIMS) and nanoscale secondary ion mass spectroscopy (NanoSIMS) determined the elemental segregation of sulphur at this interface to be up to 10 parts per million, Figure 6 (Bai et al., 2016). These findings highlighted the need for understanding the complete system (i.e. both coating and substrate) given the sulphur was also present in the nickel based superalloy (Lee et al., 1998). A solution to this sulphur segregation was found with the introduction of platinum preferentially gettering the sulphur (Gheno et al., 2010). In addition, glow discharge optical emission spectroscopy (GD-OES) has determined the elemental segregation of sulphur in TBCs given its high depth resolution and capability to detect a variety of elements (Zhao et al., 2017; Lu et al., 2019).
FIGURE 6. Nanoscale secondary ion mass spectrometry (NanoSIMS) mapping of platinum and sulphur on TBC cross-sections with EB-PVD YSZ top coat (Bai et al., 2016). (A) As-received with no heat treatment; (B) after 100 h at 1,150°C in vacuum and (C) after 100 h at 1,150°C in air. From left to right, backscattered electron (BSE) micrograph, and then secondary electron (SE) micrograph with the corresponding platinum and sulphur NanoSIMS maps. The colour scale indicates the counts per second (cps) of 32-S (0–15 cps) and 195-Pt (0–40 cps). Reprinted from Bai et al. (2016) with permission from Elsevier.
Understanding the crystallinity of the SiO2 EBC TGO is also important for failure and steam degradation. SIMs studies show that oxygen diffusivity in the amorphous state is five times greater than crystalline β-cristobalite (Lamkin et al., 1992; Rodríguez-Viejo et al., 1993). A crystalline TGO is then expected to be thinner, although a crystalline TGO is commonly reported to cause spallation failure after thermal cycling, due to the 4.5% volume shrinkage that occurs when β-SiO2 transforms to α-SiO2 upon cooling (Breneman and Halloran, 2014; Richards et al., 2016).
Identifying the crystalline SiO2 EBC TGO can be performed with TEM and Raman spectroscopy. Performing selected area electron diffraction (SAED) provides convincing evidence of α-cristobalite at ambient temperature, Figure 7 (Kane et al., 2021). However, TEM requires the time-consuming extraction of a TEM lamella with a limited region of interest. Alternatively, Raman spectroscopy provides a characteristic spectrum of α-cristobalite, without extensive sample preparation, so a crystalline TGO can be identified at several locations (Richards et al., 2015; K; Kane et al., 2021).
FIGURE 7. Characterising a single layer chemical vapour deposition (CVD) SiC/Yb2Si2O7 specimen after 100 h of 1,425°C steam exposure in an alumina reaction tube. (A) Cross-sectional backscattered electron micrograph with corresponding (B) Al, (C) Yb, (D) Si, (E) O and (F) Mg energy-dispersive X-ray spectroscopy (EDS) maps. (G) High-angle annular dark-field imaging-scanning transmission electron microscopy (HAADF-STEM) characterising the reaction layer above a silica TGO in the EBC. Selected area electron diffraction (SAED) patterns in the reaction layer (H) and silica TGO (I). Reprinted from Kane et al. (2021) with permission from Wiley Materials.
SEM and HAADF-STEM are also effective techniques to investigate the effect of top coat composition on TGO chemistry and morphology. For instance, the introduction of Al2O3 into EBC top coats can improve their steam corrosion resistance (Lee, 2018; Lee et al., 2021; Paksoy et al., 2022). Such an alumina addition to the top coat results in a reaction layer between the TGO and top coat, which is rich in aluminium and ytterbium, Figure 7 (Lee, 2018; K. A; Kane et al., 2021). The presence of aluminium then reduces the TGO growth rate (Lee, 2018; Lee et al., 2021). Furthermore, during steam testing, the formation of an 80–100 nm ytterbium-aluminium-garnet (Yb3Al5O12) outer-layer on the top coat prevented transformation of the Yb2Si2O7 to Yb2SiO5 (Paksoy et al., 2022). The dimensions and phase identification of this layer required TEM studies (Paksoy et al., 2022).
4.2.3 Bond coat
Understanding the TBC bond coat phase transformation of β-(Ni,Pt)Al to γ-(Ni,Pt)Al is important due to the difference in creep deformation. The creep strain rate of β-NiAl is about two orders of magnitude smaller than γ-Ni, so TGO rumpling is more problematic with these β-NiAl bond coats (Noebe et al., 1993; Chen et al., 2017). Furthermore, with the β-(Ni,Pt)Al bond coat, a β to γ phase transformation can occur during thermal cycling due to aluminium depletion (Chen et al., 2017). This β to γ phase transformation occurs above 600°C, evaluated with in situ variable temperature XRD (Chen et al., 2015).
Within EBC bond coats, Raman spectroscopy has identified the presence of boron. Boron is typically present in the substrate and can diffuse during thermal treatment (Costello et al., 1981; Lee et al., 2021). The increased asymmetry in the silicon 520.7 cm-1 vibrational mode, and additional vibrational mode at 620 cm-1 indicate boron is present (Nakano et al., 1992; Knauf et al., 2019; 2020; Rohbeck et al., 2019). These changes to the Raman spectrum suggest boron concentrations above 1019 cm-3 in the silicon bond coat (Dawber and Elliott, 1963). Quantitative light element detection for boron can also be performed with wavelength-dispersive X-ray spectroscopy (WDS) (Lee et al., 2021). Other alternatives include SIMS offering lower boron detection limits (1014 cm-3 (Ochiai et al., 1980)), and also Inductively Coupled Plasma Mass Spectrometry (ICP-MS) and Glow Discharge Mass Spectroscopy (GDMS), which are more regularly used in semiconductor research fields (Fauveau et al., 2016).
Identifying and quantifying boron concentrations are important as boron increases the rate of oxidation in EBCs (Lee et al., 2021; Lee et al., 2022). Furthermore, the solubility limit of boron at temperature increases, which could also contribute to accelerated oxidation at higher temperatures (Knauf et al., 2020).
4.2.4 Technique overview
Table 2 summarises the techniques informing the chemistry and phase analysis of TBCs and EBCs.
TABLE 2. Overview of techniques informing the chemistry and phase analysis of TBCs and EBCs.
XRD techniques are the most widely used for phase analysis, but limitations in gaining only surface-related information, requiring previously indexed crystalline phases, and difficulty in identifying elements in low concentrations have meant alternatives are required. TEM, coupled with STEM EDS, EELS and HAADF, is a leader in determining elemental segregation to grain boundaries, but can be time consuming if the segregation is inhomogeneous and the location is unknown. On the other hand, APT can offer unmatched 3D spatial resolution for a targeted investigation. Both these techniques suffer in that they are only representative of relatively small samples, and also require expertise in sample preparation, data processing and data analysis.
In this regard, spectroscopy techniques offer a faster alternative and can interrogate larger regions of interest. For example, Raman spectroscopy offers a much simpler identification of phases. WDS is advantageous over EDS in quantitative measurement, although NanoSIMS or glow discharge spectroscopy techniques can detect the low concentrations (i.e. parts per million) needed to identify sulphur segregation in TBCs.
5 Evaluating stress
Evaluating residual stress in TBC and EBC systems is not straightforward given there are multiple layers and defects at different length scales. The TGO, bond coat and top coat can contain processing defects (e.g. porosity, splat boundaries, EB-PVD columns) and multiple phases, resulting in residual stresses from thermal mismatch and elastic mismatch. These microstructures are prone to evolve in-service, and environmental changes, such as the penetration of CMAS or temperature, can also lead to residual stress changes. Such a complex residual stress distribution resulting from microstructures at different length scales, means that multiple evaluation techniques are required (Withers and Bhadeshia, 2001).
The following section outlines characterisation techniques that can evaluate coating strains and stresses in an entire coating system (5.1) and those which are limited to particular coating layers (5.2). Furthermore, techniques determining the thermomechanical properties necessary to evaluate the residual stress, such as the elastic modulus, are described in 5.3.
5.1 Entire coating system
A non-destructive evaluation of strain as a function of coating depth with XRD requires a transmission geometry, typically requiring an intense synchrotron light source. For instance, the measurement of residual strains in TBCs as a function of depth has been performed with heat-treated yttria-stabilised top coats (Li et al., 2016). In this experiment, the evaluated compressive strain reached a maximum about 100 µm from the TGO/bond coat interface, which linked with the likely failure location of TBCs (Li et al., 2016). Further to this experiment, residual strains were measured in as-received and heat treated APS TBC samples, again evaluating compressive strains in the top coat and also tensile strains in the NiCoCrAlY bond coat as a function of depth by synchrotron XRD (Li et al., 2019). With thermal gradients also applied, synchrotron experiments can provide the most representative strain measurements of TBCs and EBCs (Siddiqui et al., 2013).
For EBCs, synchrotron experiments have also yielded strains from Debye ring patterns and consequently evaluated residual stresses in EBCs (Harder et al., 2009; Stolzenburg et al., 2011; 2016b). In combination with in situ loading (enabling the determination of elastic constants), and also in situ heating, the residual stress distribution was determined throughout the EBC as a function of temperature (Harder et al., 2009; Stolzenburg et al., 2011). Later experiments, then introduced CMAS to perform an in situ investigation of molten CMAS with wide-angle X-ray scattering (WAXS) (Stolzenburg et al., 2016b).
Alternatively, a commonly used destructive technique for evaluating stress is the curvature method. A coating and a substrate can develop a curvature, arising from a bending moment in the coating and substrate to satisfy equilibrium. A measurement of the beam displacement can then be used to estimate the residual stress in the coating (Godoy et al., 2002; Hayase et al., 2020). The curvature method has correlated the effect of coating thickness with plasma-sprayed coatings and also the effect of heat treatment on these macro-stresses (Godoy et al., 2002; Hayase et al., 2020).
One mechanism commonly cited as causing residual stress is the thermal mismatch between coating layers. The Stoney equation offers a prediction of these macro-stresses with thermal expansion coefficients evaluated using dilatometry (Stoney, 1909; Richards et al., 2016; Begley and Hutchinson, 2017). Measuring the strain as a function of temperature of bulk samples (i.e. pressed and sintered powders, or free-standing coatings), and assuming isotropic expansion, provides a thermal expansion coefficient representative of a bulk sample (Al Nasiri et al., 2016; Kassem and Al Nasiri, 2021; Ridley et al., 2023).
However, the microstructural complexity of APS and EB-PVD deposited coatings leads to an inhomogeneous residual stress distribution within individual coating layers. For example, between different grain orientations, such as with EB-PVD top coats, or between different phases in the EBC top coat (Withers and Bhadeshia, 2001). XRD is suited to evaluate orientation dependent thermal strains of each phase and using Rietveld refinement has aided the evaluation of anisotropic thermal expansion coefficients of both TBCs (Ochrombel et al., 2010) and EBCs (Salanova et al., 2023).
A targeted evaluation of residual stress can be achieved with the micro-ring core method (µRCM) (Korsunsky et al., 2009; Korsunsky et al., 2010; Zhu et al., 2014). For this technique, a focused ion beam (FIB) or laser mills an incremental annular trench while taking SEM or optical micrographs. Performing digital image correlation (DIC) of these micrographs, can then evaluate the displacement field resulting from strain relief and hence determine the residual stress as a function of depth in a coating, Figure 8 (Zhang et al., 2019). The technique is rather intricate, reliant on experimental expertise, and also limited in the investigated region size, so is more likely used to support the findings made by other characterisation techniques.
FIGURE 8. Ring core drilling with focused ion beam (FIB) and optical microscopy to evaluate displacement fields with digital image correlation (DIC) of a heat-treated APS TBC sample. (A) Depth variation of mapped surface displacements. (B) Average surface relief strain evaluated in the pillar's radial direction. Reprinted from Zhang et al. (2019) with permission from Elsevier.
5.2 Specific layer
Evaluating strains or stresses in a particular coating layer can experimentally be favourable, given some techniques rely upon a characteristic vibrational mode or photon emission. Furthermore, many techniques also have limited penetration depths without access to a synchrotron source.
5.2.1 Thermally grown oxide
The thermally grown oxide layer has seen the most interest in residual stress evaluation given the failure of TBCs and EBCs.
Photoluminescence piezo-spectroscopy (PLPS) has extensively evaluated stresses arising in the α-Al2O3 TGO of TBCs (Christensen et al., 1996; Schlichting et al., 2000). The technique is suitable given luminescence (photon emission) arises from the Cr3+ impurity in Al2O3 and the technique is advantageous because the emitted photons can penetrate through the YSZ top coat in TBCs (Schlichting et al., 2000). The resulting spectroscopy pattern, and the peak's position, asymmetry and intensity are affected by stress, composition and defects (Shillington and Clarke, 1999; Nychka and Clarke, 2001; Selcuk and Atkinson, 2002; Zhao and Xiao, 2006; Wang and Atkinson, 2007). Correlating the peak shift and stress (He and Clarke, 1995) has enabled TBC TGO stress evaluation with PLPS, for example, as a function of top coat thickness (Lima et al., 2017), thermal cycling (Selçuk and Atkinson, 2003), and thermal oxidation time, Figure 9 (Christensen et al., 1996). Figure 9 also highlights that a relative comparison of stress evolution as a function of an environmental condition can improve understanding, even with systematic uncertainties in accurately quantifying the stress (Christensen et al., 1996; Luo and Jones, 2010; Chen et al., 2015).
FIGURE 9. (A) Comparison of photoluminescence piezo-spectroscopy (PLPS) profiles from an intact and spalled TBC initially applied on a nickel based superalloy (N5) substrate. (B) Evolution of frequency shift with oxidation time at 1,135°C. Reprinted from Christensen et al. (1996) with permission from AIP Publishing.
5.2.2 Bond coat
TBC bond coat strains can be evaluated with the sin2ψ technique using oxidised bond coat samples. X-rays can penetrate through a TGO, so stress can be evaluated as a function of oxidation time and temperature. An in situ variable temperature sin2ψ experiment with a steam oxidised NiCoCrAlY bond coat showed a notable strain change around 600°C (Chen et al., 2015). This indicated the volume expansion from γ-phase to β-phase upon cooling in the NiCoCrAlY bond coat is an important stress contributing mechanism, in addition to the thermal mismatch between TGO and bond coat (Chen et al., 2015).
Stresses in EBC silicon bond coats have been evaluated with Raman spectroscopy using calibrated stress and Raman shift relationships known for silicon (De Wolf, 1996; Rohbeck et al., 2019; Lance et al., 2023). Performing Raman spectroscopy maps across exposed steam degraded silicon samples, has identified shifting of the Raman spectra and associated this with the volume expansion of the SiO2 α-cristobalite to β-cristobalite phase transformation upon heating (Swainson et al., 2003; Lance et al., 2023).
5.2.3 Top coat
Raman spectroscopy has similarly been used for stress evaluation in the top coat of TBCs. A mapping study showed a correlation between microstructural porosity and Raman peak shift for the 640 cm-1 vibrational mode of the t-ZrO2 phase (Tanaka et al., 2006; Tanaka et al., 2009). Determining local stress is useful to understand how microstructural features affect stress concentrations that may lead to cracking. Stress evaluation studies are lesser performed in the top coats of TBCs and EBCs as the critical failure mechanism involves the bond coat and TGO.
5.3 Thermomechanical performance
Residual stresses are also dependent on thermomechanical properties. The elasticity, plasticity and creep of the coating are all temperature dependent, and these can all affect the residual stress. Furthermore, the common method of evaluating stress (i.e. using a measurement of strain) relies upon a representative evaluation of other parameters, such as the elastic modulus and Poisson's ratio. As these parameters are influenced by the characterisation technique's sampling volume, the interrogated volume of strain should also match to achieve a reliable evaluation of the stress at the desired length scale. The following section describes techniques to evaluate the elastic-plastic behaviour of the coatings (5.3.1) and also coating creep (5.3.2).
5.3.1 Elastic-plastic behaviour
Microindentation and nanoindentation performed on coating cross-sections are the most common techniques for elastic modulus and hardness evaluation (Zhong et al., 2018; Jang et al., 2020). Applying indentations in this geometry can evaluate the variation in these mechanical properties between adjacent layers and within individual layers. When performed on TBC top coats, this showed a hardness and elastic modulus gradient, with the highest values obtained at the coating's surface (Zhu and Miller, 2000). Furthermore, combining indentation with samples exposed to different sintering durations, showed the effectiveness of the technique in linking the microstructural evolution and the increase in elastic modulus (Zhu and Miller, 2000; Gao et al., 2023).
The elastic modulus of APS TBCs and EBCs can also be assessed with three-point bending. This resulted in a lower elastic modulus value when compared with indentation, which is likely given the more local interaction volume for indentation (Chen et al., 2020).
These scalar representations of elastic modulus evaluated by bending or indentation may be suitable for macroscopic stress evaluation. For example, when combined with a measured strain that is also evaluated from a region of interest that encompasses multiple grains, as with XRD. However, bending or indentation techniques are not very effective in determining anisotropic elastic properties.
Furthermore, following Neumann's Principle, the elastic modulus of TBC and EBC top coat materials is usually anisotropic, given their non-cubic crystal structures (Nye, 1957). Therefore, mechanical properties may be desirable on a much smaller length-scale to determine how APS splat orientation or EB-PVD column orientation affects coating compliance. Consequently, microcantilever bending is considered as an alternative to evaluate the elastic modulus to consider orientation effects. Figure 10 shows microcantilevers with in-plane and out-of-plane orientations (either parallel or normal to the bond coat/top coat interface), and these quantified significant anisotropy resulting from the APS splat morphology (Chen et al., 2019).
FIGURE 10. Microcantilever beam bending of APS TBC for anisotropic mechanical property evaluation. (A) Schematic of the different orientations of the microcantilever beams; (B) schematic of testing geometry with nanoindenter to apply load (refer to original publication for description of the labelled parameters); (C) Micrograph of a microcantilever beam. Reprinted from Chen et al. (2019) with permission from Elsevier.
As a function of temperature, the plasticity of coating layers also changes. Fundamental studies on single crystals of α-Al2O3 show dislocation motion on the basal plane is preferential at high temperatures with prismatic slip preferential under 600°C (Lagerlöf et al., 1994; Mitchell and Heuer, 2005). Therefore, as compressive stresses in the TBC TGO increase, stress relief from dislocation slip may occur at higher temperatures in the α-Al2O3 TGO (Chen et al., 2022). The segregation of elements and the temperature dependence of diffusion to these dislocations may also affect dislocation slip (Marquis et al., 2013; Chen et al., 2014).
5.3.2 Creep
Given the high-temperature operation of TBCs and EBCs, creep deformation (i.e. time dependent strain) can relax the stresses at oxidation temperatures (Evans et al., 2001). Evaluating these strains has been performed with DIC to evaluate a stress exponent for EBCs consistent with diffusional creep (Archer et al., 2020). DIC has also been used with optical images to evaluate the surface displacements and surface strains of a NiCoCrAlY bond coat as a function of temperature (Hemker and Thompson, 2007). The high temperatures required for creep deformation (>700°C) mean that speckle patterns are unstable at these temperatures, so the inherent microstructure is used to track deformation, limiting the correlation ability at the highest resolutions.
Displacement fields can also be evaluated using the intensity contrast in XCT tomographs with digital volume correlation (DVC) (Bay et al., 1999). For example, displacements arising due to TGO growth and bond coat creep can be evaluated at a given temperature and as a function of exposure time (Khoshkhou et al., 2016). However, the microstructure's evolution, such as the sintering of TBC EB-PVD columns and TGO growth, can prevent the same feature being tracked for successful correlation. As a result, there is a time constraint between consecutive tomographs for adequate correlation, which typically requires an intense synchrotron source (Khoshkhou et al., 2016). Furthermore, the voxel size (typically above 1 µm (Maire and Withers, 2014)), also restricts the displacement field's spatial resolution. Techniques with in situ mechanical and in situ temperature capability, and also better resolution, as achieved with SEM/EBSD (Donoghue 2022; Smith 2022), can further improve the understanding of TBC and EBC deformation.
5.4 Technique overview
Table 3 highlights the techniques evaluating the stress and thermomechanical properties of TBCs and EBCs. Many of these techniques lack a reliable quantitative evaluation of the stress given the dependence on many assumptions of measured strains and linear elastic fracture mechanics. These limitations typically lead to a reliance on qualitative comparison between samples, which may not provide a comprehensive understanding of what causes the relative changes. For instance, while XRD can provide a global measure of the coating strain, it cannot solely detect the relationship between the microstructure and residual stresses. Raman spectroscopy has started to offer insight here, although it relies upon calibration measurements, and spectra require careful interpretation to ensure the peak shift is solely from residual stress.
TABLE 3. Overview of techniques evaluating the stress and thermomechanical properties of TBCs and EBCs.
Synchrotron facilities are an excellent resource for residual stress evaluation, although their availability limits the number of studies performed. Furthermore, the speed of development of TBC and EBC systems means that laboratory-based characterisation techniques are typically favoured to reduce the development cycle time of these industrially relevant coatings. However, synchrotron investigations yield valuable experimental data, unattainable by laboratory systems, and therefore have an important role in comprehensively characterising residual stresses in TBCs and EBCs. In the future, compact light sources using inverse Compton X-ray sources may offer the advantages of synchrotron light sources with increased availability (Gunther et al., 2020).
6 Inducing failure
As described previously, the failure of coatings in the gas-turbine engine environment is complicated. TBCs and EBCs experience residual stresses, oxidation and environmental degradation. The evolving microstructure of the top coat, bond coat and TGO consequently affect the residual stresses and thermomechanical properties to complicate defining the main cause of failure. Therefore, developing accurate models that predict failure behaviour to assess coating lifetime is difficult (Hutchinson and Hutchinson, 2011).
Consequently, experimental techniques that mechanically load and reproduce the in-service failure modes of the coatings are important investigative tools. For example, these techniques can define locations that limit lifetime, while also quantifying the failure (i.e. through measurement of adhesion strength or fracture toughness). Obtaining accurate values for failure related properties can be challenging. Although a relative ranking, for example, with respect to operating conditions, can be combined with an understanding of the morphology, elemental segregation or residual stress discussed previously to understand coating failure.
Tools for inducing failure typically evaluate the coating's adhesion and cohesion. Adhesion refers to the bonding between layers at an interface, while cohesion relates to the integrity of a specific coating layer and how this resists fracture or separation. Both these concepts are crucial for TBCs and EBCs to ensure the coating remains intact and structurally stable under various operating conditions, such as thermal cycling, steam oxidation, and ingested contaminants.
Given these harsh operating conditions and the lifetime limiting failure mechanisms of TBC and EBCs, typically involving cracks propagating along or adjacent to a TGO interface, tools for investigating adhesive failure, rather than cohesive failure, have seen greater use. Furthermore, for TBCs there has been a focus on room-temperature testing, with thermal mismatch contributing to the largest residual stresses, making ambient temperature testing the most relevant for investigating failure (Hutchinson and Hutchinson, 2011).
6.1 Cohesive failure
Investigating cohesive failure considers the inherent fracture resistance of a particular coating layer. Techniques such as XCT, SEM or TEM can all identify initiating defects and crack paths, and from these observations, failure mechanisms can be inferred. However, characterising these cracks occurs without a measurement of the boundary conditions (i.e. displacement or load), making it difficult for a quantitative comparison by evaluating the fracture toughness between coatings exposed to different environments. Therefore, using specific testing geometries, at least with analytical solutions, is favourable to evaluate the fracture toughness.
Additionally, the anisotropic microstructure adds complexity to a representative evaluation of the fracture toughness. As a result, common techniques such as bulk bending tests (i.e. three-point and four-point bending tests) are unsuitable. Therefore, there's a requirement for evaluation of fracture toughness on a scale appropriate to defect size and coating thickness.
As such, Vickers indentations can initiate well-defined radial crack lengths and evaluate fracture toughness (Anstis et al., 1981). Despite challenges posed by non-homogeneous microstructures, such as obtaining Palmqvist cracks from a Vickers indenter in plasma-sprayed microstructures, several studies on TBCs and EBCs have shown consistency with other characterisation techniques (Zhou et al., 2017; Song et al., 2019; Yuan et al., 2022). For example, Mao et al. (2012) conducted a notable study using a modified indentation technique to evaluate fracture toughness and residual stress in plasma sprayed TBCs (Mao et al., 2012). By analysing crack lengths, indentation loads, and deformation patterns, these mechanical properties were load dependent and influenced by both the microstructure and layer thickness.
Another option for fracture toughness evaluation is microcantilever beam bending, Figure 10 (Di Maio and Roberts, 2005; Chen et al., 2019). This provides a local evaluation of the fracture toughness, although the population of defects within the sampled volume can influence the value obtained. For instance, cantilevers manufactured in TBC APS microstructures without splat boundaries, have significantly higher in-plane fracture toughness values (Chen et al., 2019). The ability to machine cantilevers in different regions can also evaluate the fracture toughness variation as a function of depth within the coating (Jaya B et al., 2012). Therefore, it is critical that fracture toughness is evaluated on different length scales to evaluate the impact of different defects contributing to failure, although this is difficult given the inherently inhomogeneous thermal spray microstructure.
6.2 Adhesive failure
Many techniques can investigate a coating's interfacial adhesion. Techniques such as the tensile pull-out test, bending tests (i.e. three or four-point bending) and transverse scratch testing provide good qualitative, comparative understanding of how an in-service environment affects the coating. Other testing methods such as the shear test are also used for quantitative measurement of interfacial properties.
6.2.1 Tensile pull-off and buckle methods
The pull-off test is one of the most widely used tests for comparing the bonding strength between the bond coat/top coat or bond coat/substrate interfaces (Karaoglanli et al., 2013; Ghasemi and Vakilifard, 2017). It involves applying an adhesive bonding agent to attach both sides of the test specimen (i.e. coating and substrate) to a loading fixture. Consequently, a tensile load is applied normal to the coating layer plane. Due to the complex microstructure and composition of both TBCs and EBCs, and also the strength limits of the adhesive agents, multiple tests are required to obtain a statistically significant measure of the coating adhesion or cohesion. Therefore, it is often used for quality control or a qualitative comparison, rather than an accurate quantification of interfacial strength.
The strain to buckle test is commonly used to assess the bonding strength of ceramic coatings (Zhao et al., 2011). This test involves subjecting the coated substrate to a compressive stress until the coating buckles. However, the strain to buckle test has similar limitations to the pull-off test. For example, only a comparative analysis of similar coatings and it also requires many samples. Moreover, the applicability of this technique to EBCs is limited due to the stiffer EBC SiC substrate compared to the nickel based superalloys used with TBCs.
Given these limitations, laser spallation tests have attracted interest. Named the laser shock adhesion test (LASAT), a laser generates shock waves at the substrate or coating surface, and the resulting vibrations can result in delamination at the coating/substrate interface. Although LASAT is currently a rare method for assessing the interfacial properties of TBC or EBCs, its ability to rank different coatings in terms of adhesion, and to evaluate the evolution of adhesion over time with thermal exposure (such as thermal cycling or aging), suggests it has potential for future adoption (Guipont et al., 2019; Maurel et al., 2019; Mahfouz et al., 2023).
6.2.2 Bending tests
Bending tests are also conducted widely for the interfacial fracture behaviour and the bonding strength of TBCs and EBCs. In these tests, a coating is applied to a substrate, and the coated sample is subjected to a bending load with standard testing equipment. The bonding strength of the coating can then be calculated based on the maximum load that the sample withstands before the coating delaminates or fractures.
Bending tests are typically performed with three-point or four-point geometries. In three-point bending, the maximum stress occurs at the loading point, so a notch is usually manufactured prior to testing, to create a stress concentration at a specific location in the coating thickness, for instance, at the top coat/TGO interface for TBCs (Martins et al., 2021). On the other hand, the four-point bending test involves two loading points and two holding points. This configuration results in a uniform stress between the loading points and results in a larger sample volume under a uniform stress, therefore making the four-point bending geometry more suitable to determining the weakest region of the coating system (Chen et al., 2020).
One example of a four-point bending test that has influenced TBC and EBC development is conducted on the fracture characteristics and damage evolution of coatings with a notch (Zhao et al., 2012; Martins et al., 2021). Using this geometry, the critical strain energy release rate could be linked with the bond coat morphology (previously specified in 4.1.3) to determine preferential bond coat morphologies to resist failure (Martins et al., 2021). Another modified four-point bending geometry can also generate a single interface crack (Zhao et al., 2010). Advantageously, this propagates a crack along the TGO/top coat interface to determine the preferential crack path.
Overall, bending tests have played a significant role in understanding the failure mechanisms and mechanical properties of TBC and EBC systems. They have provided valuable data for modelling the cracking mode, analysing the energy release rate, and developing numerical lifetime models for the durability of these coatings.
6.2.3 Transverse scratch testing
An alternative method evaluating the adhesion strength is transverse scratch testing (Johnstone et al., 1997). Primarily used to investigate APS deposited TBCs, the method involves scratching a polished cross-section with a cone shaped intender under a constant load from the substrate to coating surface. The resulting scratched areas can then evaluate the failure modes (i.e. cohesive or adhesive failure) (Lopez and Beltzung, 1989). While not appropriate for thin EBCs (Al Nasiri et al., 2019; Yilmaz and Xiao, 2022), it is suitable for thicker coatings (i.e. processed using plasma spraying). Studies comparing the technique with the pull-off test suggest it could be useful as a simple and reliable quality control technique (Nohava et al., 2010; Vencl et al., 2011). For example, it has been used to evaluate the effect of heat treatment temperature on TBC bond coat adhesion (Morrell and Rickerby, 1998).
6.2.4 Indentation
Microindentation and nanoindentation techniques can also be used for the analysis of bonding behaviour (Vidakis et al., 2003; Mao et al., 2012). Indentation from the top surface of the coating to evaluate the interfacial toughness is widely used for TBCs (Vasinonta and Beuth, 2001). Cross-sectional indentation at or near the interface has also attracted interest to examine the crack paths and evaluate the interfacial toughness (Lesage and Chicot, 2002; Wang et al., 2012). However, the non-uniform topography of the interfaces resulting from thermal spraying or EB-PVD processing, means that indentation placement is important.
Furthermore, a suitable choice of indenter geometry is required. For example, using indenters with sharp corner geometries (i.e. Vickers and Berkovich) provides an inhomogeneous stress distribution that preferentially initiates cracks near to the corners. Therefore, overcoming this issue is macro-scale spherical indentations near the interface. This generates, large displacements at the interface, with the crack paths examined to determine preferential sites of crack propagation, Figure 11 (Wang et al., 2012).
FIGURE 11. (A) Cross-sectional indentation in nickel based superalloy (CMSX4) substrate near TBC interface. (B) Resulting crack path along the bond coat/TGO interface. Reprinted from Wang et al. (2012) with permission from Elsevier.
An appreciation of the residual stresses is also necessary in relationship to the failure test chosen (Clyne and Gill, 1996). All test geometries have an associated sampling volume, which will affect the contribution of residual stresses, and may alter the propagation of a crack (Araujo et al., 2005). Therefore, caution is required when comparing results between different testing geometries.
6.2.5 Shear test
Replicating the in-service mode of failure during testing is important to obtain representative behaviour. In-service, cracks in multi-layered coatings are subject to mixed mode loading (Suo and Hutchinson, 1990; Hutchinson and Suo, 1991; Begley and Hutchinson, 2017). For example, a crack initiated in mode I (tensile stress normal to the crack) may then propagate into adjacent layers with compressive stress, subjecting these cracks to shear stresses and mode II loading. As a result, the crack changes direction and propagates along an interface which can then lead to coating delamination.
The shear test is a testing geometry with these mode II loading conditions, which lead to delamination failure of TBCs. Although, requiring a specific sample geometry, the shear test has quantified a coating's interfacial adhesion by evaluating the critical strain energy release rate in EB-PVD TBCs (Guo et al., 2005; Guo et al., 2007).
6.3 Technique overview
Table 4 contains an array of techniques for initiating the failure of TBCs and EBCs, and details their advantages and disadvantages. Even though some techniques have been standardised by industrial stakeholders performing TBC and EBC research and development (i.e. pull off test, etc.), obtaining reliable results to understand failure has proven a difficult feat, given the size and geometry of the coatings. Furthermore, the actual failure of TBCs and EBCs often occurs due to residual stresses evolving in environmental conditions for extended periods, such as thermal cycling or high-temperature steam oxidation. Therefore, work is still needed to ensure in-service failure modes are reproduced with failure testing that is effective and efficient.
TABLE 4. Overview of techniques used to initiate the failure of TBCs and EBCs.
7 Conclusion
Characterisation techniques facilitate the understanding of thermal barrier coating and environmental barrier coating failure by revealing the evolution of microstructure, residual stresses and thermomechanical properties throughout their lifetime. All these factors contribute to coating performance; an appreciation of each factor and their interconnections is crucial to successful coating development. This understanding, in combination with techniques to initiate failure, can then inform of desirable microstructures and compositions to improve coating performance and prolong coating lifetime.
A wide variety of characterisation techniques are presented for investigating thermal barrier coating and environmental barrier coating failure. While the understanding of coating performance is complex, it is essential that the advantages and disadvantages of characterisation techniques in respect to thermal barrier coatings and environmental barrier coatings are also understood. This enables the effective choice of a combination of techniques to investigate these coatings on a variety of length scales.
The need for characterising thermal barrier coatings and environmental barrier coatings will only grow into the future. Improving the efficiency and lowering the environmental impact of gas-turbine engines remains desirable, and consequently this requires substrates and coatings to operate in increasingly extreme environments. Therefore, understanding coating failure will continue to rely upon characterisation techniques to further develop high-performance thermal barrier coatings and environmental barrier coatings.
Author contributions
DS: Writing–original draft, Writing–review and editing. AP: Writing–original draft, Writing–review and editing. PX: Funding acquisition, Project administration, Supervision, Writing–review and editing, Conceptualization, Resources.
Funding
The author(s) declare financial support was received for the research, authorship, and/or publication of this article. DS acknowledges funding support from the Engineering and Physical Sciences Research Council (EPSRC) Doctoral Training Partnership (DTP) (grant number EP/R513131/1).
Acknowledgments
PX would like to acknowledge support from the Royal Academy of Engineering and Rolls-Royce for the appointment of Rolls-Royce/Royal Academy of Engineering Research Chair in Advanced Coating Technology.
Conflict of interest
The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.
Publisher’s note
All claims expressed in this article are solely those of the authors and do not necessarily represent those of their affiliated organizations, or those of the publisher, the editors and the reviewers. Any product that may be evaluated in this article, or claim that may be made by its manufacturer, is not guaranteed or endorsed by the publisher.
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Keywords: characterisation techniques, environmental barrier coating, EBC, thermal barrier coating, TBC, failure
Citation: Scotson D, Paksoy AH and Xiao P (2024) Characterisation techniques for investigating TBC and EBC failure: a review. Front. Ceram. 1:1307437. doi: 10.3389/fceic.2023.1307437
Received: 04 October 2023; Accepted: 29 November 2023;
Published: 08 January 2024.
Edited by:
Catalin Croitoru, Transilvania University of Brașov, RomaniaReviewed by:
Ali Talimian, Alexander Dubcek University in Trencin, SlovakiaKen Lee, National Aeronautics and Space Administration (NASA), United States
Copyright © 2024 Scotson, Paksoy and Xiao. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
*Correspondence: Ping Xiao, p.xiao@manchester.ac.uk