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Spark plasma sintering of a multilayer thermal barrier coating on Inconel 738 superalloy: Microstructural development and hot corrosion behavior
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CERAMICS INTERNATIONAL
Ceramics International 42 (2016) 2770–2779 www.elsevier.com/locate/ceramint
Spark plasma sintering of a multilayer thermal barrier coating on Inconel 738 superalloy: Microstructural development and hot corrosion behavior A.H. Paksereshta, A.H. Javadib, M. Bahramic, F. Khodabakhshid, A. Simchic,e,n a
Department of Ceramics, Materials and Energy Research Center, P.O. Box 31787-316, Karaj, Iran b Engineering Research Institute, Tehran, Iran c Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran d Department of Materials Science and Engineering, School of Engineering, Shiraz University, Zand Boulevard, Shiraz, Iran e Institute for Nanoscience and Nanotechnology, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran Received 20 August 2015; received in revised form 2 November 2015; accepted 2 November 2015 Available online 10 November 2015
Abstract In the present work, spark plasma sintering (SPS) process was employed to prepare a nanostructured yttria-stabilized zirconia (8YSZ) coating on a nickel-based superalloy (INCONEL 738) with functionally graded structure. A stack layer of INCONEL 738/NiCrAlY powder/Al foil/ NiCrAlY þYSZ powder/YSZ powder was SPSed in a graphite die at an applied pressure of 40 MPa under an vacuum atmosphere (8 Pa). The sintering temperature was 1040 1C. For comparison purpose, the air plasma spray (APS) technique was employed to prepare the thermal barrier coating (TBC). Microstructural studies by scanning electron microscopy showed that the SPSed coating was sound and free of interfacial cracks and large pores. Energy dispersive X-ray spectroscopy determined a limited inter-diffusion between the layers. A significant improvement ( 26%) in the Vickers micro-hardness was also measured. In order to examine the functionality of the coatings, hot corrosion test was performed at 1050 1C in a molten bath of Na2SO4 and V2O5 (45 to 55 weight ratio). The results showed that spallation occurred after 12 and 20 h for the SPSed and APS coatings, respectively. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Multilayer thermal barrier coating; Spark plasma sintering; Microstructure; Hot corrosion; Thermally grown oxide
1. Introduction Thermal barrier coatings (TBCs) are broadly utilized as a protection layer for the superalloys in order to reduce thermal oxidation and corrosion rate, for example, in gas engine turbines and aeronautic diesel engine chambers [1–4]. TBCs have a multi-layer composite structure which commonly consist of a superalloy substrate covered with one or more metallic bond coat (BC) followed by a thermally grown oxide (TGO) layer and finally a ceramic top layer [5–8]. MCrAlY (M ¼ Ni or Co or both) alloys or a platinum aluminized coating n Corresponding author at: Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran. Tel.: þ98 21 6616 5261; fax: þ98 21 6600 5717. E-mail address: [email protected] (A. Simchi).
http://dx.doi.org/10.1016/j.ceramint.2015.11.008 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
are typically used as BC to enhance the high temperature oxidation and corrosion [5,9,10]. The TGO layer operates as a diffusion barrier against hot oxidation which enhances the system thermal stability under service conditions [11]. Yttriastabilized zirconia (YSZ) is the most common ceramic material used in TBCs due to its low thermal conductivity that reduces the peak temperature for the internally cooled underlying substrates [6]. Recent studies have demonstrated the advantages of nano-sized YSZ for TBCs [12–14]. As compared with conventional YSZ coatings, a higher coefficient of thermal expansion, lower thermal diffusivity, higher hardness and toughness, and better wear resistance have been reported [12,15,16]. Physical vapor deposition (PVD) [17,18], low pressure plasma spray (LPPS) [19], vacuum/air plasma spray (VPS/APS) [20– 23], high velocity oxyfuel spray (HVOF) [24], and electrolytic
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deposition (ED) [25] have already been employed for the preparation of TBCs. Among these methods, APS is more practical and has been utilized broadly for different systems [26– 30]. Nevertheless, processing of nanostructured TBCs by APS is challenging due to a poor flowability of nanoparticles (which requires proper granulation) and grain growth caused by partial melting of the particles [16,31,32]. The deposited material by APS often displays a bi-modal grain structure induced by consolidation of melted and non-melted regions [33]. Recently, SPS has attracted great interest for the fabrication of functionally graded materials and nanostructured TBCs [1,34,35]. Jiang et al. [36] employed this process to prepare WC coating on tungsten substrates. Boidota et al. [1] and Song et al. [5] prepared TBCs on superalloys by SPS. The process has a number of advantages over the conventional methods such as a fast processing cycle (high heating rates with short dwell time), energy and material saving, limited grain growth and its efficiency to sinter materials with limited change of their initial microstructures [5,37]. Therefore, SPS is potentially very attractive for the fabrication of nanostructured coatings. Recently, multilayered or functionally-graded structures have been utilized to prepare TBCs [38–41]. Gradual compositional variations can reduce the residual and the thermal stresses between the top layer and BC, which improves the bonding strength [38–41]. However, fabrication of multilayered structures often requires multi-step processing methods. The aim of the present work is to prepare multilayered TBCs by SPS through a strightforward proces. To the best knowledge of the authors, no results on development of functionally-graded TBCs by SPS method have been reported so far; though FGM TBC is as a multi-layered structure with gradiently variation of YSZ volume fraction from the ceramicrich upper-layer toward the alloy-rich TBC bond coat interface. Nanostructured YSZ powder was utilized as the raw material to reduce the sintering temperature and to enhance the performance of the multi-layered TBC on INCONEL 738 (IN738). For comparison purpose, a TBC with the same composition and layer structure was fabricated by APS. Since hot oxidation and corrosion are the main phenomena restricting the application of TBCs in harsh atmospheres [42,43], hot corrosion resistance of the multilayered TBCs was studied in a molten V2O5/Na2SO4 salt. It is well known that the molten salt reacts with the top ceramic layer or even penetrates through its defects (pores, micro-cracks and columnar gaps) during thermal cycling under high temperature, which leads to the tetragonal to monoclinic phase transformation, delamination and spallation of the coating [8,42, 44–47].
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2. Experimental procedure 2.1. Spark plasma sintering of multilayered coating The characteristics of the materials used in the present work are shown in Table 1. Discs of Inconel 738 with a nominal chemical composition of Ni–15.48Cr–8.57Co–2.99Al– 2.12Mo–3.17Ti–2.99W–1.09Ta (wt%) in diameter of 30 mm and height of 5 mm were prepared by machining followed by surface grinding (2000 SiC papers) and acetone washing. To prepare the multilayered TBC by SPS, the superalloy substrate was inserted into a graphite die with an inner diameter of 30 mm. In order to prevent carbon diffusion and reaction, a thin layer of hex-BN powder (TERIO HBN, China) was deposited on the die surface by spraying. The NiCrAlY powder ( 1 g) was then poured on the surface of the substrate and leveled off by hand. A foil of commercially pure aluminum with a thickness of 4 μm was then placed on the top and a mixture of YSZ and NiCrAlY powder at a weight ratio of 1:1 was spread. The aim of this layer was to generate a thermally grown Al2O3 oxide layer to keep the interface flat and homogeneous, and to enhance the adhesion to the top layer. Finally, nanostructured 8YSZ (8 wt% Y2O3 þ 92 wt% ZrO2) powder was inserted on the top layer. A schematic imge of the multilayered TBC on IN738 superalloy is shown in Fig. 1a. As it is seen, the multilayered coating is composed of five layers: (i) Inconel substrate; (ii) a bond coat layer from NiCrAlY powders with a thickness of 100 μm; (iii) a thin aluminum layer (for TGO layer formation) with a thickness in the range of 6–8 μm; (iv) an intermediate functionally graded layer (50% 8YSZ and 50%NiCrAlY) with a thickness of 200 mm; (v) YSZ top layer with a thickness of 200 mm. It is pertinent to point out that this structure has been determined by various efforts to attain specimens without having defects such as unbounded layers, large cracks, and delamination. A commercial SPS apparatus (EF 20T-10, Easy Fashion Metals Products Co.) was utilized for the sample preparation. A schematic view of the machine operation is shown in Fig. 1b. The SPS route was carried out in a multi-stage thermal/pressure cycling under vacuum (8 Pa), as shown in Fig. 1c. In the first stage, pre-heating to 600 1C and holding for one minute was performed for thermal stabilization. A pressure of 10 MPa was applied at the start and then increased to 20 MPa at 600 1C. After that, the temperature raised to 1040 1C at a rate of 35 1C/min and the sample was held for 10 min. The applied pressure for sintering was 40 MPa. The specimens were then cooled to 600 1C within 10 min. Visual
Table 1 Characteristics of materials used in this study. Materials
Commercial name
Chemistry (in wt%)
Form
Fabrication route
Nickel base superalloy NiCrAlY Yettria-stabilized zirconia Aluminum foil
INCONEL 738 Metco 442 8YSZ (US3630) 1050Al
Ni–15.48Cr–8.57Co–2.99Al-–2.12Mo–3.17Ti–2.99W–1.09Ta Ni–8.5Cr–7Al–5Mo–2Si–2B–2Fe–3Al2O3 92ZrO2–8Y2O3 (high purity499.9%) Al–0.05Cu–0.05Mg–0.25Si–0.4Fe–0.05Mn–0.07Zn–0.05Ti
Plate (substrate) Powder ( 70 μm) Powder ( 40 nm) Sheet
Vacuum melting Mechanically clad Agglomerated/sintered Rolling/annealing
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Fig. 1. (a) A schematic representation of multilayer TBC on Incolnel 738 superalloy. (b) Schematic of spark plasma sintering accomplishment. (c) Termal/pressure cycling used for SPS of the multilayered TBC.
and 325 US mesh) to attin granulates with 44–90 mm sizes. Finally, heat treatmrnt was performed at 1100 1C for 2 h.
Table 2 Processing parameters used for APS of different layers. Parameter
Current (A) Voltage (V) Primary gas, Ar (l/min) Secondary gas, H2 (l/min) Powder feed rate (g/min) Spraying distance (cm)
Different layers
2.3. Microstructural evaluation and hardness measureemt
NiCrAlY
Graded layer
YSZ
450 75 55 17 35 11
550 70 36 17 30 8
550 70 36 17 30 8
examination revealed that the samples are sound but in a dark color (Fig. 1d). It seemed that surface carbon soating and/or carbonization of the substrate was taken place, albeit a barrier coating (hex-BN) was utilized. Therefore, the SPSed samples were heat treated in an electrical furnace at 1000 1C for 2 h. The mentioned time for post annealing phemomenon was selected based on previous studies [1,5], which aimed to eliminat carborization as well as detecting the γ and β phases.
For microstructural studies, the multilayered samples were cross-sectioned by a diamond cutting machine (DCS 150, Synova). The specimens were then cold mounted in an epoxy resin and grined by SiC emergy papers down to 2000 mesh. Diamond pastes (10, 3, and 1 μm sizes) were used for mirror polishing. A field emission-scanning electron microscope (FESEM, VEGA//TESCAN-XMU, Russia) equipped by an energy dispersive X-ray spectroscopy (EDS) was employed for microstructural examines. Possible phase transformations were analyzed by using an X-ray diffractometer (XRD: GNR, Explore, Analytical Instruments Group, Netherlands) with Cu Kα radiation. Variations in the micro-hardness profile along the cross-section were determined by a Vickers indenter (MVK-H21, AKASHI, Japan) at an applied load of 50 g for a dwell time of 15 s. The indentation testing was repeated ten times for each region and the average value was reported with standard deviation (SD).
2.2. Air plasma spraying
2.4. Hot corrosion test
Plasma spraying was performed with a Sulzer-Metco F4MB plasma gun (Sulzer-Metco, Switzerland) in air. Argon was used as a primary and carrier gas and hydrogen was utillized as a secondary gas. The processing parameters for different layers are shown in Table 2. To improve the flowability of thefeed stock powders, the YSZ particles were granulated by utilizing polyvinyl alcohol (Junsel Chemical Co. Ltd., Japan) as a binder [48]. The aggromerates were screened by sieving (180
A batch of molten V2O5/Na2SO4 salt (55/45 wt%) was prepared by mixing of the constituents at room temperature followed by heating at 1050 1C. The corrosive salt was then spread over the TBC in a concentration of 30 mg/cm2. To avoid the edge effect [42,44,45,49], no covering was considered near the edges for apart 3 mm. The prepared setup was placed in an electric furnace with a fixed temperature of 1050 1C in atmospheric condition. Every four hours, the
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Fig. 2. A cross-sectional FE-SEM image combined with elemental EDS maping for the multilayer TBC produced by SPS followed by post-heating at 1000 ºC for 2 h.
specimens were taken out from the furnace and analyzed for possible cracking, spallation, and severe formation of corrosion products. 3. Results 3.1. Microstructure A typical cross-sectional FE-SEM image of the multilayered coating combined with EDS elemental mapping of layers is shown in Fig. 2. The specimens consist of five layers of IN738/NiCrAlY/Al/NiCrAlY þ 8YSZ/8YSZ. The layers are dense and free of common defects like pores and microcracks. Formation of a TGO layer with a mixed composition, mainly alumina (Al2O3) with some zirconia (ZrO2) inclusions,
was also detected. To determine the composition of the bond layer, EDS analysis of selected regions (A, B and C in Fig. 2) was performed. The results are shown in Fig. 3. The formation of γ(Ni) solid solution with dark gray color and β(NiAl) phase with bright gray color was noticed Typical cross-sectional FE-SEM image of the TBC coating prepared by APS is shown in Fig. 4a. Unlike the SPSed coating, non-bonded interface, voids, micro-cracks and intersplats boundaries were observed. These defects were partly formed due to presence of un-melted/semi-melted particles, weak adhesion of splats to the underlayers, evaporation of condensate, loss of solid particles and also a pull-out of particles in sample preparation [50,51]. To better show the microstructure of the prepared TBCs, the YSZ layer was broken and the resulting cross-section was observed by FE-SEM. Fig. 4b shows a typical
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Fig. 3. Results of EDS analysis performed on locations (a) A, (b) B, and (c) C shown in Fig. 2.
fracture surface of the APS sample (near the top YSZ layer region). A lamellar and bi-modal structure with micro- and submicro-sized grains (melted and non-melted regions) is seen. It was concluded that APS changed the initial nanostructure of the particles to a coarse structure with pore formation. This microstructural changes occurred due to high processing temperature, partial or complete melting of particles, and formation of individual splats during solidification of molten droplet due to collision. Moreover, some degrees of cross-sectional inhomogeneities are noticeable, which are attributed to the functionally graded structures of the APSed and SPSed TBCs. Fig. 4c shows the fracture surface of the SPSed specimen. A dense structure
with nanometric grains (average size of about 60 nm) is visible. The short sintering cycle at relatively low temperature in the presence of applied pressure inhibited significant grain growth while preceding rapid densification. As it will be shown in following, the presence of pores and defects in the top coating layer facilitates penetration of the molten salt into the underlayers, degrading the hot corrosion resistance. 3.2. Mechanical properties Variations of micro-hardness along the cross-section of the prepared specimens are shown in Fig. 4d. The results indicate a
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Fig. 4. (a) Cross-sectional FE-SEM image of the multilayer TBC prepared by APS. Cross-sectional FE-SEM images of the deposited YSZ by (b) APS and (b) SPS after fracturing. (d) Mean values of Vickers micro-hardness across thickness section for the prepared TBCs.
relatively smooth increase in the hardness from the superalloy substrate to the ceramic layer. Interestingly, the hardness of the multilayered coating prepared by the SPS is higher than the APS one, particularly in the top ceramic layer. While the SPSed intermediate layers exhibited a small improvement in the mean hardness value (about 10%) compared with the APSed ones, the hardness of the top coating significantly increased from 620 HV to 780 HV (26% enhancement). The capacity of the SPS process in full densification while retaining nanostructured grains improved the indentation resistance of the coatings. 3.3. Hot corrosion behavior The hot corrosion response of the TBCs was evaluated by V2O5/Na2SO4 melt reaction at 1050 1C. The corrosion products were analyzed by XRD. Fig. 5a shows XRD patterns of
the top coatings prepared by APS and SPS. Both coatings have a tetragonal zirconia crystallographic structure. The pattern of the SPSed specimen exhibits peak-line broadening which is an indicator of its finer grain structure [52,53]. XRD patterns of these samples after hot corrosion test are shown in Fig. 5b. A mixture of monoclinic (m-ZrO2) and tetragonal (t-ZrO2) zirconia together with corrosion product (YVO4) were noticed. Results of microstructural studies after hot corrosion test are shown in Fig. 6. It appears that elongated phases are formed due to reactions at the elevated temperature (Figs. 6a and 6b). EDS analysis support the XRD results, indicating that the main product is yttrium-vanadium oxide (Figs. 6c and 6d). Large and needle-like crystals (with a length of 70 mm and larger) were detected for the APS coating while smaller flower-like crystals (with a length of 30 mm and smaller) were observed for the SPS layer. Both XRD analysis (peak intensity) and
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studying of the phase transformation before and after air annealing is useful to get more insight on the possible mechanisms of TGO layer. Many studies have already been performed on this issue, for example [1,5,55], and it is determined that reactions of the aluminum folil with both oxidizing atmosphere and zirconia are feasible. It is well known that hot corrosion resistance of TBCs strongly depends on their chemical compostion, grain size, and structure homogeneity [56]. Pores and micro-cracks are suitable sites for molten salt penetration through the coating during hot corrosion testing. XRD analysis showed that after hot corrosion test, monoclinic zirconia (m-phase) was formed. To determine the amount of m-phase in the coatings, the following equation was utilized [37–40,48]: %M ¼ ðM 1 þ M 2 Þ=ðM 1 þ M 2 þ I Þ
Fig. 5. XRD patterns of the prepared TBCs (a) before and (b) after hot corrosion test.
microstructural examinations revealed that the volume of the corrosion product on the SPS specimen was lower than that of APS sample.
4. Discussion The formation of TGO layer during SPS is a results of reaction between the aluminum thin foil and ZrO2 through oxygen diffusion from the zironia ceramic towards the interface and/or trapped oxygen in the sintering chamber, which can form oxidizing atomosphere with the graphite die (CO/ CO2) gas mixture [5]. Microstructural analysis showed that γ and β phases were formed. The presence of β phase may improve the coating lifetime [54] as this phase acts as the Al storage to produce the protective layer. The γ phase may enhance the high temperature performance of the coating due to its refrectory nature (high melting temperature). Meanwhile, the presence of other oxides such as NiO and NiAl2O4, which could affect the life span of TBCs by crack formation and delamination, was not detected. This is probably due to presence of β phase in a layer behind of TGO, which provides the required Al concentration for the formation and growing of the protective Al2O3 layer. As reported by Selezneff et al. [55], the minimum required Al concentration for the growth of the TGO layer before spallation is about 7 at%. Obviously,
ð1Þ
where M 1 and M 2 are the peak intensity of 111 plane (2θ ¼ 28:218)and ð111Þ plane (2θ ¼ 31:502) for the m-phase, respectively. I is the peak intensity of ð101Þ plane (2θ ¼ 30:271) for tetragonal-ZrO2 (t-zirconia). The volume fraction of the m-zirconia was obtained 95% and 60% for the APSed and SPSed specimens, respectively. The t-ZrO2 to mZrO2 phase transformation is a martensitic or a shear type transformation which is accompanied by a large volume expansion [56–59]. This phase transformation during cooling stage of the thermal cycling is destructive that can lead to spallation and cracking of the coating. As a result, the remained corrosion products impose stresses and intensify the coating damage. On the other hand, experimental results determined the formation of YVO4 crystals after the hot corrosion test. The possible reaction pathway is dissociation of Na2SO4 and formation of sodium metavanadate, which in reaction with YSZ forms YVO4 and monoclinic ZrO2 [41,46]. This reaction is probably caused by Y3 þ migration outward the coating surface by diffusion into the YSZ lattice structure. Therefore, it was expected that the dense and almost defectfree structure of the SPSed specimens resulted in better corrosion performance. However, experimental results revealed that the SPSed coating was also under severe corrosion attack after 12 h testing. This weakness may be attributed to the thickness of the TGO layer. As examined by Song et al. [5], the structure of SPS coating becomes close to the plasma-sprayed TBC system after exposure to a mild temperature for several hours. In general, the occurrnce of spalling mechanism in TBCs can be performed by the following three mechanisms [55]: (i) the thermal expansion mismatch between the ceramic and metallic layers, (ii) defects formation and surface deformation caused by the stress gradient during thermal cycling which lead to the coating delamination, (iii) the growth rate of the TGO layer which is the most important factor in the coating life span. Albite the lower fraction of defects and m-phase in the SPSed coating, it is therefore suggestible that the spalling is attributed to the excessive growth of TGO layer. Since the growth kinetics of TGO layer is an important factor, future
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Fig. 6. FE-SEM micrographs and EDS spectra of TBCs afte hot corrosion test. The EDS analyses were performed on the corrosion products, as indicated by points A and B. (a, b) APSed TBC; (c, d) SPSed TBC.
studies are required to optimize the thickness of this layer for enhanced hot corrosion resistance. 5. Conclusions A multilayered TBC on Inconel 738 superalloy was prepared by the SPS process at relatively low temperature of 1040 1C and applied pressure of 40 MPa. It was shown that as compared with conventional APS coatings, the multilayered structure was dense and almost free of defects and delaminated layers. The micro-hardness of the coating was also higher, particularly on the top ceramic layer. Microstructural studies showed that the grain structure of the zirconia layer was remained in nanometric region ( 60 nm), providing more hardness and corrosion resistance. Formation of γ (Ni) solid solution, β phase, and a thin TGO layer after SPS was also detected. Under hot corrosion test in a molten salt, the corrosion products was in flower-like structure with smaller feature sizes compared with the needle-like products for the APSed coating. The volume fraction of the formed m-phase was also lower. While about 95% of the t-zirconia was
transformed to the m-zirconia in the APSed specimen, this transformation was about 60% for the SPSed coating. Nevertheless, it was found that the corrosion resistance of the multilayered SPSed coating was not as high as it was expected. It is envisaged that the thickness of the TGO layer plays a critical role, which should be optimized in future work.
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CERAMICS INTERNATIONAL
Ceramics International 42 (2016) 2770–2779 www.elsevier.com/locate/ceramint
Spark plasma sintering of a multilayer thermal barrier coating on Inconel 738 superalloy: Microstructural development and hot corrosion behavior A.H. Paksereshta, A.H. Javadib, M. Bahramic, F. Khodabakhshid, A. Simchic,e,n a
Department of Ceramics, Materials and Energy Research Center, P.O. Box 31787-316, Karaj, Iran b Engineering Research Institute, Tehran, Iran c Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran d Department of Materials Science and Engineering, School of Engineering, Shiraz University, Zand Boulevard, Shiraz, Iran e Institute for Nanoscience and Nanotechnology, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran Received 20 August 2015; received in revised form 2 November 2015; accepted 2 November 2015 Available online 10 November 2015
Abstract In the present work, spark plasma sintering (SPS) process was employed to prepare a nanostructured yttria-stabilized zirconia (8YSZ) coating on a nickel-based superalloy (INCONEL 738) with functionally graded structure. A stack layer of INCONEL 738/NiCrAlY powder/Al foil/ NiCrAlY þYSZ powder/YSZ powder was SPSed in a graphite die at an applied pressure of 40 MPa under an vacuum atmosphere (8 Pa). The sintering temperature was 1040 1C. For comparison purpose, the air plasma spray (APS) technique was employed to prepare the thermal barrier coating (TBC). Microstructural studies by scanning electron microscopy showed that the SPSed coating was sound and free of interfacial cracks and large pores. Energy dispersive X-ray spectroscopy determined a limited inter-diffusion between the layers. A significant improvement ( 26%) in the Vickers micro-hardness was also measured. In order to examine the functionality of the coatings, hot corrosion test was performed at 1050 1C in a molten bath of Na2SO4 and V2O5 (45 to 55 weight ratio). The results showed that spallation occurred after 12 and 20 h for the SPSed and APS coatings, respectively. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Multilayer thermal barrier coating; Spark plasma sintering; Microstructure; Hot corrosion; Thermally grown oxide
1. Introduction Thermal barrier coatings (TBCs) are broadly utilized as a protection layer for the superalloys in order to reduce thermal oxidation and corrosion rate, for example, in gas engine turbines and aeronautic diesel engine chambers [1–4]. TBCs have a multi-layer composite structure which commonly consist of a superalloy substrate covered with one or more metallic bond coat (BC) followed by a thermally grown oxide (TGO) layer and finally a ceramic top layer [5–8]. MCrAlY (M ¼ Ni or Co or both) alloys or a platinum aluminized coating n Corresponding author at: Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Avenue, 14588 Tehran, Iran. Tel.: þ98 21 6616 5261; fax: þ98 21 6600 5717. E-mail address: [email protected] (A. Simchi).
http://dx.doi.org/10.1016/j.ceramint.2015.11.008 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
are typically used as BC to enhance the high temperature oxidation and corrosion [5,9,10]. The TGO layer operates as a diffusion barrier against hot oxidation which enhances the system thermal stability under service conditions [11]. Yttriastabilized zirconia (YSZ) is the most common ceramic material used in TBCs due to its low thermal conductivity that reduces the peak temperature for the internally cooled underlying substrates [6]. Recent studies have demonstrated the advantages of nano-sized YSZ for TBCs [12–14]. As compared with conventional YSZ coatings, a higher coefficient of thermal expansion, lower thermal diffusivity, higher hardness and toughness, and better wear resistance have been reported [12,15,16]. Physical vapor deposition (PVD) [17,18], low pressure plasma spray (LPPS) [19], vacuum/air plasma spray (VPS/APS) [20– 23], high velocity oxyfuel spray (HVOF) [24], and electrolytic
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deposition (ED) [25] have already been employed for the preparation of TBCs. Among these methods, APS is more practical and has been utilized broadly for different systems [26– 30]. Nevertheless, processing of nanostructured TBCs by APS is challenging due to a poor flowability of nanoparticles (which requires proper granulation) and grain growth caused by partial melting of the particles [16,31,32]. The deposited material by APS often displays a bi-modal grain structure induced by consolidation of melted and non-melted regions [33]. Recently, SPS has attracted great interest for the fabrication of functionally graded materials and nanostructured TBCs [1,34,35]. Jiang et al. [36] employed this process to prepare WC coating on tungsten substrates. Boidota et al. [1] and Song et al. [5] prepared TBCs on superalloys by SPS. The process has a number of advantages over the conventional methods such as a fast processing cycle (high heating rates with short dwell time), energy and material saving, limited grain growth and its efficiency to sinter materials with limited change of their initial microstructures [5,37]. Therefore, SPS is potentially very attractive for the fabrication of nanostructured coatings. Recently, multilayered or functionally-graded structures have been utilized to prepare TBCs [38–41]. Gradual compositional variations can reduce the residual and the thermal stresses between the top layer and BC, which improves the bonding strength [38–41]. However, fabrication of multilayered structures often requires multi-step processing methods. The aim of the present work is to prepare multilayered TBCs by SPS through a strightforward proces. To the best knowledge of the authors, no results on development of functionally-graded TBCs by SPS method have been reported so far; though FGM TBC is as a multi-layered structure with gradiently variation of YSZ volume fraction from the ceramicrich upper-layer toward the alloy-rich TBC bond coat interface. Nanostructured YSZ powder was utilized as the raw material to reduce the sintering temperature and to enhance the performance of the multi-layered TBC on INCONEL 738 (IN738). For comparison purpose, a TBC with the same composition and layer structure was fabricated by APS. Since hot oxidation and corrosion are the main phenomena restricting the application of TBCs in harsh atmospheres [42,43], hot corrosion resistance of the multilayered TBCs was studied in a molten V2O5/Na2SO4 salt. It is well known that the molten salt reacts with the top ceramic layer or even penetrates through its defects (pores, micro-cracks and columnar gaps) during thermal cycling under high temperature, which leads to the tetragonal to monoclinic phase transformation, delamination and spallation of the coating [8,42, 44–47].
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2. Experimental procedure 2.1. Spark plasma sintering of multilayered coating The characteristics of the materials used in the present work are shown in Table 1. Discs of Inconel 738 with a nominal chemical composition of Ni–15.48Cr–8.57Co–2.99Al– 2.12Mo–3.17Ti–2.99W–1.09Ta (wt%) in diameter of 30 mm and height of 5 mm were prepared by machining followed by surface grinding (2000 SiC papers) and acetone washing. To prepare the multilayered TBC by SPS, the superalloy substrate was inserted into a graphite die with an inner diameter of 30 mm. In order to prevent carbon diffusion and reaction, a thin layer of hex-BN powder (TERIO HBN, China) was deposited on the die surface by spraying. The NiCrAlY powder ( 1 g) was then poured on the surface of the substrate and leveled off by hand. A foil of commercially pure aluminum with a thickness of 4 μm was then placed on the top and a mixture of YSZ and NiCrAlY powder at a weight ratio of 1:1 was spread. The aim of this layer was to generate a thermally grown Al2O3 oxide layer to keep the interface flat and homogeneous, and to enhance the adhesion to the top layer. Finally, nanostructured 8YSZ (8 wt% Y2O3 þ 92 wt% ZrO2) powder was inserted on the top layer. A schematic imge of the multilayered TBC on IN738 superalloy is shown in Fig. 1a. As it is seen, the multilayered coating is composed of five layers: (i) Inconel substrate; (ii) a bond coat layer from NiCrAlY powders with a thickness of 100 μm; (iii) a thin aluminum layer (for TGO layer formation) with a thickness in the range of 6–8 μm; (iv) an intermediate functionally graded layer (50% 8YSZ and 50%NiCrAlY) with a thickness of 200 mm; (v) YSZ top layer with a thickness of 200 mm. It is pertinent to point out that this structure has been determined by various efforts to attain specimens without having defects such as unbounded layers, large cracks, and delamination. A commercial SPS apparatus (EF 20T-10, Easy Fashion Metals Products Co.) was utilized for the sample preparation. A schematic view of the machine operation is shown in Fig. 1b. The SPS route was carried out in a multi-stage thermal/pressure cycling under vacuum (8 Pa), as shown in Fig. 1c. In the first stage, pre-heating to 600 1C and holding for one minute was performed for thermal stabilization. A pressure of 10 MPa was applied at the start and then increased to 20 MPa at 600 1C. After that, the temperature raised to 1040 1C at a rate of 35 1C/min and the sample was held for 10 min. The applied pressure for sintering was 40 MPa. The specimens were then cooled to 600 1C within 10 min. Visual
Table 1 Characteristics of materials used in this study. Materials
Commercial name
Chemistry (in wt%)
Form
Fabrication route
Nickel base superalloy NiCrAlY Yettria-stabilized zirconia Aluminum foil
INCONEL 738 Metco 442 8YSZ (US3630) 1050Al
Ni–15.48Cr–8.57Co–2.99Al-–2.12Mo–3.17Ti–2.99W–1.09Ta Ni–8.5Cr–7Al–5Mo–2Si–2B–2Fe–3Al2O3 92ZrO2–8Y2O3 (high purity499.9%) Al–0.05Cu–0.05Mg–0.25Si–0.4Fe–0.05Mn–0.07Zn–0.05Ti
Plate (substrate) Powder ( 70 μm) Powder ( 40 nm) Sheet
Vacuum melting Mechanically clad Agglomerated/sintered Rolling/annealing
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Fig. 1. (a) A schematic representation of multilayer TBC on Incolnel 738 superalloy. (b) Schematic of spark plasma sintering accomplishment. (c) Termal/pressure cycling used for SPS of the multilayered TBC.
and 325 US mesh) to attin granulates with 44–90 mm sizes. Finally, heat treatmrnt was performed at 1100 1C for 2 h.
Table 2 Processing parameters used for APS of different layers. Parameter
Current (A) Voltage (V) Primary gas, Ar (l/min) Secondary gas, H2 (l/min) Powder feed rate (g/min) Spraying distance (cm)
Different layers
2.3. Microstructural evaluation and hardness measureemt
NiCrAlY
Graded layer
YSZ
450 75 55 17 35 11
550 70 36 17 30 8
550 70 36 17 30 8
examination revealed that the samples are sound but in a dark color (Fig. 1d). It seemed that surface carbon soating and/or carbonization of the substrate was taken place, albeit a barrier coating (hex-BN) was utilized. Therefore, the SPSed samples were heat treated in an electrical furnace at 1000 1C for 2 h. The mentioned time for post annealing phemomenon was selected based on previous studies [1,5], which aimed to eliminat carborization as well as detecting the γ and β phases.
For microstructural studies, the multilayered samples were cross-sectioned by a diamond cutting machine (DCS 150, Synova). The specimens were then cold mounted in an epoxy resin and grined by SiC emergy papers down to 2000 mesh. Diamond pastes (10, 3, and 1 μm sizes) were used for mirror polishing. A field emission-scanning electron microscope (FESEM, VEGA//TESCAN-XMU, Russia) equipped by an energy dispersive X-ray spectroscopy (EDS) was employed for microstructural examines. Possible phase transformations were analyzed by using an X-ray diffractometer (XRD: GNR, Explore, Analytical Instruments Group, Netherlands) with Cu Kα radiation. Variations in the micro-hardness profile along the cross-section were determined by a Vickers indenter (MVK-H21, AKASHI, Japan) at an applied load of 50 g for a dwell time of 15 s. The indentation testing was repeated ten times for each region and the average value was reported with standard deviation (SD).
2.2. Air plasma spraying
2.4. Hot corrosion test
Plasma spraying was performed with a Sulzer-Metco F4MB plasma gun (Sulzer-Metco, Switzerland) in air. Argon was used as a primary and carrier gas and hydrogen was utillized as a secondary gas. The processing parameters for different layers are shown in Table 2. To improve the flowability of thefeed stock powders, the YSZ particles were granulated by utilizing polyvinyl alcohol (Junsel Chemical Co. Ltd., Japan) as a binder [48]. The aggromerates were screened by sieving (180
A batch of molten V2O5/Na2SO4 salt (55/45 wt%) was prepared by mixing of the constituents at room temperature followed by heating at 1050 1C. The corrosive salt was then spread over the TBC in a concentration of 30 mg/cm2. To avoid the edge effect [42,44,45,49], no covering was considered near the edges for apart 3 mm. The prepared setup was placed in an electric furnace with a fixed temperature of 1050 1C in atmospheric condition. Every four hours, the
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Fig. 2. A cross-sectional FE-SEM image combined with elemental EDS maping for the multilayer TBC produced by SPS followed by post-heating at 1000 ºC for 2 h.
specimens were taken out from the furnace and analyzed for possible cracking, spallation, and severe formation of corrosion products. 3. Results 3.1. Microstructure A typical cross-sectional FE-SEM image of the multilayered coating combined with EDS elemental mapping of layers is shown in Fig. 2. The specimens consist of five layers of IN738/NiCrAlY/Al/NiCrAlY þ 8YSZ/8YSZ. The layers are dense and free of common defects like pores and microcracks. Formation of a TGO layer with a mixed composition, mainly alumina (Al2O3) with some zirconia (ZrO2) inclusions,
was also detected. To determine the composition of the bond layer, EDS analysis of selected regions (A, B and C in Fig. 2) was performed. The results are shown in Fig. 3. The formation of γ(Ni) solid solution with dark gray color and β(NiAl) phase with bright gray color was noticed Typical cross-sectional FE-SEM image of the TBC coating prepared by APS is shown in Fig. 4a. Unlike the SPSed coating, non-bonded interface, voids, micro-cracks and intersplats boundaries were observed. These defects were partly formed due to presence of un-melted/semi-melted particles, weak adhesion of splats to the underlayers, evaporation of condensate, loss of solid particles and also a pull-out of particles in sample preparation [50,51]. To better show the microstructure of the prepared TBCs, the YSZ layer was broken and the resulting cross-section was observed by FE-SEM. Fig. 4b shows a typical
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Fig. 3. Results of EDS analysis performed on locations (a) A, (b) B, and (c) C shown in Fig. 2.
fracture surface of the APS sample (near the top YSZ layer region). A lamellar and bi-modal structure with micro- and submicro-sized grains (melted and non-melted regions) is seen. It was concluded that APS changed the initial nanostructure of the particles to a coarse structure with pore formation. This microstructural changes occurred due to high processing temperature, partial or complete melting of particles, and formation of individual splats during solidification of molten droplet due to collision. Moreover, some degrees of cross-sectional inhomogeneities are noticeable, which are attributed to the functionally graded structures of the APSed and SPSed TBCs. Fig. 4c shows the fracture surface of the SPSed specimen. A dense structure
with nanometric grains (average size of about 60 nm) is visible. The short sintering cycle at relatively low temperature in the presence of applied pressure inhibited significant grain growth while preceding rapid densification. As it will be shown in following, the presence of pores and defects in the top coating layer facilitates penetration of the molten salt into the underlayers, degrading the hot corrosion resistance. 3.2. Mechanical properties Variations of micro-hardness along the cross-section of the prepared specimens are shown in Fig. 4d. The results indicate a
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Fig. 4. (a) Cross-sectional FE-SEM image of the multilayer TBC prepared by APS. Cross-sectional FE-SEM images of the deposited YSZ by (b) APS and (b) SPS after fracturing. (d) Mean values of Vickers micro-hardness across thickness section for the prepared TBCs.
relatively smooth increase in the hardness from the superalloy substrate to the ceramic layer. Interestingly, the hardness of the multilayered coating prepared by the SPS is higher than the APS one, particularly in the top ceramic layer. While the SPSed intermediate layers exhibited a small improvement in the mean hardness value (about 10%) compared with the APSed ones, the hardness of the top coating significantly increased from 620 HV to 780 HV (26% enhancement). The capacity of the SPS process in full densification while retaining nanostructured grains improved the indentation resistance of the coatings. 3.3. Hot corrosion behavior The hot corrosion response of the TBCs was evaluated by V2O5/Na2SO4 melt reaction at 1050 1C. The corrosion products were analyzed by XRD. Fig. 5a shows XRD patterns of
the top coatings prepared by APS and SPS. Both coatings have a tetragonal zirconia crystallographic structure. The pattern of the SPSed specimen exhibits peak-line broadening which is an indicator of its finer grain structure [52,53]. XRD patterns of these samples after hot corrosion test are shown in Fig. 5b. A mixture of monoclinic (m-ZrO2) and tetragonal (t-ZrO2) zirconia together with corrosion product (YVO4) were noticed. Results of microstructural studies after hot corrosion test are shown in Fig. 6. It appears that elongated phases are formed due to reactions at the elevated temperature (Figs. 6a and 6b). EDS analysis support the XRD results, indicating that the main product is yttrium-vanadium oxide (Figs. 6c and 6d). Large and needle-like crystals (with a length of 70 mm and larger) were detected for the APS coating while smaller flower-like crystals (with a length of 30 mm and smaller) were observed for the SPS layer. Both XRD analysis (peak intensity) and
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studying of the phase transformation before and after air annealing is useful to get more insight on the possible mechanisms of TGO layer. Many studies have already been performed on this issue, for example [1,5,55], and it is determined that reactions of the aluminum folil with both oxidizing atmosphere and zirconia are feasible. It is well known that hot corrosion resistance of TBCs strongly depends on their chemical compostion, grain size, and structure homogeneity [56]. Pores and micro-cracks are suitable sites for molten salt penetration through the coating during hot corrosion testing. XRD analysis showed that after hot corrosion test, monoclinic zirconia (m-phase) was formed. To determine the amount of m-phase in the coatings, the following equation was utilized [37–40,48]: %M ¼ ðM 1 þ M 2 Þ=ðM 1 þ M 2 þ I Þ
Fig. 5. XRD patterns of the prepared TBCs (a) before and (b) after hot corrosion test.
microstructural examinations revealed that the volume of the corrosion product on the SPS specimen was lower than that of APS sample.
4. Discussion The formation of TGO layer during SPS is a results of reaction between the aluminum thin foil and ZrO2 through oxygen diffusion from the zironia ceramic towards the interface and/or trapped oxygen in the sintering chamber, which can form oxidizing atomosphere with the graphite die (CO/ CO2) gas mixture [5]. Microstructural analysis showed that γ and β phases were formed. The presence of β phase may improve the coating lifetime [54] as this phase acts as the Al storage to produce the protective layer. The γ phase may enhance the high temperature performance of the coating due to its refrectory nature (high melting temperature). Meanwhile, the presence of other oxides such as NiO and NiAl2O4, which could affect the life span of TBCs by crack formation and delamination, was not detected. This is probably due to presence of β phase in a layer behind of TGO, which provides the required Al concentration for the formation and growing of the protective Al2O3 layer. As reported by Selezneff et al. [55], the minimum required Al concentration for the growth of the TGO layer before spallation is about 7 at%. Obviously,
ð1Þ
where M 1 and M 2 are the peak intensity of 111 plane (2θ ¼ 28:218)and ð111Þ plane (2θ ¼ 31:502) for the m-phase, respectively. I is the peak intensity of ð101Þ plane (2θ ¼ 30:271) for tetragonal-ZrO2 (t-zirconia). The volume fraction of the m-zirconia was obtained 95% and 60% for the APSed and SPSed specimens, respectively. The t-ZrO2 to mZrO2 phase transformation is a martensitic or a shear type transformation which is accompanied by a large volume expansion [56–59]. This phase transformation during cooling stage of the thermal cycling is destructive that can lead to spallation and cracking of the coating. As a result, the remained corrosion products impose stresses and intensify the coating damage. On the other hand, experimental results determined the formation of YVO4 crystals after the hot corrosion test. The possible reaction pathway is dissociation of Na2SO4 and formation of sodium metavanadate, which in reaction with YSZ forms YVO4 and monoclinic ZrO2 [41,46]. This reaction is probably caused by Y3 þ migration outward the coating surface by diffusion into the YSZ lattice structure. Therefore, it was expected that the dense and almost defectfree structure of the SPSed specimens resulted in better corrosion performance. However, experimental results revealed that the SPSed coating was also under severe corrosion attack after 12 h testing. This weakness may be attributed to the thickness of the TGO layer. As examined by Song et al. [5], the structure of SPS coating becomes close to the plasma-sprayed TBC system after exposure to a mild temperature for several hours. In general, the occurrnce of spalling mechanism in TBCs can be performed by the following three mechanisms [55]: (i) the thermal expansion mismatch between the ceramic and metallic layers, (ii) defects formation and surface deformation caused by the stress gradient during thermal cycling which lead to the coating delamination, (iii) the growth rate of the TGO layer which is the most important factor in the coating life span. Albite the lower fraction of defects and m-phase in the SPSed coating, it is therefore suggestible that the spalling is attributed to the excessive growth of TGO layer. Since the growth kinetics of TGO layer is an important factor, future
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Fig. 6. FE-SEM micrographs and EDS spectra of TBCs afte hot corrosion test. The EDS analyses were performed on the corrosion products, as indicated by points A and B. (a, b) APSed TBC; (c, d) SPSed TBC.
studies are required to optimize the thickness of this layer for enhanced hot corrosion resistance. 5. Conclusions A multilayered TBC on Inconel 738 superalloy was prepared by the SPS process at relatively low temperature of 1040 1C and applied pressure of 40 MPa. It was shown that as compared with conventional APS coatings, the multilayered structure was dense and almost free of defects and delaminated layers. The micro-hardness of the coating was also higher, particularly on the top ceramic layer. Microstructural studies showed that the grain structure of the zirconia layer was remained in nanometric region ( 60 nm), providing more hardness and corrosion resistance. Formation of γ (Ni) solid solution, β phase, and a thin TGO layer after SPS was also detected. Under hot corrosion test in a molten salt, the corrosion products was in flower-like structure with smaller feature sizes compared with the needle-like products for the APSed coating. The volume fraction of the formed m-phase was also lower. While about 95% of the t-zirconia was
transformed to the m-zirconia in the APSed specimen, this transformation was about 60% for the SPSed coating. Nevertheless, it was found that the corrosion resistance of the multilayered SPSed coating was not as high as it was expected. It is envisaged that the thickness of the TGO layer plays a critical role, which should be optimized in future work.
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