Essay: Coatings for increased turbine engine efficiency

Recently used turbines for power generation and aero-jet aircraft engines applications experience incessant efforts towards increasing the thermal efficiency [1]. Greater efficiency and higher output inescapably need to increase the turbine operating temperature, which needs novel methods to overcome the restraints of hot sections materials [2]. Early, before 25 years ago, the increase in the turbines operating temperatures was depended mainly on alloy design to produce alloy with higher creep and oxidation resistance. Internal cooling using intricate channels inside the blades had been used, also, to increase the gas operating temperature [1001, 1002]. For the last 25 years, The application of Thermal barrier Coating (TBC)s on Ni-base superalloys lead to a significant increase in engine operating temperature. TBC systems are the most promising technology to meet the incessant gas turbine development. Fig.000 is showing the maximum gas turbines operating temperatures over the last decades [1003].
Recently, research efforts were focused to develop the ceramic TBC on the parts of the turbines because the traditional coating materials have reached their temperature limits capabilities. TBCs are deposited on the hot-path components of gas turbines like combustion lines, blades vanes and other transition parts either to increase the inlet temperature in order to improve the thermal efficiency or to reduce the cooling system requirements. For instance, it has been reported that an increase by 170°C in the operating temperature of the engine can improve the thrust of the engines by 5 % and enhance the efficiency by 1.0 % [a].TBC has the capability of increasing the limit operating temperature of gas turbines by 90~170°C above what is achievable with the most advanced super alloys [3].
Typically TBC systems consist of a Ni-based superalloy substrate coated with MCrAlY or nickel platinum aluminide Ni (Pt,Al) Bond Coat . Yttria-Stabilized Zirconia (YSZ) is deposited on the bond coat as Top Coat layer. The BC develops a Thermally Grown Oxide layer (TGO), ideally consists of α-alumina, at high temperatures, which protects the bond coat from excessive oxidation. [2]. [b]
While dramatic improvements in power output and efficiency have been achieved by the use of TBCs, further demands for better engine efficiency, reliability and durability raise new problems and challenges to the existing TBCs. First, TBCs are currently not considered “prime reliant” to the extent that they can be used with assurance that spallation failure will not occur [10, 17]. The failure of TBCs can endanger the engine because it exposes the underlying load-bearing alloy to temperatures above its melting point.
The present study objective is to investigate the behavior of the TBC with a thin intermediate α-alumina layer was pre-deposited by sol gel dip-coating techniques on the top surface of APS- CoNiCrAlY metallic bond coat. The APS-YSZ ceramic top coat was deposited on the top of alumina layer. This intermediate alumina layer acts as barrier for oxygen diffusion to protect the metallic bond coat from oxidation. The new thermal barrier coating system was compared with the standard TBC system which comprising of super alloy substrate, bond coat and top coat concerning the microstructure and the behavior under thermal cycling at 1150˚C. The microstructure of the samples and the properties of the coating were investigated before and after thermal cycling by optical microscopy, Scanning Electron Microscopy (SEM), Energy Dispersive X-ray spectroscopy (ED), X-Ray Diffraction (XRD) and Raman spectroscopy. Image analysis software was used compare between the samples by measuring the cracks lengths and the generated (TGO) layer thickness after thermal cycling.
More uniform and compart TGO was observed for the new TBC system. It was found that the alumina intermediate layer has the potential to reduce the average and the maximum length of the cracks and the average thickness of TGO by suppressing the detrimental mixed oxides, like Cr2O3, CoO, NiO, and (Ni, Co) (Cr, Al) 2O4 (spinel phases). Thus, improve the TBC life-time, which is correlated with the increase in the amount of α-Al2O3 in the TGO layer.
2. LITERATURE REVIEW
2.1 COATINGS
Coatings are used to protect the materials for hundreds of years. With the increasing demands that modern technology is placing on materials, coatings are finding more and more uses. In high temperature oxidation field, coatings are used for 2 purposes. The 1st one is to protect substrates, which have high mechanical properties, from high temperature environments. The 2nd purpose is to protect substrates from thermal degradation and oxidation. The protective coatings mainly contain elements such as aluminum, chromium or silicon that can form a protective oxide layer.
2.2. Composition of TBC System
Thermal barrier coatings (TBCs) are commonly used in the fabrication of components which exposed to high temperature in gas turbine for enhancing performance, efficiency and or extended component life. Fig000 shows a Siemens SGT6-4000F gas turbine rotor and part of the ignition chamber coated with TBC with output power 180 MW with efficiency 38%. Fig. 1b shows a turbine blade coated with TBC on the metallic Ni-super alloy substrate with additional coaling film and operates at 1400°C gas inlet temperature, which is more than 100°C above the melting point of the substrate. The selection of TBC materials is constrained by some elementary requirements: 1- relatively high melting point, 2- no phase transformation during heating and cooling, 3- relatively low thermal conductivity, 4- chemically inert, 5- matching thermal expansion coefficient with the metallic substrate, and 6- adherent to the metallic substrate. So far, only a limited materials have been found to satisfy these requirements [4].
Generally, over the years, several ceramic materials have been proposed and tested as TBCs or anti-corrosion coatings in engine parts. However, stabilized zirconium dioxide (ZrO2) acts to be optimal choice as suggested originally by NASA [4]. High temperature tetragonal phase ZrO2 should be stabilized down to ambient temperature. Over decades of research, Y2O3 (7~8wt. %) partially stabilized zirconia (YSZ) met several high temperature applications requirements [5].
TBC system, traditionally, consists of 4 layers as shown in Fig.2.1
(1) A Ni-based super alloy substrate to sustain strength and toughness at high operating temperatures.
(2) An oxidation resistant metallic BC, usually M Cr Al Y (where M= Ni, Co, or both of them) or a platinum aluminide coating. BC is used to overcome the mismatching in thermal expansion coefficient between the ceramic top coat and metallic substrate.
(3) Thermally grown oxide (TGO) layer; some researchers includes the TGO layer as a part of the system despite of it being formed due to due to BC oxidation during processing as well as in service as oxygen goes through the inter-connected pores of the ceramic top coat [3].
(4) The thick insulating ceramic TC, usually 7~8 wt % yttria- partially stabilized or stabilized zirconia (YSZ) [4].
2.2.1 Ni-Based Superalloys Substrate.
Researchers -through the past 50 years- have developed Ni-based super alloys to be used in the high temperature sections of modern stationary turbines [6]. Ni-Based super alloys are able to combine between high melting temperature (~1300˚C), hot strength, high fatigue resistance, superior creep resistance [1005] and high oxidation resistance as well as high stiffness. Other commercial available metallic alloy can’t offer such properties as Ni- based super [7]. Ni-based superalloy matrix has a greater content of Ni, Ti and other refractory elements, such as Mo, Ta, W, and Re mainly to enhance the mechanical properties at elevated operating temperatures.
The microstructure of Ni-base superalloys is based on a γ (Ni) – matrix that usually contains Co, Fe, Cr, Mo, and W in solid solution. The second dominant phase is intermetallic γ’ Ni3 (Al, Ti), which coherently precipitates within the austenitic γ -matrix.
Strengthening of Ni-based superalloys can be done by 2 know mechanisms. The 1st mechanism is solid solution hardening. The 2nd mechanism is precipitation hardening. The theory of hardening in solid solution mechanism depends on preventing the motion of the dislocation by the matrix of the crystal lattice [8]. Chromium and aluminum are refractory elements, while Tungsten, Molybdenum, and Tantalum are important solid solution hardeners elements for Ni-base superalloys, because those elements have an extensive range of solubility in γ -matrix. However the creep strength of the Ni-based super alloy suffers reduction with increasing Chromium content [9]. As soon as the maximum solubility of the austenitic matrix achieved, precipitation of topologically close packed TCP phases, such as ϭ, , χ or Laves detrimental phases can occur [10]. The mechanical properties are affected harmfully by the presence of these detrimental phases. The theory of the precipitation hardening depends on precipitating of strengthening phases like γ’- phase and borides or carbides.
There are 2 main types of Ni-based super alloys: cast alloys and wrought alloys. Wrought super alloys are usually more ductile than cast alloys. This advantage allows producing certain larger shapes, such as turbine discs. Both, solid solution and precipitation hardening (carbide, γ’ and γ\’\'(Ni3Nb) phases) are usually applied for this type of superalloys. This phase provides very high strength at low to intermediate temperatures, but it is unstable above 650°C. Among commercially available wrought Ni-base superalloys there are several with excellent long term performance at 700°C and above. Ni-Fe base polycrystalline alloy IN706, containing about 20% of γ \’\’ is widely used in gas turbine disc application. Cast alloys are normally used, as a structural material for the turbine components with a complicated shape, like blades and vanes. Cast alloys can be polycrystalline (conventionally cast – CC) as well as directionally solidified (DS) or single crystal (SC). [1]
2.2.2 Metallic Bond Coat
The metallic bond coat protects the substrate material against oxidation and hot gases corrosion, also the bond coat is responsible for the adhesion between the Ni-based super alloy as a substrate and the ceramic top coat. Moreover, the bond coat compensates the thermal expansion coefficient mismatch between the metallic super alloy substrate and the ceramic top coat [11].A thin oxide layer called Thermally Grown Oxide (TGO) layer is formed on the top of bond coat under working conditions. This layer protects the substrate from further oxidation.
The commercially used bond coat types are diffusion Platinum aluminide bond coat and overlay MCrAlY bond coat. The Both categories of bond coat super alloys have a high percentage of Al, Cr, and Co. The Al allows the formation of a stable alumina. The alumina offers a protection from excess bond coat oxidation. α-Al2O3 (corundum) has low growth rate and a CPH dense structure, which is stable between room temperature and melting point. α-Al2O3 is, also, thermodynamically compatible with the ceramic top coat, which helps in maintaining their the interface adhesion upon high temperature. Chromium, also, is a common element between the 2 types of bond coat. Chromia protects the substrate at temperatures higher than 850°C from hot gases corrosion. Fig. 000 shows the relationship between the Cr content of different bond coat materials and the capability of the oxidation resistance.
The 1st type of bond coat alloys is diffusion Platinum aluminide coatings. The intermetallic phase β-NiAl (B.C.C -nickel aluminide) is the base phase of this type, which is stable between room temperature and melting point. The Ni content is between 45 and 59 atomic %, as shown on the Ni-Al phase diagram (Fig.000).
β-NiAl can be deposited using pack cementation, slurry coating and chemical vapour deposition [1006, 1007]. The pack cementation is the most widely used due to its simplicity and cost effectiveness [1008] [12]. The pack cementation process for aluminizing contains the substrate, the master alloy (a powder containing Al particles) to be deposited, salt activator which could be NaCl or NaF or NH4Cl and an inert alumina filler powder to prevent the pack from sintering. The substrate is immersed in mixture of master alloy powder, filler and activator. After that, the chamber is heated in hydrogen or argon atmosphere to 1150oC. The Al reacts with the activator at the elevated temperature to produce volatile metal halides. The Al halides diffuses through the gas to deposit on the Ni substrate [13].
There are 2 types of aluminides. The 1st type is low Al-activity bond coat, which happens when the Al content in the master powder is less than 50%. The diffusion in that type occurs outward from the Ni substrate to the Al layer forming β-NiAl layer. The 2nd type is high Al-activity bond coat, which happens when the Al content in the master powder is higher than 50%. The diffusion in that type occurs inward from the Al layer to the Ni substrate forming Ni2Al3 brittle layer. A further heat treatment is required to be converted into required β-NiAl layer.
In commercially used diffusion aluminide bond coat, Platinum is introduced as a diffusion barrier. (Ni,Pt)Al bond coats are deposited by the same technique as aluminides, but a thin Pt layer is electroplated with thickness 5~10 µm before aluminizing as shown in Fig 000. The Pt enhances the oxidation and corrosion resistance. Inward diffusion of AL and outward diffusion of Ni and other refractory elements like Mo, Ta, V, W are reduced by adding the Pt, which increases the coating stability [14]. Pt increases the possibility of forming pure Al2O3 TGO, which protects the bond coat from further oxidation by enhancing the diffusivity of Al in the outer region of bond coat [15].
Fig.00 Cross-sectional microstructure of an as-deposited platinum modified nickel aluminide bond coat fabricated by chemical vapour deposition [1009].
Due to long operation of the turbine parts at high temperature, the Al suffers depletion due to formation of Al2O3 and diffustion through the superalloy. This depletion leads to transform the as-deposited phase β-(Ni,Pt) Al to the martensite phase γ´ -Ni3Al [16]. The ternary phase diagram of Ni-Pt-Al system is shown in Fig. 000 [17]. The martensite γ´ -Ni3Al is a brittle phase in addition to the volumetric change due to phase transformation, may lead to contribute in rumpling of the bond coat layer under thermal cycling [18].
The 2nd type of the bond coat layer is MCrAlY coatings. The “M” of MCrAlY stands for either Co or Ni, or mixture of both, which could be determined according to the type and the purpose of super alloy. Hemker reported that NiCoCrAlY coating has a very high tensile strength ~ 1400 MPa, but the strength is dramatically reduced at elevated temperatures [1010]. Commercial MCrAlY bond coat is deposited by low pressure high velocity oxy fuel spraying (HVOF), plasma spraying (LPPS), vacuum plasma spraying (VPS) and EP-PVD. For NiCrAlY, the possible phases depending on the chemical composition are, as shown in Fig000, the brittle β phase with the ductile γ phase matrix [19]. The β-phase can be considered as a reservoir for AL and the γ-phase enhances ductility of the coating. The commercial bond coat is NiCoCrAlY coatings elements, composition and the purpose of each element is shown in table 1.
Fig.2.9 Cross-sectional microstructure of a NiCoCrAlY bond coat deposited by HVOF [108]: (a) optical image and (b) high-magnification backscattered electron (BSE) image showing that the bond coat consists of β-phase (grey contrast), γ-phase (white contrast) and interfacial pores (dark contrast) between the metal particles
Element Weight % Purpose
Co 20-25 wt.% – Improve ductility at lower temperatures and enhance the creep resistance [19]
– Affects the microstructure by enlarging the β-NiAl phase field
– Increases the amount of Ni, which could be dissolved in β-phase.
Al 8-15 wt.% – Guarantees the formation of alumina scale.
Cr Below 20 wt%. – Improve hot corrosion resistance [1011].
Y Below 1 wt% – Improves the adherence of the oxide layer [20].
– increasing plasticity of the oxides.[21]
– Enhance the high temperature oxidation resistance [d]
Hf (optional) Below 1 wt% – Improve TGO adherence.
– Enhance the high temperature oxidation resistance [d]
The existence of Y oxidizes and Al oxides together at high temperature, as in the case of NiCoCrAlY’s, several compounds can be formed between Y2O3 and Al2O3. At elevated temperatures, as shown in fig000, it is observed that 3 intermediate phases can be formed. The Y3Al5O12 compound is known as yttrium alumina garnet (YAG). YAG is less protective than Al2O3, because of its high oxygen diffusivity of oxygen. The YAlO3 phase is known as yttria alumina perovskite (YAP) and the Y4Al2O9 phase is known as yttria alumina monoclinic (YAM) [e].
Physical and mechanical properties of bond coat are listed in table 00
NiCoCrAlY (Ni,Pt)Al
Young’s modulus 140 and 230 MPa at room temperature and between 60 and 130 GPa at 1000˚C [23]. 115 GPa at room temperature [51].
thermal expansion coefficients at room temperature 13-16*10-6 1/K 13-16*10-6 1/K
thermal expansion coefficients at higher temperature Fig000
the ductile-to-brittle-transition temperature
(DBTT) ranges from 200˚C for low Al coatings (6-9 wt.% Al) to 700 ˚C for Al-rich coatings (12 wt.% Al) [25]. from 600 to about 750 °C
11111
At higher temperatures they show significant variation in thermal expansion behavior as a function of chemical composition and temperature (Fig.00000, 00000).
2.2.3 Ceramic Topcoat
The topcoat is a ceramic layer, which faces the thrust hot gases. This layer is mandatory to create the temperature gradient across it’s thickness to protect the substrate from the high temperature. Therefore, the top coat layer should have high melting point, high surface emissivity, low thermal conductivity, low vapor pressure, strain tolerance, high resistance to thermal shock, chemical stability, high hardness, resistance to oxidation or chemical degradation, and relatively high coefficient of thermal expansion. The 1st use of ceramic coating for aerospace applications was frit enamels. This ceramic coat was developed by the National Advisory Committee for Aeronautics (NACA). Then, the coating of calcia stabilized zirconia on the nozzle of the exhaust of the X-15 manned rocket in 1960s is believed to be the 1st usage of TBC system in manned flight. Other ceramics like, CaO/ MgO+ZrO2, CeO2+YSZ, zircon and La2Zr2O7, etc. had been evaluated as TBC materials. Currently, Yttrium Stab YSZ layer with thickness 100~400 μm, is the most preferred as TBC ceramic top coat [14].
ZrO2 has a low thermal conductivity at elevated temperature. As well as, a rather high coefficient of thermal expansion ~ 9*10-6 ~10*10-6 K-1 fig000, which is almost close to the underlying bond coat coefficient of thermal expansion 12*10-6 K-1. ZrO2 is the base material of choice for the ceramic top-coats of high temperature components [4, 26]. ZrO2 has very melting temperature ~ 2700˚C and the high hardness ~ 14 GPa at room temperature. The main crystallographic structures of zirconia are cubic, tetragonal and monoclinic. The change between the phases depends on the temperature as shown in the phase diagram of ZrO2-Y2O3 (Fig. 000). Pure zirconia has 3 main phases, cubic structure occurs between 2680˚C (melting temperature) and 2370˚C, Tetragonal structure between 2370˚C and 1170˚C and monoclinic structure at temperatures below 1170˚C. The transformation to the monoclinic structure is associated with 4-5 % volumetric expansion [26]. To prevent this transformation, dopants, like Y2O3, CaO, and Yb2O3, are added to stabilize cubic or tetragonal phases at lower temperatures. The most commercially applied top coat is 7wt % Y2O3 partially stabilized ZrO2 (7YSZ) [27, 28].
Fig.2.2 Coefficients of thermal expansion and thermal conductivities of a range of engineering materials [1005]
TBCs should be strain tolerant to protect the coating from delamination and spallation by incorporating either micro cracks or aligned porosity [7]. The top coat could be deposited either by Air Plasma Spray (APS), or Physical Vapor Deposition- Electron Beam (EB-PVD).
i. Plasma sprayed TBC
Plasma spraying is the most common technique for manufacturing ceramic coatings. Plasma spraying can deposit till several millimeters thickness of ceramic materials. The theory of the plasma spray is shown in fig000. The powder are melted and accelerated towards the target by injecting from the injection nozzle into high-temperature plasma gas stream (plasma jet), which created inside a plasma gun. An electrostatic field controls the direction of the powders towards the target. Then, the molten metal spreads out and solidifies onto the substrate surface .The main types of the plasma spraying technique are Air Plasma Spray (APS), Low Pressure Plasma Spray (LPPS) and Vacuum Plasma Spray (VPS).
APS is the most suitable method for YSZ deposition. There are many parameters affecting the microstructure of the deposited layer, such as; the density of the coating powder, the size of the particles, particles morphology, melting temperature, particles velocity, etc. Generally, the morphology of the deposited plasma sprayed coatings exhibit a lamellar structure of the spraying splats with intersplat boundaries and cracks oriented parallel to the bond coat surface interface (Fig. 000). Hence, APS-TBC is less strain tolerance compared to EB-PVD TBC and has shorter thermal cycling life. The deposited layer porosity is about 20-25 vol.% [29]. The porosity of the layers strongly affects the thermal conductivity of the YSZ. As the thermal conductivity is reduced when increasing the porosity %. The thermal conductivity of APS-YSZ layer is about 0.8 Wm-1K-1, while the fully dense YSZ is about 2.5 Wm-1K-1. The top surface of the bond coat should be rough enough before the applying the top coat to provide good mechanical serration between the top coat and the bond coat. On the other side, the surface roughness induces higher internal stresses between the adjacent layers. Generally, APS- TBC is used, only, on stationary parts, like combustor cans, fuel vaporizers, vanes and shrouds.
ii. EB-PVD TBC
Physical vapor deposition (PVD) had been developed in the 1960 as one of the main coating deposition techniques. PVD refers to deposition of layers by vapor transport in a vacuum without any chemical reaction. The theory of operation of EB-PVD technique is transformation of the kinetic energy of a high power density heat source electron beam (100 – 200 kW) into thermal energy to evaporate the ingot (coating material). The process is operated under vacuum conditions of 10-2 Pa or higher. The target (substrate) is pre-heated and positioned above the cloud of the vapor as shown in Fig.000. EB-PVD can deposit up to 300 µm thickness. There are 2 different modes of the PVD technique, the 1st mode is stationary mode, and the 2nd one is rotating mode. The rotating mode is applied with the aim of getting a homogeneous thickness through the coating layer. Rotating mode is, specially, used with complex geometry parts, like turbine blades [30].
The morphology of the EB-PVD-YSZ deposited layers, as shown in Fig.000, has a columnar structure with fine scale of porosity through the discrete columns [2]. The columnar structure is because of the atomistic nature of the deposition technique, in which the growth of the grains are governed through the condensation of the ceramic ingot vapor [1009]. The columnar structure offers a strain tolerance for the layer [15], because the columns can separate, accommodating thermal expansion mismatch, as described by Strangman [31], so that , EP-PVD TBC has higher thermal cycling life.
Rotating EB-PVD method is the applied technique for TBC ceramic top coat deposition. The microstructure of rotating EB-PVD top coat can’t be homogenous. The 1st region of the top coat layer close to the interface is equiaxed grains structure. Most of TBC failures occurs in this region [31]. Directly over the equiaxed region a columnar structure region is apparent. The thickness of the layers depends on the rotation conditions [32]. EB-PVD topcoat has the following properties: high resistance for thermal shock, smooth aerodynamically surfaces and good erosion resistance. The main physical and mechanical properties of EB-PVD TBC top coat are:
– Thermal conductivity 1.5 ~ 1.9 Wm-1K-1,
– Coefficient of thermal expansion 9*10-6 K-1 at room temperature and 12*10-6K-1 at 1100˚C,
– Young’s modulus 150 GPa at room temperature and 22-54 GPa at 1100˚C [33].
EB-PVD top coat is more expensive and durable than APS-TBC. PVD is primarily used in most severe applications, such as jet engines turbine blades and vanes.
iii. Comparison of EB-PVD and APS TBCs
APS EB-PVD
less expensive More expensive
thickness up to 3 mm not thicker than 300 μm
Better reproducibility.
Lamellar porous structure columnar structure
Lower thermal conductivity higher thermal conductivity
Higher thermal shock resistance Lower thermal shock resistance
2.3 TBC failure by thermal cyclic loading
TBC systems failure under thermal cycling is an intricate process. The failure process is affected by the change of material properties because of the operation temperature and the internal stresses. Generally, there are 2 types of fracture for TBC ceramic layer, segmentation and delamination. The Segmentation cracks propagates normal to interface between the top coat and the bond coat (vertical orientation). These vertical cracks can be considered as a favorable defect with point of view of the strain tolerance of the top coat [34]. As the vertical cracks behaves like the gaps between the columnar grains in EB-PVD top layer. On the other hand, delamination cracks (horizontal orientation), which occurs parallel to bond coat/top coat interface, can cause spallation of TBC coating, hence failure of the system [35].
There are many reasons, which are responsible for the change of physical and mechanical properties of TBC system during thermal cycling:
– Oxidation of the bond coat
– Diffusion of elements between Ni super alloy and bond coat;
– Ceramic Top coat sintering;
– Residual stresses, which is typically associated with oxide scale formation, oxides growth and cooling stresses due to coefficient of thermal expansion mismatch.
– Crack formation and propagation.
– Stresses due to substrate geometry. e.g. substrate curvature [36].
– Bond coat relaxation.
The above aspects have been focused in the literature for several years. Some of the obtained are described in the following sections.
2.3.1 Oxidation of the bond coat (Bond Coat degradation).
Under operation conditions the TBC suffers bond coat oxidation. The ceramic top coat is transparent for the hot gases; especially oxygen; due to its porosity and high ionic diffusivity. Incessant transport of oxygen ions through the ceramic top coat resulting in formation of a compact TGO layer along the bond coat / top coat interface Fig000. Some types of the generated oxides protects the metallic bond coat from further oxidation. In a TBC systems, the chemical composition of the bond coat affects the oxidation behavior of it [40]. For both commercially used bond coats types ((Ni,Pt)Al, or NiCoCrAlY). The grown oxide scales consist mainly of alumina (Al2O3), Cr2O3 (chromium oxides), NiO (nickel oxides), and (Ni, Co) (Cr, Al) 2O4 (spinel phases). The Al2O3 has several phases, like α (Rhombohedral), δ (tetragonal), θ (monoclinic) and γ (cubic). The most desirable, stable, slow growing phase is α-Al2O3, which originates at temperature higher than 900˚C [16]. α-Al2O3 has a very low oxygen ionic diffusivity and provides an excellent diffusion barrier, retarding further bond coat oxidation The other alumina phases are metastable and non-protective so that they are called transient oxides.
There are 2 main factors affecting the growth rate of the TGO layer. The 1st factor is the transportation of the oxygen through the grain boundaries of the top coat towards the interface with the bond coat. This type of oxidation is called internal oxidation [41]. The 2nd factor is the outward diffusion of Al cations through the generated alumina layer to react with oxygen anions. This type of oxidation is called external oxidation [42]. Consequently, the depletion of the Al in the bond coat will results in the development of β-depleted zone. This depletion zone occurs usually close to the interface between the bond coat and the top coat [43]. A volumetric expansion occurs with the internal oxidation of the bond coat, which lead to high compression stress within the interface (> 1GPa) [44, C], which play an important role for the adhesion of oxide scale [45].
After the depletion of the aluminum in the MCrAlY bond coat, other non- protective scales will start to form. Ulrike Täck [46] found that, MCrAlY bond coats with Al-content less than 5 wt. % wouldn’t form a protective alumina scale layer. The non-protective oxides, such as, CoO, Cr2O3, NiO, and Ni (Cr, Al)2O4 spinels appear at the interface [47, 48]. A rapid volumetric expansion combined with the formation of these oxides, which, led to massive local internal stresses, then failure occurs [49].
V.K. Tolpygo introduced a method for enhancing the durability of EB-PVD bond coat in TBC system. The method is based on pre-oxidation of the bond coat. Peroxidation means getting an oxide layer on the bond coat before depositing the top coat as shown in fig.000. It was found that the pre-deposited α-alumina scale is less susceptible to separation along the bond coat/top coat interface, as well exhibiting a low growth rate, good adhesion comparing to the developed TGO on a similar bond coat without pre-oxidation. The method produces a twofold TBC service life time [37-old]
W.R. Chen mentioned that the formation and growth of a TGO layer is a major factor influencing the durability of TBC system. The TGO has been identified to be (Cr, Al) 2O3, Ni (Cr,Al)2O4 and NiO, which grows in a parabolic manner. The oxide clusters were found to be the preferred sites for crack nucleation and propagation into the ceramic coating, causing premature TBC separation. W.R. Chen modified the microstructure of TGO by forming a thin α-Al2O3 layer by a low-pressure oxidation treatment (LPOT) for the bond coat before top coat deposition, which acts as a diffusion barrier to suppress the formation of other detrimental oxides during service. A heat treatment in LPOT promotes the formation of a nearly continuous alumina layer and partially reduces the detrimental oxides. More research is required to further understand and subsequently reduce the formation of Ni (Cr,Al)2O4 and NiO particles [38-old]
The surface condition of the coatings when they are exposed to oxidizing environments has been shown to play an important role in their lifetimes. Different deposition methods can produce wide variations in the surface roughness of MCrAlY coatings. Surface treatments that may be performed on coating surfaces include hand-polishing, rough grinding, and grit blasting.
Recent studies have shown that in NiCoCrAlY coatings exposed in their rough, as-sprayed state at 1100°C, the alumina scales had a completely different morphology than those grown on a coating surface that had been ground flat [3]. The scale formed on the flat surface had a uniform thickness and a homogenous distribution of Y-Al mixed oxide precipitates beneath the surface of the coating.
2.3.2 Inter-Diffusion between substrate base material and bond coat
The definition of the inter-diffusion is the transportation of the material by atomic motion. The inter-diffusion occurs because of the atomic concentration gradient. In TBC system, MCrAlY bond layer has a high content of Cr and Al content comparing to the Ni-based superalloy substrates to stimulate the development of a protective oxides. This high difference in concentration results in elemental diffusion of Al and Cr from the bond coat to the Ni-based superalloy at high temperatures. As a result of Al diffusion from the bond coat at working temperature, the zone of the bond coat near to the interface with the substrate, will be β-depleted zone. γ’phase will be formed instead of β-phase.[fig 00 phase diagram of bond coat], . Busso found that the thickness of the β-depleted zone is about 80 µm near to the bond coat/substrate interface. Busso measured the thickness after thermal cycling at ambient air for TBC system with maximum temperature 1090°C [50].
Furthermore the underlying Ni-based superalloy matrix has a greater content of Ni, Ti and other refractory elements, such as Mo, Ta, W, and Re. Consequently these elements, also, diffuse from the Ni- based super alloy substrate into the coating as shown in Fig.000 [46]. The inter-diffusion from the Ni-based superalloys to the bond coat results in the formation of harmful phases, like σ (sigma)-phase, laves phases [50], cavities and porosity [51]. O.Knotek observed that, applying a barrier layer from (W, Al-O-N, or PtAl2) can confine the inter-diffusion process between the substrate and the MCrAlY bond coat [52].
2.3.3 Sintering of YSZ
i. APS TBC
During the operation, the TBC system is exposed to high temperature. This temperature causes densification and sintering for the APS-ceramic top coat. This process affects the durability of TBC negatively [53]. The densification of the ceramic layer decreases the volume fraction of voids and defects within the layer, which increases the stiffness of the ceramic. The stiffness of APS-YSZ increases because of healing of the micro-cracks above 900˚C [54]. The stiffness has a direct effect on the stress distribution in TBC multilayer system by reducing the compliance of the ceramic layer. Moreover, the decrease of the voids through the ceramic layer leads to increase thermal conductivity of it [55].
WANG Kai studied the effect of sintering behavior with the detailed data of the ceramic coating porosity, thermal conductivity and thermal barrier effects to provide quantitative analysis of the sintering effect on the life time [34-old]. It was reported that the microstructure of the coat changed significantly with heat treatment time due to coating sintering. Remarkable increases of the thermal conductivity occur to both typical coatings after heat treatment. The grain size of the nanostructure zirconia coating increases drastically with annealing time, which indicates that coating sintering makes more contribution to the thermal conductivity of the nanostructured coating than that of the conventional coating [34-old].
A Portinhaa developed a new system of TBC that consist of a conventional NiCoCrAlY bond coat and APS ZrO2–8 wt. %Y2O3 top coat graded in porosity along the corss section by modifying the deposition parameters, as shown in fig 000. The aim is to produce coatings with lower thermal conductivity and higher thermal shock resistance thermo-mechanical properties, because of the gradient in porosity which reflects a gradient in the elastic properties [35-old]. It was observed that the top coat hardness drops from the interface to the surface and increases after annealing as expected. The new TBC system has higher service life comparable to the conventional one [35-old].
ii. EB-PVD TBC
The EB-PVD ceramic YSZ top coat suffers impairment due to the sintering process, while exposing to high temperature operation [56, 57]. The structure of the EB-PVD ceramic top coat is columnar with gaps around the separate columns. These columns shouldn’t be in contact to provide strain tolerance. Due to the plumy structure of the deposited columns, still there is a contact sites exist due to the sintering process called “necks” as shown in fig 000. The necks decreases the stiffness, hence; reducing strain tolerance behavior [58].
2.3.4 Residual stresses
The TBC systems are subjected to 2 types of stresses. The 1st type is the applied mechanical stresses, such as tensile stresses due to the high rotational speeds of the turbines. The 2nd type is the residual stresses of the multilayer TBC systems. The residual stresses affect the thermo-mechanical integrity of the system significantly. There are a lot of studies examined the failure of the TBC systems to get a better understanding of the stresses state in each layer and to study the effect of the stresses on the ceramic layer at the form of cracks, spallation, delamination [15, 43, 59, 60].
Mostly the residual stresses in the ceramic layer generated either from quenching or operation. The quenching residual stresses arise because of the quick solidification and contraction of the sprayed molten particles of the top coat. The quenching stresses always make the top coat under tension. The operational residual stresses are independent of the technique of layer deposition. The operation stresses occurs due to the thermal expansion mismatches between the separate adjacent layers of the thermal barrier system. The operation stresses are changed cycle by cycle in case of thermal cycling exposure because of the microstructural changes in the layers, type and thickness of the generated TGO. The life time of TBC systems in several cases can be correlated with TGO phases and average TGO thickness [42, 61].
Conversely, the lateral growth of TGO and the thermal cycling conditions induce a huge out-of-plane stresses, which results in generating and propagating of the delamination cracks [37].
The residual stresses at the layers can be tension or compression, depending on the values of the coefficient of thermal expansion of the adjacent layers. Usually, the ceramic top coat is under compression (up to 6 GPa) [f, g] at low temperatures and under tension at high temperatures [62].
Initial phase constituents of the TGO are also considered to be a critical factor influencing the integrity of TBCs. The formation of the metastable γ- and θ-Al2O3, and their transformations to the equilibrium α-Al2O3 in the TGO scale have been reported to have a profound effect on the structural integrity of TGO/YSZ interface during thermal cycles.24,79 These polymorphic transformations of Al2O3 may be responsible for additional residual stress from the volumetric constraint in the TGO, and nucleation of sub-critical defects, eventually leading to the spallation of the TGO. For example, there is about a 4.7% volume change for θ- to α-Al2O3 phase transformation. Thus, the formation of a uniform and defect-free α-Al2O3 scale prior to the deposition of topcoat can help to enhance the durability and reliability of TBCs.
Thera are many successful methods had been used to determine the residual stresses in the TBC systems. Scardi, P. used X-ray diffraction (XRD) to measure at high accuracy the stresses close to the surfaces in TBC systems because of its limited depth of XRD [63]. XRD is preferred a method to evaluate the stresses in dense layers like the TGO [63]. Tomimatsu, T. used Raman photoluminescence piezo-spectroscopy to determine the stresses in the ceramic layers [64].There are several methods based on the change in the strain of the layer due to material removal, such as hole drilling [65]. Specimen curvature measurement method can only be used for large area samples [66]. Residual stresses can be, also, predicted by analytical models by defining the system main parameters like; the thermos-elastic properties, geometry parameters, the interface topography between the layers, stiffness, thermal expansion coefficient, the thickness and the temperature [67, 68].