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Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys
Journal Pre-proof Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys Bin Yin, Guang Xie, Langhong Lou, Jian Zhang PII:
S0925-8388(20)30803-3
DOI:
https://doi.org/10.1016/j.jallcom.2020.154440
Reference:
JALCOM 154440
To appear in:
Journal of Alloys and Compounds
Received Date: 15 October 2019 Revised Date:
6 February 2020
Accepted Date: 19 February 2020
Please cite this article as: B. Yin, G. Xie, L. Lou, J. Zhang, Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154440. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys Bin Yina,b, Guang Xiea,*, Langhong Loua, Jian Zhanga a. Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China b. School of Materials Science and Engineering, University of Science and Technology of China, Hefei 230026, China
Abstract The effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys at 1000 oC and 1100 oC has been investigated. Interdiffusion occurs between the NiCrAlYSi coating and the substrate, leading to the formation of inter-diffusion zone (IDZ) and secondary reaction zone (SRZ). After exposure at 1000 o
C for 100h, the initial γ/γ' transforms to γ' phases in the IDZ, and SRZ with σ and µ
topologically close-packed (TCP) phases is formed primarily due to the inward diffusion of Al. Different behavior of SRZ formation in the alloys with various Ta content can be observed after exposure at 1000 oC, which highlights the potency of Ta addition in promoting SRZ formation. This may be ascribed to the fact that Ta promotes the precipitation of TCP phases. After exposure at 1100 oC for 100h, the initial γ/γ' transforms to γ layer and coarse γ' phases in the IDZ, and no SRZ can be found. The formation of IDZ is significantly suppressed by Ta addition, which may be related to the retardation of interdiffusion between the coating and the substrate due to the increased γ' volume fraction. Furthermore, γ' rafting parallel to and/or 1
perpendicular to the coating/substrate interface in the substrate beneath IDZ are found in the alloys after exposure at 1000 oC and 1100 oC. Microstructural observation shows that Ta has a potent influence on the development of perpendicular γ' rafting. The mechanism of parallel γ' rafting and perpendicular γ' rafting has been discussed, respectively. Key words: Single crystal superalloys, NiCrAlYSi coating, Ta, Secondary reaction zone (SRZ), Rafting 1. Introduction Ni-based single crystal superalloys are widely used in industrial gas turbines, aircrafts and jet engines due to their excellent mechanical strength at elevated temperature. Usually, MCrAlY (M=Ni and/or Co) coatings both as stand-alone protective overlay coating and as a bond coat for thermal barrier coatings (TBCs) are widely applied on the superalloys to protect them from oxidation and corrosion attack [1-3]. However, interdiffusion between the coating and the underlying substrate inevitably occurs during thermal exposure at high temperature. Then Al in the coating is consumed, which accelerates the degradation of the coating. Moreover, microstructure degradation of the underneath substrate occurs as a result of local chemistry changes in the substrate [4]. On one hand, it has been extensively reported that inter-diffusion zone (IDZ) and even secondary reaction zone (SRZ) form beneath the coating/substrate interface, which causes considerable deterioration of mechanical properties in the alloy substrate [5-8]. SRZ located beneath the IDZ is comprised of a 2
three-phase transformation product involving γ, γ' and TCP phases [5]. The formation of SRZ is supposed to be dominantly controlled by the inward diffusion of aluminum from the coating layer to the substrate [7, 8]. Moreover, the morphology of SRZ strongly depends on Ni-base superalloy composition [9-11]. On the other hand, oriented γ' rafting in the region beneath the coating occurs during thermal exposure and the rafts are parallel to the coating/substrate interface, which strongly influences the deformation and failure of materials [12-16]. Furthermore, the composition of alloy has a potent influence on the parallel rafting [15]. However, to the knowledge of the authors, the γ' rafting perpendicular to the coating/substrate interface has not been reported. Ta is one of the major elements of Ni-base superalloys, which is benefitial to improve creep performance. However, few works have focused on the effect of Ta in the substrate on the microstructure degradation of the coated substrate. It has shown that Ta addition in the alloy significantly retards the propagation of the SRZ (consists of γ and β-NiAl with no TCP phases) during high-temperature exposure [17]. It is important for better alloy and coating design to reveal the effect of Ta on IDZ, SRZ (consists of γ, γ' and TCP phases) formation and γ' rafting of the coated superalloys. In this paper, the influence of Ta on the microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys associated with IDZ, SRZ formation and γ' rafting was investigated. In addition, the mechanisms of the above microstructural degradations were also discussed. 2. Experimental procedures 3
The three experimental single crystal (SX) superalloys, namely, 2Ta, 6Ta, and 8Ta, with nominal composition listed in Table 1 was directionally solidified into SX bars (16 mm in diameter and 220 mm in length) in a Bridgman furnace. As-cast SX alloy rods were solution heat treated followed by aging to develop the γ/γ' structure. Samples with the size of Φ15 × 2 mm were sliced from the heat-treated rods by using an electric discharge machine. The samples were ground with 800-mesh SiC paper, grit blasted in a wet atmosphere, and then ultrasonically cleaned in acetone. The NiCrAlYSi coating with nominal compositions Ni-18.1Cr-11.5Al-0.6Y-0.9Si (wt. %) were deposited on samples by arc ion plating. The thickness of the coating is about 40 µm. The coated samples were vacuum annealed in a sealed silica tube at 870 oC for 3.5 h. After that, isothermal oxidation of the coated samples was carried out in a muffle furnace at 1000 ℃ and 1100 ℃ for 1000 h respectively in static air. At intervals of 100, 300 and 500 h, some samples were taken out for microstructure analysis. The cross sectional microstructures of the samples were examined by scanning electron microscopy (SEM, Philips FEI-Inspet F) equipped with energy dispersive spectroscopy (EDS, Oxford X-man). An etchant (40 g CuSO4, 100ml HCl, 200 ml H2O) was used to reveal the microstructure. Compositional information and elemental mapping were obtained using a Cameca SX100 electron probe microanalyzer (EPMA, Shimadzu EPMA1610) operating at 20 kV. Besides, the JEOL JEM-2100 (200 kV) transmission electron microscope (TEM) was used to determine and analyze the phases. 3. Results 4
3.1. Microstructure of the as-coated alloys SEM micrographs of the cross-sectional as-coated samples (after annealed in a sealed silica tube at 870 oC for 3.5 h) are shown in Fig. 1. An interdiffusion zone (IDZ) formed beneath the coating/substrate interface in 2Ta, 6Ta and 8Ta alloys, which was mainly composed of γ' phases and Cr-rich precipitates. The IDZ thicknesses of 2Ta, 6Ta and 8Ta alloy are about 5.4±0.6, 3.7±0.2 and 3.1±0.3µm, respectively. The γ' precipitates were found to be regular and cuboidal beneath the IDZ in 2Ta, 6Ta and 8Ta alloys. 3.2. Interdiffusion and microstructural evolution at 1000 oC The cross-sectional morphologies of coated alloys after oxidation at 1000 oC for 100 h are displayed in Figs. 2-4. Further diffusion had occurred. The initial γ/γ' had transformed to γ' phases and rod-like and granular carbides were precipitated in the IDZ. The thicknesses of IDZ in 2Ta, 6Ta and 8Ta alloys are about 21±1.0, 9±0.8, 13 ±1.7 µm, respectively. Beneath the IDZ, secondary reaction zone (SRZ) with topologically close-packed (TCP) phases was observed. However, the morphology of SRZ varied greatly for three alloys. Almost no SRZ could be found in 2Ta alloy (Fig. 2), continuous SRZ formed in 6Ta and 8Ta alloys with the depth of 16±1.5 and 21± 3.5 µm, respectively (Figs. 3 and 4). It indicates that higher content of Ta in substrate encourages more SRZ formation. The rod-like and needle-like TCP phases were present in the SRZ of 6Ta and 8Ta alloys. The compositions of carbides and TCP phases are shown in Table 2. The Cr(Re)-rich carbide precipitated in the IDZ was determined as M23C6 by TEM (Figs. 5
5a, 5d). The rod-like TCP phase enriched in Cr, W and Re was determined as σ phase (Figs. 5b, 5e). While the needle-like TCP phase enriched in W and Re was determined as µ phase (Figs. 5c, 5f). Apart from the above microstructure changes observed in IDZ and SRZ, the bulk γ'/γ substrate degraded due to the interdiffusion during thermal exposure. Degradation of substrate associated with γ' rafting parallel to the coating/substrate interface was observed in 2Ta, 6Ta and 8Ta alloys (Figs. 2c, 3c and 4c), which was consistent with previous studies [14-16]. However, γ' rafting perpendicular to the coating/substrate interface beneath parallel rafts was also observed in 6Ta alloy (Fig. 3d). Fig. 6 shows the TEM micrograph of the parallel rafts and perpendicular rafts in 6Ta alloy. Plenty of dislocations were observed in the γ matrix of the parallel rafts region (Fig. 6a). While dislocations were not observed in the γ matrix of the perpendicular rafts region (Fig. 6b). The SEM with EDS was used to measure the chemistry profiles, and the experiment was carried out from the coating to the substrate in a step of 10 µm. The typical elemental concentration profiles with the depth of 6Ta alloy after thermal exposure at 1000 oC for 100 h are shown in Fig. 7. Chemical gradient of Al, Re and W was present beneath the IDZ due to the interdiffusion between the coating and the substrate. EPMA element maps of 2Ta, 6Ta and 8Ta alloys after exposure at 1000 ◦C for 100 h are shown in Figs. 8-10, respectively. Due to the chemical gradient between NiCrAlYSi coating and substrate, Ni, Cr and Al diffused from the coating into the substrate, while W, Mo, Ta, Co, Re and Ti were along the opposite direction. As 6
illustrated in Figs. 8-10, the concentration of Al and Ni in the IDZ is much higher than that in the substrate. A thin layer of Ti-rich phase formed between the coating and the substrate. The thickness of Ti-rich layer decreased as the content of Ta in the substrate increased. Granular carbides distributed along the coating/substrate interface were observed. The cross-sectional morphologies of 2Ta, 6Ta and 8Ta alloys after oxidation at 1000 oC for 1000 h are displayed in Fig. 11. The IDZ and SRZ significantly grew. The thicknesses of IDZ+SRZ in 2Ta, 6Ta and 8Ta alloys are about 42±3.3, 66±4.6 and 73±6.5µm, respectively. A significant amount of TCP phases are precipitated in the dendrite of 8Ta alloy, which in turn promotes the formation of SRZ. The parallel γ' rafting almost disappeared, and perpendicular γ' rafting was observed in 6Ta and 8Ta alloys. 3.3. Interdiffusion and microstructural evolution at 1100 oC The cross-sectional morphologies of coated alloys after exposure at 1100 oC for 100 h are displayed in Figs. 12-14. The IDZ was composed of a thin layer of γ phase beneath the coating/substrate interface and some coarse γ' phases located beneath the γ layer. These coarse γ' phases were dissolved during etching and only voids have left (Figs. 12b, 13b, 14b). Moreover, it is noted that the thickness of γ layer decreased and the content of coarse γ' phases increased with increasing the content of Ta. No SRZ was observed in three alloys. The depth of IDZ in 2Ta, 6Ta and 8Ta alloys was about 19±1.7, 16±2.1 and 12±1.1 µm, respectively. Ta was shown to suppress the formation of IDZ. 7
Degradation of the substrate associated with γ' rafting parallel to the coating/substrate interface was observed beneath the IDZ (Figs. 12c, 13c, 14c). Moreover, the perpendicular γ' rafting were present in both 6Ta and 8Ta alloys (Figs. 13d, 14d). EPMA element maps after exposure at 1100 oC for 100 h are shown in Figs. 15-17. Ni and Cr diffused from the coating into the substrate, while W, Mo, Ta, Co, Re and Ti were along the opposite direction. However, the inward diffusion of Al hardly occurred, due to the fact that the concentration of Al in the coating was not higher than that in the substrate after oxidation consumption (Figs. 15-17). In the 6Ta and 8Ta alloys, some granular bright phases were present in the substrate beneath the IDZ. According to the EPMA results shown in Figs. 16 and 17, the granular phase was Ti, Ta-rich carbide. A thin layer of Ti-rich phase was also found between the coating and substrate. Moreover, the thickness of Ti-rich layer decreased with the increase of Ta content in the substrate. The cross-sectional morphologies of 2Ta, 6Ta and 8Ta alloys after oxidation at 1100 oC for 1000 h are displayed in Fig. 18. The IDZ significantly grew. The thicknesses of IDZ in 2Ta, 6Ta and 8Ta alloys are about 73±6.8, 30±1.3 and 23±2.1 µm, respectively. The parallel γ' rafting almost disappeared, and perpendicular γ' rafting was observed in 6Ta and 8Ta alloys. The 8Ta alloy showed poor microstructure stability and TCP phases precipitated in the substrate after thermal exposure. However, the TCP phases disappeared as the interdiffusion between the coating and the substrate further occurred (Fig. 19). The depth of TCP-free region 8
significantly increased with increasing the exposure time. 4. Discussions 4.1.
Interdiffusion
and
microstructural
evolution
behavior
near
the
coating/substrate interface The formation of IDZ and SRZ beneath coating has been ascribed to the interdiffusion of elements between the coating and the substrate [5-8]. After exposure at 1000 oC, the content of aluminum is significantly increased in the IDZ due to the inward diffusion of Al (Figs. 8-10). This leads to the transformation from γ/γ' to γ' phases in the substrate adjacent to the coating and the resulting precipitation of carbides to accommodate excess chromium and rhenium (Fig. 2). As a forming element of γ' phase, Ta significantly promotes the above transformation. However, the growth of IDZ has been suppressed by the formation of SRZ in 6Ta and 8Ta alloys. Therefore, the depth of IDZ in 2Ta alloy without SRZ is the largest in the three alloys, while the depth of IDZ in 8Ta alloy is larger than that in 6Ta alloy. It has been suggested that SRZ formation was driven by the outward diffusion of Ni from the substrate into the coating [18]. In the present study, however, the diffusion of Ni is along the opposite direction (Figs. 8-10). It is likely ascribed to that the concentration of Ni in the NiCrAlYSi coating (68.9 wt. %) is significantly higher than that in 2Ta (65.8 wt. %), 6Ta (61.8 wt. %) and 8Ta (59.8 wt. %) alloys. Therefore, it has been suggested that the inward of Al resulted in the SRZ formation after exposure at 1000 oC [7, 8]. Two types of TCP phases are precipitated in the SRZ of 6Ta and 8Ta alloys. The σ phase enriched in Cr and Re prefers to precipitate nearby 9
the IDZ probably due to the enrichment of Cr in this region. While, the W, Re-rich µ phase is distributed nearby the substrate probably because of high content of W. Different SRZ formation behavior in 2Ta, 6Ta and 8Ta alloys is observed after exposure at 1000 oC, which highlights the potency of Ta addition in promoting SRZ formation (Figs. 2-4). Generally, the formation of TCP phases in superalloys can be predicted by the average electron-hole concentration (Nv) in the substrate [19]. As a refractory element, Ta significantly contributes to the Nv and promotes the precipitation of TCP phases. This is in good agreement with the difference in the microstructure instability of substrate far away from the coating /substrate interface in 2Ta, 6Ta and 8Ta alloys (Figs. 2-4). Therefore, Ta promotes the precipitation of TCP phase and leads to more SRZ formation. After exposure at 1100 oC, the consumption of Al in the coating increased due to higher temperature oxidation, resulting in less inward diffusion of Al. In contrast, the inward diffusion of Ni accelerates with increasing the exposure temperature. Therefore, a layer of γ phase formed beneath the coating/substrate interface. The elemental diffusivities in the ordered γ'-Ni3Al phase are generally 1-2 orders of magnitude lower than that in γ-Ni [21]. Therefore, the formation of γ layer beneath coating/substrate interface promotes the diffusion of Re and W from the substrate to the coating, leading to the depletion of W and Re in the section beneath the γ layer, which promotes the growth of γ' phases. Therefore, the IDZ is composed of γ layer and coarse γ' phases. Ta is shown to significantly suppress the IDZ formation after exposure at 1100 10
o
C (Figs. 12-14). It has been reported that Ta decreases the inter-diffusivity between γ
and γ' phases and improves the stability of γ' phase [17, 20]. So Ta addition in the substrate suppresses transformation from γ/γ' to γ phase layer, and retards the growth of IDZ. Moreover, Ta significantly increases the volume fractions of γ' phases, which impedes the interdiffusion between the coating and the substrate [17]. Therefore, the growth of IDZ is further suppressed. Moreover, the thickness of Ti-rich layer decreases with increasing Ta content (Figs. 15-17). No SRZ is observed in 2Ta, 6Ta and 8Ta alloys after exposure at 1100 oC, which is ascribed to the inward diffusion of Ni and outward diffusion of Re and W. The TCP phases precipitated in the substrate of 8Ta alloy disappear due to the outward diffusions of Re and W further proceeding (Fig. 19). However, the inward diffusion of Cr accelerates, leading to the precipitation of Cr-rich carbide beneath the IDZ in 6Ta and 8Ta alloys after exposure at 1100 oC for 100 h (Figs. 13 and 14). 4.2. Degradation of substrate during interdiffusion In addition to the microstructure changes observed in the IDZ and SRZ, the bulk γ/γ' substrate degrades after exposure at 1000 oC and 1100 oC. Parallel γ' rafting and perpendicular γ' rafting in the substrates beneath IDZ were observed. The γ' rafting parallel to the coating/substrate interface is caused by the externally applied stresses as well as the positive or negative sense of the misfit [14-16]. Furthermore, the compositions of substrate has a potent influence on the parallel γ' rafting, due to the probably change of γ/γ' misfit parameter during the interdiffusion between the coating and substrate [15]. However, in the present study, 11
parallel γ' rafting occurs in the 2Ta, 6Ta and 8Ta alloys both at 1000 oC and 1100 oC. Ta addition has little effect on the parallel γ' rafting. The parallel γ' rafting may be primarily ascribed to the strain induced by grit blasting. The threshold strain for rafting to continue in the absence of applied stress is as small as 0.1%, as reported by Reed et al. [22]. Therefore, the strain induced by grit blasting may provide the driving force for parallel γ' rafting of 2Ta, 6Ta and 8Ta alloys. As shown in Fig. 6a, Plenty of dislocations are distributed in the γ matrix of parallel rafts region, which is similar with that in the deformation-induced rafting during creep [23]. However, the perpendicular γ' rafting beneath the IDZ has not been reported. According to the results shown in Figs. 2-4 and 12-14, Ta has a potent influence on the perpendicular rafting. Few dislocations are present near the perpendicular γ' rafts (Fig. 6b). The formation mechanism of perpendicular γ' rafting may be different from that of parallel γ' rafting. The γ' rafting during load-free annealing has been reported in materials like AM1 and CMSX-4 [24-26]. It was suggested that the local chemical heterogeneities would lead to corresponding chemical gradients on the scale of the dendrite structure, which may play a crucial role in orientating the γ' coarsening. In the present study, the mechanism of perpendicular rafting can be schematically summarized as shown in Fig. 20. With no chemical gradients, the γ′ may grow by equally moving any of the three equivalent {001} γ/γ′ interfaces, and the γ and γ′ forming elements diffuse and exchange (Fig. 20a). Due to the interdiffusion between the coating and substrate, the chemical gradients of Al, Re and W are formed in the substrate, which may lead to the 12
directional movement of the corresponding elements into the same directions as the gradients. More alloying elements for phase transformation would diffuse in the perpendicular direction (Fig. 20b). The γ′ phase is prone to coalesce vertically. The three experimental alloys display negative misfit at high temperatures. It is generally known that Al with low solubility in the γ phase is mainly segregated in the γ' phase. The lattice parameters of γ decreases due to the outward diffusion of Re and W, leading to the relax of misfit stresses in the vertical matrix channels. While, the misfit stresses in the horizontal matrix channels persist relatively (Fig. 20b). It has been reported that the difference in misfit stresses on the horizontal faces and vertical faces of γ' phase, causing by the formation of horizontal dislocation networks, results in the γ' rafting [23, 27]. Thus, this promotes further vertical growth of γ' phase. Therefore, the γ′ is prone to coalesce vertically rather than laterally, which results in the formation of perpendicular rafting in 6Ta alloy after exposure at both 1000 oC and 1100 oC (Fig. 20c). Furthermore, the magnitude of the misfit is also found to affect the rate at which rafts develops [23, 25]. Ta has been shown to increases the lattice parameters of both γ' and γ phases, but increase the lattice misfit [28]. The misfit stresses increase with increasing Ta addition, resulting in the increase of driving force for rafting. In the 2Ta alloy, the misfit stresses are not enough for perpendicular rafting at both 1000 oC and 1100 oC. Reed has reported that higher γ' volume fraction with lower γ channel width restricting dislocation movement, resulting in greater resistance to rafting [25]. There is no perpendicular rafts in the 8Ta alloy after exposure at 1000 oC for 100h, although 13
the misfit of 8Ta alloy is larger. When the exposure time prolongs to 1000h or the exposure temperature increases to 1100 oC, the interdiffusion accelerates and the chemical gradient of Al, Re and W beneath the IDZ significantly increases, which promotes the perpendicular rafting. Moreover, the misfit stresses increased with increasing temperature, resulting in the increase of driving force for rafting. Therefore, perpendicular rafting in 8Ta alloy occurs after exposure at 1000 oC for 1000h and at 1100 oC. 5. Conclusions In summary, the following conclusions can be drawn from this work: 1. After exposure at 1000 oC, the initial γ/γ' transformed to γ' and Cr, Re-rich M23C6 carbides in the IDZ of 2Ta, 6Ta and 8Ta alloys. Moreover, the SRZ with σ and µ topologically close-packed (TCP) phase formed in the 6Ta and 8Ta alloys. It indicates that Ta addition promotes the formation of SRZ, which may be ascribed to the fact that Ta promotes the precipitation of TCP. 2. After exposure at 1100 oC, the initial γ/γ' transformed to a γ layer and some coarse γ' phases in the IDZ. No SRZ was present in the 2Ta, 6Ta and 8Ta alloys. Ta was shown to suppress the growth of IDZ, which may be related with the increase of γ' volume fraction, retarding the interdiffusion between the coating and the substrate. 3. The γ' rafts parallel to the coating/substrate interface in the substrates beneath IDZ were found in the 2Ta, 6Ta and 8Ta alloys both at 1000 oC and 1100 oC. While the perpendicular rafts were present in the 6Ta alloy after exposure at both 1000 oC and 1100 oC and 8Ta alloy after exposure at 1000 oC for 1000h and at 1100 oC. It is 14
believed that the parallel rafting can be attributed to the strain induced by grit blasting. While, the perpendicular rafting are probably ascribed to the high chemical gradient of Al, Re and W caused by the interdiffusion between the coating and the substrate. Acknowledgements This work was financially supported by the National Natural Science Foundation of China under grant No. 51771204, No. U1732131, No. 51631008, No. 91860201 and No. 51911530154.The authors are grateful for these supports. References [1] G.W. Goward, Progress in Coatings for Gas Turbine Airfoils, Surf. Coat. Technol. 108 (1998) 73-79. [2] J. Ma, S.M. Jiang, J. Gong, C. Sun, Composite Coatings with and without an in Situ Forming Cr-Based Interlayer: Preparation and Oxidation Behaviour, Corros. Sci. 53 (2011) 2894-2901. [3] Y.Z. Liu, X.B. Hu, S.J. Zheng, Y.L. Zhu, H. Wei, X.L. Ma, Microstructural Evolution of the Interface between NiCrAlY Coating and Superalloy During Isothermal Oxidation, Mater. Des. 80 (2015) 63-69. [4] J.L. Wang, M.H. Chen, L.L. Yang, S.L. Zhu, F.H. Wang, Comparative Study of Oxidation and Interdiffusion Behavior of AIP NiCrAlY and Sputtered Nanocrystalline Coatings on a Nickel-Based Single-Crystal Superalloy, Corros. Sci. 98 (2015) 530-40. [5] W.S. Walston, J.C. Schaeffer, W.H. Murphy, A New Type of Microstructural Instability in Superalloys - SRZ, Superalloys1996 (1996) 9-18. [6] X.Y. Gong, Y.H. Yang, Y. Ma, H. Peng, H.B. Guo, Microstructures and 15
Mechanical Properties of β-NiAlHf Coated Single Crystal Superalloy, Mater. Sci. Eng. A 673 (2016) 39-46. [7] Y. Matsuoka, Y. Aoki, K. Matsumoto, A. Satou, T. Suzuki, K. Chikugo, K. Murakami, The Formation of SRZ on a Fourth Generation Single Crystal Superalloy Applied with Aluminide Coating, Superalloys 2004 (2004) 637-642. [8] L. Shi, L. Xin, X.Y. Wang, X.L. Wang, H. Wei, S.L. Zhu, F.H. Wang, Influences of MCrAlY Coatings on Oxidation Resistance of Single Crystal Superalloy DD98M and Their Inter-Diffusion Behaviors, J. Alloy. Comp. 649 (2015) 515-530. [9] A. Suzuki, C.M.F. Rae, Secondary Reaction Zones in Coated 4th Generation Ni-Based Blade Alloys, Superalloys 2008 (2008) 777-786. [10] A.S. Suzuki, C.M.F. Rae, R.A. Hobbs, H. Murakami, Secondary Reaction Zone Formations in Pt-Aluminised Fourth Generation Ni-Base Single Crystal Superalloys, Adv. Mater. Res. 278 (2011) 78-83. [11] A.S. Suzuki, K. Kawagishi, T. Yokokawa, H. Harada, Effect of Cr on Microstructural Evolution of Aluminized Fourth Generation Ni-Base Single Crystal Superalloys, Surf. Coat. Technol. 206 (2012) 2769-2773. [12] X.Y. Gong, H. Peng, Y. Ma, H.B. Guo, S.K. Gong, Microstructure Evolution of an EB-PVD NiAl Coating and Its Underlying Single Crystal Superalloy Substrate, J. Alloy. Comp. 672 (2016) 36-44. [13] M.M. Kirka, K.A. Brindley, R.W. Neu, S.D. Antolovich, S.R. Shinde, P.W. Gravett,
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18
Table 1. Nominal compositions of the experimental alloys (wt. %). Alloy
Cr
Co
W
Mo
Al
Ti
Ta
Re
C
Ni
2Ta 6Ta 8Ta
9.5 9.5 9.5
8.5 8.5 8.5
4.5 4.5 4.5
0.5 0.5 0.5
4.5 4.5 4.5
2.3 2.3 2.3
2 6 8
2.3 2.3 2.3
0.05 0.05 0.05
bal bal bal
Table 2. Compositions of carbides and TCP phases determined by TEM with EDS (wt. %). Phase
Al
Ti
Cr
Co
Ni
Ta
W
Re
Mo
M23C6 σ µ
0.3 0.6 0.6
0.1 0.1 0.1
70.8 16.4 10.8
0.6 3.8 6.7
7.1 7.7 9
0.1 0.3 2.1
4 22.6 35.9
16.5 43.7 30.3
0.5 4.8 4.4
Fig.1. Cross-sectional microstructure of as-coated (a) 2Ta, (b) is the magnification of IDZ in (a) , (c) 6Ta and (d) 8Ta alloys. Fig. 2. Cross-sectional microstructure of 2Ta alloy after exposure at 1000 oC for 100 h, (a) back scattered electron (BSE) image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 3. Cross-sectional microstructure of 6Ta alloy after exposure at 1000 oC for 100 h, (a) BSE image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 4. Cross-sectional microstructure of 8Ta alloy after exposure at 1000 oC for 100 h, (a) BSE image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 5. TEM micrograph and respective electron diffraction patterns of M23C6 (a, d), σ ( b, e) and µ (c, f) in the IDZ and SRZ after exposure at 1000 oC for 100h. Fig. 6.TEM micrograph of (a) parallel rafts and (b) perpendicular rafts in 6Ta alloy after exposure 19
at 1000 oC for 100h. Fig.7. Elemental concentration profiles with depth from coating to substrate in 6Ta alloy after exposure at 1000 oC for 100 h. Fig. 8. EPMA element maps of 2Ta alloy after exposure at 1000 oC for 100 h. Fig. 9. EPMA element maps of 6Ta alloy after exposure at 1000 oC for 100 h. Fig. 10. EPMA element maps of 8Ta alloy after exposure at 1000 oC for 100 h. Fig. 11. Cross-sectional microstructure of alloys after exposure at 1000 oC for 1000 h, (a) 2Ta, (c) 6Ta, and (e) 8Ta, (b), (d) and (f) are magnified morphology of marked zone in (a), (c) and (e) respectively. Fig. 12. Cross-sectional microstructure of 2Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 13. Cross-sectional microstructure of 6Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 14. Cross-sectional microstructure of 8Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 15. EPMA element maps of 2Ta alloy after exposure at 1100 oC for 100 h. Fig. 16. EPMA element maps of 6Ta alloy after exposure at 1100 oC for 100 h. Fig. 17. EPMA element maps of 8Ta alloy after exposure at 1100 oC for 100 h. Fig. 18. Cross-sectional microstructure of alloys after exposure at 1100 oC for 1000 h, (a) 2Ta, (b) 20
6Ta, and (c) 8Ta alloys. Fig. 19. Microstructure of 8Ta alloys after exposure at 1100 oC for (a) 100 h, (b) is magnified morphology in zone marked A in (a), (c) 500 h and (d) 1000 h. Fig. 20. Schematics of perpendicular γ' rafting under the influence of the interdiffusion in the 6Ta alloy after thermal exposure.
21
Highlights
The influence of Ta on the microstructure degradation of NiCrAlYSi coated Ni-base single crystal superalloys at 1000 oC and 1100
o
C has been
investigated.
At 1000 oC, Ta encourages more inter-diffusion zone (IDZ) and secondary reaction zone (SRZ) formation.
At 1100 oC, Ta suppresses the formation of IDZ, and no SRZ is formed.
Ta has a potent influence on the γ' rafting perpendicular to the coating/substrate interface in the substrate beneath IDZ.
Declaration of interest statement The authors declare no conflict of interest.
Credit author statement We declare that this manuscript entitled “Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys” is original and integrated, and has not been published before, and is not currently being considered for publication elsewhere. We confirm that it will not be published elsewhere in the same form, in English and in any other language, including electronically without the written consent of the copyright holder. We confirm that the manuscript has been read and approved by all authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us.
S0925-8388(20)30803-3
DOI:
https://doi.org/10.1016/j.jallcom.2020.154440
Reference:
JALCOM 154440
To appear in:
Journal of Alloys and Compounds
Received Date: 15 October 2019 Revised Date:
6 February 2020
Accepted Date: 19 February 2020
Please cite this article as: B. Yin, G. Xie, L. Lou, J. Zhang, Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154440. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys Bin Yina,b, Guang Xiea,*, Langhong Loua, Jian Zhanga a. Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China b. School of Materials Science and Engineering, University of Science and Technology of China, Hefei 230026, China
Abstract The effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys at 1000 oC and 1100 oC has been investigated. Interdiffusion occurs between the NiCrAlYSi coating and the substrate, leading to the formation of inter-diffusion zone (IDZ) and secondary reaction zone (SRZ). After exposure at 1000 o
C for 100h, the initial γ/γ' transforms to γ' phases in the IDZ, and SRZ with σ and µ
topologically close-packed (TCP) phases is formed primarily due to the inward diffusion of Al. Different behavior of SRZ formation in the alloys with various Ta content can be observed after exposure at 1000 oC, which highlights the potency of Ta addition in promoting SRZ formation. This may be ascribed to the fact that Ta promotes the precipitation of TCP phases. After exposure at 1100 oC for 100h, the initial γ/γ' transforms to γ layer and coarse γ' phases in the IDZ, and no SRZ can be found. The formation of IDZ is significantly suppressed by Ta addition, which may be related to the retardation of interdiffusion between the coating and the substrate due to the increased γ' volume fraction. Furthermore, γ' rafting parallel to and/or 1
perpendicular to the coating/substrate interface in the substrate beneath IDZ are found in the alloys after exposure at 1000 oC and 1100 oC. Microstructural observation shows that Ta has a potent influence on the development of perpendicular γ' rafting. The mechanism of parallel γ' rafting and perpendicular γ' rafting has been discussed, respectively. Key words: Single crystal superalloys, NiCrAlYSi coating, Ta, Secondary reaction zone (SRZ), Rafting 1. Introduction Ni-based single crystal superalloys are widely used in industrial gas turbines, aircrafts and jet engines due to their excellent mechanical strength at elevated temperature. Usually, MCrAlY (M=Ni and/or Co) coatings both as stand-alone protective overlay coating and as a bond coat for thermal barrier coatings (TBCs) are widely applied on the superalloys to protect them from oxidation and corrosion attack [1-3]. However, interdiffusion between the coating and the underlying substrate inevitably occurs during thermal exposure at high temperature. Then Al in the coating is consumed, which accelerates the degradation of the coating. Moreover, microstructure degradation of the underneath substrate occurs as a result of local chemistry changes in the substrate [4]. On one hand, it has been extensively reported that inter-diffusion zone (IDZ) and even secondary reaction zone (SRZ) form beneath the coating/substrate interface, which causes considerable deterioration of mechanical properties in the alloy substrate [5-8]. SRZ located beneath the IDZ is comprised of a 2
three-phase transformation product involving γ, γ' and TCP phases [5]. The formation of SRZ is supposed to be dominantly controlled by the inward diffusion of aluminum from the coating layer to the substrate [7, 8]. Moreover, the morphology of SRZ strongly depends on Ni-base superalloy composition [9-11]. On the other hand, oriented γ' rafting in the region beneath the coating occurs during thermal exposure and the rafts are parallel to the coating/substrate interface, which strongly influences the deformation and failure of materials [12-16]. Furthermore, the composition of alloy has a potent influence on the parallel rafting [15]. However, to the knowledge of the authors, the γ' rafting perpendicular to the coating/substrate interface has not been reported. Ta is one of the major elements of Ni-base superalloys, which is benefitial to improve creep performance. However, few works have focused on the effect of Ta in the substrate on the microstructure degradation of the coated substrate. It has shown that Ta addition in the alloy significantly retards the propagation of the SRZ (consists of γ and β-NiAl with no TCP phases) during high-temperature exposure [17]. It is important for better alloy and coating design to reveal the effect of Ta on IDZ, SRZ (consists of γ, γ' and TCP phases) formation and γ' rafting of the coated superalloys. In this paper, the influence of Ta on the microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys associated with IDZ, SRZ formation and γ' rafting was investigated. In addition, the mechanisms of the above microstructural degradations were also discussed. 2. Experimental procedures 3
The three experimental single crystal (SX) superalloys, namely, 2Ta, 6Ta, and 8Ta, with nominal composition listed in Table 1 was directionally solidified into SX bars (16 mm in diameter and 220 mm in length) in a Bridgman furnace. As-cast SX alloy rods were solution heat treated followed by aging to develop the γ/γ' structure. Samples with the size of Φ15 × 2 mm were sliced from the heat-treated rods by using an electric discharge machine. The samples were ground with 800-mesh SiC paper, grit blasted in a wet atmosphere, and then ultrasonically cleaned in acetone. The NiCrAlYSi coating with nominal compositions Ni-18.1Cr-11.5Al-0.6Y-0.9Si (wt. %) were deposited on samples by arc ion plating. The thickness of the coating is about 40 µm. The coated samples were vacuum annealed in a sealed silica tube at 870 oC for 3.5 h. After that, isothermal oxidation of the coated samples was carried out in a muffle furnace at 1000 ℃ and 1100 ℃ for 1000 h respectively in static air. At intervals of 100, 300 and 500 h, some samples were taken out for microstructure analysis. The cross sectional microstructures of the samples were examined by scanning electron microscopy (SEM, Philips FEI-Inspet F) equipped with energy dispersive spectroscopy (EDS, Oxford X-man). An etchant (40 g CuSO4, 100ml HCl, 200 ml H2O) was used to reveal the microstructure. Compositional information and elemental mapping were obtained using a Cameca SX100 electron probe microanalyzer (EPMA, Shimadzu EPMA1610) operating at 20 kV. Besides, the JEOL JEM-2100 (200 kV) transmission electron microscope (TEM) was used to determine and analyze the phases. 3. Results 4
3.1. Microstructure of the as-coated alloys SEM micrographs of the cross-sectional as-coated samples (after annealed in a sealed silica tube at 870 oC for 3.5 h) are shown in Fig. 1. An interdiffusion zone (IDZ) formed beneath the coating/substrate interface in 2Ta, 6Ta and 8Ta alloys, which was mainly composed of γ' phases and Cr-rich precipitates. The IDZ thicknesses of 2Ta, 6Ta and 8Ta alloy are about 5.4±0.6, 3.7±0.2 and 3.1±0.3µm, respectively. The γ' precipitates were found to be regular and cuboidal beneath the IDZ in 2Ta, 6Ta and 8Ta alloys. 3.2. Interdiffusion and microstructural evolution at 1000 oC The cross-sectional morphologies of coated alloys after oxidation at 1000 oC for 100 h are displayed in Figs. 2-4. Further diffusion had occurred. The initial γ/γ' had transformed to γ' phases and rod-like and granular carbides were precipitated in the IDZ. The thicknesses of IDZ in 2Ta, 6Ta and 8Ta alloys are about 21±1.0, 9±0.8, 13 ±1.7 µm, respectively. Beneath the IDZ, secondary reaction zone (SRZ) with topologically close-packed (TCP) phases was observed. However, the morphology of SRZ varied greatly for three alloys. Almost no SRZ could be found in 2Ta alloy (Fig. 2), continuous SRZ formed in 6Ta and 8Ta alloys with the depth of 16±1.5 and 21± 3.5 µm, respectively (Figs. 3 and 4). It indicates that higher content of Ta in substrate encourages more SRZ formation. The rod-like and needle-like TCP phases were present in the SRZ of 6Ta and 8Ta alloys. The compositions of carbides and TCP phases are shown in Table 2. The Cr(Re)-rich carbide precipitated in the IDZ was determined as M23C6 by TEM (Figs. 5
5a, 5d). The rod-like TCP phase enriched in Cr, W and Re was determined as σ phase (Figs. 5b, 5e). While the needle-like TCP phase enriched in W and Re was determined as µ phase (Figs. 5c, 5f). Apart from the above microstructure changes observed in IDZ and SRZ, the bulk γ'/γ substrate degraded due to the interdiffusion during thermal exposure. Degradation of substrate associated with γ' rafting parallel to the coating/substrate interface was observed in 2Ta, 6Ta and 8Ta alloys (Figs. 2c, 3c and 4c), which was consistent with previous studies [14-16]. However, γ' rafting perpendicular to the coating/substrate interface beneath parallel rafts was also observed in 6Ta alloy (Fig. 3d). Fig. 6 shows the TEM micrograph of the parallel rafts and perpendicular rafts in 6Ta alloy. Plenty of dislocations were observed in the γ matrix of the parallel rafts region (Fig. 6a). While dislocations were not observed in the γ matrix of the perpendicular rafts region (Fig. 6b). The SEM with EDS was used to measure the chemistry profiles, and the experiment was carried out from the coating to the substrate in a step of 10 µm. The typical elemental concentration profiles with the depth of 6Ta alloy after thermal exposure at 1000 oC for 100 h are shown in Fig. 7. Chemical gradient of Al, Re and W was present beneath the IDZ due to the interdiffusion between the coating and the substrate. EPMA element maps of 2Ta, 6Ta and 8Ta alloys after exposure at 1000 ◦C for 100 h are shown in Figs. 8-10, respectively. Due to the chemical gradient between NiCrAlYSi coating and substrate, Ni, Cr and Al diffused from the coating into the substrate, while W, Mo, Ta, Co, Re and Ti were along the opposite direction. As 6
illustrated in Figs. 8-10, the concentration of Al and Ni in the IDZ is much higher than that in the substrate. A thin layer of Ti-rich phase formed between the coating and the substrate. The thickness of Ti-rich layer decreased as the content of Ta in the substrate increased. Granular carbides distributed along the coating/substrate interface were observed. The cross-sectional morphologies of 2Ta, 6Ta and 8Ta alloys after oxidation at 1000 oC for 1000 h are displayed in Fig. 11. The IDZ and SRZ significantly grew. The thicknesses of IDZ+SRZ in 2Ta, 6Ta and 8Ta alloys are about 42±3.3, 66±4.6 and 73±6.5µm, respectively. A significant amount of TCP phases are precipitated in the dendrite of 8Ta alloy, which in turn promotes the formation of SRZ. The parallel γ' rafting almost disappeared, and perpendicular γ' rafting was observed in 6Ta and 8Ta alloys. 3.3. Interdiffusion and microstructural evolution at 1100 oC The cross-sectional morphologies of coated alloys after exposure at 1100 oC for 100 h are displayed in Figs. 12-14. The IDZ was composed of a thin layer of γ phase beneath the coating/substrate interface and some coarse γ' phases located beneath the γ layer. These coarse γ' phases were dissolved during etching and only voids have left (Figs. 12b, 13b, 14b). Moreover, it is noted that the thickness of γ layer decreased and the content of coarse γ' phases increased with increasing the content of Ta. No SRZ was observed in three alloys. The depth of IDZ in 2Ta, 6Ta and 8Ta alloys was about 19±1.7, 16±2.1 and 12±1.1 µm, respectively. Ta was shown to suppress the formation of IDZ. 7
Degradation of the substrate associated with γ' rafting parallel to the coating/substrate interface was observed beneath the IDZ (Figs. 12c, 13c, 14c). Moreover, the perpendicular γ' rafting were present in both 6Ta and 8Ta alloys (Figs. 13d, 14d). EPMA element maps after exposure at 1100 oC for 100 h are shown in Figs. 15-17. Ni and Cr diffused from the coating into the substrate, while W, Mo, Ta, Co, Re and Ti were along the opposite direction. However, the inward diffusion of Al hardly occurred, due to the fact that the concentration of Al in the coating was not higher than that in the substrate after oxidation consumption (Figs. 15-17). In the 6Ta and 8Ta alloys, some granular bright phases were present in the substrate beneath the IDZ. According to the EPMA results shown in Figs. 16 and 17, the granular phase was Ti, Ta-rich carbide. A thin layer of Ti-rich phase was also found between the coating and substrate. Moreover, the thickness of Ti-rich layer decreased with the increase of Ta content in the substrate. The cross-sectional morphologies of 2Ta, 6Ta and 8Ta alloys after oxidation at 1100 oC for 1000 h are displayed in Fig. 18. The IDZ significantly grew. The thicknesses of IDZ in 2Ta, 6Ta and 8Ta alloys are about 73±6.8, 30±1.3 and 23±2.1 µm, respectively. The parallel γ' rafting almost disappeared, and perpendicular γ' rafting was observed in 6Ta and 8Ta alloys. The 8Ta alloy showed poor microstructure stability and TCP phases precipitated in the substrate after thermal exposure. However, the TCP phases disappeared as the interdiffusion between the coating and the substrate further occurred (Fig. 19). The depth of TCP-free region 8
significantly increased with increasing the exposure time. 4. Discussions 4.1.
Interdiffusion
and
microstructural
evolution
behavior
near
the
coating/substrate interface The formation of IDZ and SRZ beneath coating has been ascribed to the interdiffusion of elements between the coating and the substrate [5-8]. After exposure at 1000 oC, the content of aluminum is significantly increased in the IDZ due to the inward diffusion of Al (Figs. 8-10). This leads to the transformation from γ/γ' to γ' phases in the substrate adjacent to the coating and the resulting precipitation of carbides to accommodate excess chromium and rhenium (Fig. 2). As a forming element of γ' phase, Ta significantly promotes the above transformation. However, the growth of IDZ has been suppressed by the formation of SRZ in 6Ta and 8Ta alloys. Therefore, the depth of IDZ in 2Ta alloy without SRZ is the largest in the three alloys, while the depth of IDZ in 8Ta alloy is larger than that in 6Ta alloy. It has been suggested that SRZ formation was driven by the outward diffusion of Ni from the substrate into the coating [18]. In the present study, however, the diffusion of Ni is along the opposite direction (Figs. 8-10). It is likely ascribed to that the concentration of Ni in the NiCrAlYSi coating (68.9 wt. %) is significantly higher than that in 2Ta (65.8 wt. %), 6Ta (61.8 wt. %) and 8Ta (59.8 wt. %) alloys. Therefore, it has been suggested that the inward of Al resulted in the SRZ formation after exposure at 1000 oC [7, 8]. Two types of TCP phases are precipitated in the SRZ of 6Ta and 8Ta alloys. The σ phase enriched in Cr and Re prefers to precipitate nearby 9
the IDZ probably due to the enrichment of Cr in this region. While, the W, Re-rich µ phase is distributed nearby the substrate probably because of high content of W. Different SRZ formation behavior in 2Ta, 6Ta and 8Ta alloys is observed after exposure at 1000 oC, which highlights the potency of Ta addition in promoting SRZ formation (Figs. 2-4). Generally, the formation of TCP phases in superalloys can be predicted by the average electron-hole concentration (Nv) in the substrate [19]. As a refractory element, Ta significantly contributes to the Nv and promotes the precipitation of TCP phases. This is in good agreement with the difference in the microstructure instability of substrate far away from the coating /substrate interface in 2Ta, 6Ta and 8Ta alloys (Figs. 2-4). Therefore, Ta promotes the precipitation of TCP phase and leads to more SRZ formation. After exposure at 1100 oC, the consumption of Al in the coating increased due to higher temperature oxidation, resulting in less inward diffusion of Al. In contrast, the inward diffusion of Ni accelerates with increasing the exposure temperature. Therefore, a layer of γ phase formed beneath the coating/substrate interface. The elemental diffusivities in the ordered γ'-Ni3Al phase are generally 1-2 orders of magnitude lower than that in γ-Ni [21]. Therefore, the formation of γ layer beneath coating/substrate interface promotes the diffusion of Re and W from the substrate to the coating, leading to the depletion of W and Re in the section beneath the γ layer, which promotes the growth of γ' phases. Therefore, the IDZ is composed of γ layer and coarse γ' phases. Ta is shown to significantly suppress the IDZ formation after exposure at 1100 10
o
C (Figs. 12-14). It has been reported that Ta decreases the inter-diffusivity between γ
and γ' phases and improves the stability of γ' phase [17, 20]. So Ta addition in the substrate suppresses transformation from γ/γ' to γ phase layer, and retards the growth of IDZ. Moreover, Ta significantly increases the volume fractions of γ' phases, which impedes the interdiffusion between the coating and the substrate [17]. Therefore, the growth of IDZ is further suppressed. Moreover, the thickness of Ti-rich layer decreases with increasing Ta content (Figs. 15-17). No SRZ is observed in 2Ta, 6Ta and 8Ta alloys after exposure at 1100 oC, which is ascribed to the inward diffusion of Ni and outward diffusion of Re and W. The TCP phases precipitated in the substrate of 8Ta alloy disappear due to the outward diffusions of Re and W further proceeding (Fig. 19). However, the inward diffusion of Cr accelerates, leading to the precipitation of Cr-rich carbide beneath the IDZ in 6Ta and 8Ta alloys after exposure at 1100 oC for 100 h (Figs. 13 and 14). 4.2. Degradation of substrate during interdiffusion In addition to the microstructure changes observed in the IDZ and SRZ, the bulk γ/γ' substrate degrades after exposure at 1000 oC and 1100 oC. Parallel γ' rafting and perpendicular γ' rafting in the substrates beneath IDZ were observed. The γ' rafting parallel to the coating/substrate interface is caused by the externally applied stresses as well as the positive or negative sense of the misfit [14-16]. Furthermore, the compositions of substrate has a potent influence on the parallel γ' rafting, due to the probably change of γ/γ' misfit parameter during the interdiffusion between the coating and substrate [15]. However, in the present study, 11
parallel γ' rafting occurs in the 2Ta, 6Ta and 8Ta alloys both at 1000 oC and 1100 oC. Ta addition has little effect on the parallel γ' rafting. The parallel γ' rafting may be primarily ascribed to the strain induced by grit blasting. The threshold strain for rafting to continue in the absence of applied stress is as small as 0.1%, as reported by Reed et al. [22]. Therefore, the strain induced by grit blasting may provide the driving force for parallel γ' rafting of 2Ta, 6Ta and 8Ta alloys. As shown in Fig. 6a, Plenty of dislocations are distributed in the γ matrix of parallel rafts region, which is similar with that in the deformation-induced rafting during creep [23]. However, the perpendicular γ' rafting beneath the IDZ has not been reported. According to the results shown in Figs. 2-4 and 12-14, Ta has a potent influence on the perpendicular rafting. Few dislocations are present near the perpendicular γ' rafts (Fig. 6b). The formation mechanism of perpendicular γ' rafting may be different from that of parallel γ' rafting. The γ' rafting during load-free annealing has been reported in materials like AM1 and CMSX-4 [24-26]. It was suggested that the local chemical heterogeneities would lead to corresponding chemical gradients on the scale of the dendrite structure, which may play a crucial role in orientating the γ' coarsening. In the present study, the mechanism of perpendicular rafting can be schematically summarized as shown in Fig. 20. With no chemical gradients, the γ′ may grow by equally moving any of the three equivalent {001} γ/γ′ interfaces, and the γ and γ′ forming elements diffuse and exchange (Fig. 20a). Due to the interdiffusion between the coating and substrate, the chemical gradients of Al, Re and W are formed in the substrate, which may lead to the 12
directional movement of the corresponding elements into the same directions as the gradients. More alloying elements for phase transformation would diffuse in the perpendicular direction (Fig. 20b). The γ′ phase is prone to coalesce vertically. The three experimental alloys display negative misfit at high temperatures. It is generally known that Al with low solubility in the γ phase is mainly segregated in the γ' phase. The lattice parameters of γ decreases due to the outward diffusion of Re and W, leading to the relax of misfit stresses in the vertical matrix channels. While, the misfit stresses in the horizontal matrix channels persist relatively (Fig. 20b). It has been reported that the difference in misfit stresses on the horizontal faces and vertical faces of γ' phase, causing by the formation of horizontal dislocation networks, results in the γ' rafting [23, 27]. Thus, this promotes further vertical growth of γ' phase. Therefore, the γ′ is prone to coalesce vertically rather than laterally, which results in the formation of perpendicular rafting in 6Ta alloy after exposure at both 1000 oC and 1100 oC (Fig. 20c). Furthermore, the magnitude of the misfit is also found to affect the rate at which rafts develops [23, 25]. Ta has been shown to increases the lattice parameters of both γ' and γ phases, but increase the lattice misfit [28]. The misfit stresses increase with increasing Ta addition, resulting in the increase of driving force for rafting. In the 2Ta alloy, the misfit stresses are not enough for perpendicular rafting at both 1000 oC and 1100 oC. Reed has reported that higher γ' volume fraction with lower γ channel width restricting dislocation movement, resulting in greater resistance to rafting [25]. There is no perpendicular rafts in the 8Ta alloy after exposure at 1000 oC for 100h, although 13
the misfit of 8Ta alloy is larger. When the exposure time prolongs to 1000h or the exposure temperature increases to 1100 oC, the interdiffusion accelerates and the chemical gradient of Al, Re and W beneath the IDZ significantly increases, which promotes the perpendicular rafting. Moreover, the misfit stresses increased with increasing temperature, resulting in the increase of driving force for rafting. Therefore, perpendicular rafting in 8Ta alloy occurs after exposure at 1000 oC for 1000h and at 1100 oC. 5. Conclusions In summary, the following conclusions can be drawn from this work: 1. After exposure at 1000 oC, the initial γ/γ' transformed to γ' and Cr, Re-rich M23C6 carbides in the IDZ of 2Ta, 6Ta and 8Ta alloys. Moreover, the SRZ with σ and µ topologically close-packed (TCP) phase formed in the 6Ta and 8Ta alloys. It indicates that Ta addition promotes the formation of SRZ, which may be ascribed to the fact that Ta promotes the precipitation of TCP. 2. After exposure at 1100 oC, the initial γ/γ' transformed to a γ layer and some coarse γ' phases in the IDZ. No SRZ was present in the 2Ta, 6Ta and 8Ta alloys. Ta was shown to suppress the growth of IDZ, which may be related with the increase of γ' volume fraction, retarding the interdiffusion between the coating and the substrate. 3. The γ' rafts parallel to the coating/substrate interface in the substrates beneath IDZ were found in the 2Ta, 6Ta and 8Ta alloys both at 1000 oC and 1100 oC. While the perpendicular rafts were present in the 6Ta alloy after exposure at both 1000 oC and 1100 oC and 8Ta alloy after exposure at 1000 oC for 1000h and at 1100 oC. It is 14
believed that the parallel rafting can be attributed to the strain induced by grit blasting. While, the perpendicular rafting are probably ascribed to the high chemical gradient of Al, Re and W caused by the interdiffusion between the coating and the substrate. Acknowledgements This work was financially supported by the National Natural Science Foundation of China under grant No. 51771204, No. U1732131, No. 51631008, No. 91860201 and No. 51911530154.The authors are grateful for these supports. References [1] G.W. Goward, Progress in Coatings for Gas Turbine Airfoils, Surf. Coat. Technol. 108 (1998) 73-79. [2] J. Ma, S.M. Jiang, J. Gong, C. Sun, Composite Coatings with and without an in Situ Forming Cr-Based Interlayer: Preparation and Oxidation Behaviour, Corros. Sci. 53 (2011) 2894-2901. [3] Y.Z. Liu, X.B. Hu, S.J. Zheng, Y.L. Zhu, H. Wei, X.L. Ma, Microstructural Evolution of the Interface between NiCrAlY Coating and Superalloy During Isothermal Oxidation, Mater. Des. 80 (2015) 63-69. [4] J.L. Wang, M.H. Chen, L.L. Yang, S.L. Zhu, F.H. Wang, Comparative Study of Oxidation and Interdiffusion Behavior of AIP NiCrAlY and Sputtered Nanocrystalline Coatings on a Nickel-Based Single-Crystal Superalloy, Corros. Sci. 98 (2015) 530-40. [5] W.S. Walston, J.C. Schaeffer, W.H. Murphy, A New Type of Microstructural Instability in Superalloys - SRZ, Superalloys1996 (1996) 9-18. [6] X.Y. Gong, Y.H. Yang, Y. Ma, H. Peng, H.B. Guo, Microstructures and 15
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Influence
of
Coarsened
and
Rafted
Microstructures
on
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Thermomechanical Fatigue of a Ni-Base Superalloy, Int J Fatigue 81 (2015) 191-201. 16
[14] Q.M. Wang, H. Li, M.H. Guo, P.L. Ke, J. Gong, C. Sun, L.S. Wen, Thermal Shock Cycling Behavior of NiCoCrAlYSiB Coatings on Ni-Base Superalloys II. Microstructure Evolution, Mater. Sci. Eng. A 406 (2005) 350-357. [15] M.Z. Alam, D.V.V. Satyanarayana, D. Chatterjee, R. Sarkar, D.K. Das, Creep Behavior of Pt-Aluminide (PtAl) Coated Directionally Solidified Ni-Based Superalloy CM-247LC After Thermal Exposure, Mater. Sci. Eng. A 641 (2015) 84-95. [16] D. Texier, D. Monceau, Z. Hervier, E. Andrieu, Effect of Interdiffusion on Mechanical and Thermal Expansion Properties at High Temperature of a MCrAlY Coated Ni-Based Superalloy, Surf. Coat. Technol. 307 (2016) 81-90. [17] T. Galiullin, A. Chyrkin, R. Pillai, R. Vassen, W.J. Quadakkers, Effect of Alloying Elements in Ni-Base Substrate Material on Interdiffusion Processes in MCrAlY -Coated Systems, Surf. Coat. Technol. 350 (2018) 359-368. [18] D.K. Das, K.S. Murphy, S.W. Ma, T.M. Pollock, Formation of Secondary Reaction Zones in Diffusion Aluminide-Coated Ni-Base Single-Crystal Superalloys Containing Ruthenium, Metall. Mater. Trans. A 39A (2008) 1647-1657. [19] B. Seiser, R. Drautz, D.G. Pettifor, TCP Phase Predictions in Ni-Based Superalloys: Structure Maps Revisited, Acta Mater. 59 (2011) 749-763. [20] C.B. Morrison, R.D. Noebe, D.N. Seidman, Effects of Tantalum on the Temporal Evolution of a Model Ni-Al-Cr Superalloy During Phase Decomposition, Acta Mater. 57 (2009) 909-20. [21] C.E. Campbell, Assessment of the Diffusion Mobilites in the Gamma-Prime and B2 Phases in the Ni-Al-Cr System, Acta Mater. 56 (2008) 4277-4290. 17
[22] N. Matan, D.C. Cox, C.M.F. Rae, R.C. Reed, On the Kinetics of Rafting in CMSX-4 Superalloy Single Crystals, Acta Mater. 47 (1999) 2031-2045. [23] T.M. Pollock, A.S. Argon, Directional Coarsening in Nickel-Base Single Crystals with High Volume Fractions of Coherent Precipitates, Acta Mater. 42 (1994) 1859-1874. [24] A. Hazotte, J. Lacaze, Chemically oriented γ′ plate development in a nickel base superalloy, Scripta Mater. 23 (1989) 1877-1882. [25] R.C. Reed, D.C. Cox, C.M.F. Rae, Kinetics of Rafting in a Single Crystal Superalloy: Effects of Residual Microsegregation, Mater. Sci. Technol. 23 (2007) 893-902. [26] K.Y. Cheng, C.Y. Jo, D.H. Kim, T. Jin, Z.Q. Hu, Influence of local chemical segregation on the γ′ directional coarsening behavior in single crystal superalloy CMSX-4, Mater. Char. 60(2009) 210-218. [27] J. Y. Buffiere, M. Ignat, A Dislocation Based Criterion for the Raft Formation in Nickel-Based Superalloys Single Crystals, Acta Mater. 43 (1995) 1791-1797. [28] Y. Ru, H. Zhang, Y.L. Pei, S.S. Li, X.B. Zhao, S.K. Gong, H.B. Xu, Improved 1200°C Stress Rupture Property of Single Crystal Superalloys by γ'-Forming Elements Addition, Scripta Mater. 147 (2018) 21-26.
18
Table 1. Nominal compositions of the experimental alloys (wt. %). Alloy
Cr
Co
W
Mo
Al
Ti
Ta
Re
C
Ni
2Ta 6Ta 8Ta
9.5 9.5 9.5
8.5 8.5 8.5
4.5 4.5 4.5
0.5 0.5 0.5
4.5 4.5 4.5
2.3 2.3 2.3
2 6 8
2.3 2.3 2.3
0.05 0.05 0.05
bal bal bal
Table 2. Compositions of carbides and TCP phases determined by TEM with EDS (wt. %). Phase
Al
Ti
Cr
Co
Ni
Ta
W
Re
Mo
M23C6 σ µ
0.3 0.6 0.6
0.1 0.1 0.1
70.8 16.4 10.8
0.6 3.8 6.7
7.1 7.7 9
0.1 0.3 2.1
4 22.6 35.9
16.5 43.7 30.3
0.5 4.8 4.4
Fig.1. Cross-sectional microstructure of as-coated (a) 2Ta, (b) is the magnification of IDZ in (a) , (c) 6Ta and (d) 8Ta alloys. Fig. 2. Cross-sectional microstructure of 2Ta alloy after exposure at 1000 oC for 100 h, (a) back scattered electron (BSE) image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 3. Cross-sectional microstructure of 6Ta alloy after exposure at 1000 oC for 100 h, (a) BSE image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 4. Cross-sectional microstructure of 8Ta alloy after exposure at 1000 oC for 100 h, (a) BSE image, (b) magnified morphology of IDZ and substrate after etching, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 5. TEM micrograph and respective electron diffraction patterns of M23C6 (a, d), σ ( b, e) and µ (c, f) in the IDZ and SRZ after exposure at 1000 oC for 100h. Fig. 6.TEM micrograph of (a) parallel rafts and (b) perpendicular rafts in 6Ta alloy after exposure 19
at 1000 oC for 100h. Fig.7. Elemental concentration profiles with depth from coating to substrate in 6Ta alloy after exposure at 1000 oC for 100 h. Fig. 8. EPMA element maps of 2Ta alloy after exposure at 1000 oC for 100 h. Fig. 9. EPMA element maps of 6Ta alloy after exposure at 1000 oC for 100 h. Fig. 10. EPMA element maps of 8Ta alloy after exposure at 1000 oC for 100 h. Fig. 11. Cross-sectional microstructure of alloys after exposure at 1000 oC for 1000 h, (a) 2Ta, (c) 6Ta, and (e) 8Ta, (b), (d) and (f) are magnified morphology of marked zone in (a), (c) and (e) respectively. Fig. 12. Cross-sectional microstructure of 2Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 13. Cross-sectional microstructure of 6Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 14. Cross-sectional microstructure of 8Ta alloy after exposure at 1100 oC for 100 h, (a) BSE image, (b) etched image, (c) and (d) are magnified morphology in zone marked A and B in (b) respectively. Fig. 15. EPMA element maps of 2Ta alloy after exposure at 1100 oC for 100 h. Fig. 16. EPMA element maps of 6Ta alloy after exposure at 1100 oC for 100 h. Fig. 17. EPMA element maps of 8Ta alloy after exposure at 1100 oC for 100 h. Fig. 18. Cross-sectional microstructure of alloys after exposure at 1100 oC for 1000 h, (a) 2Ta, (b) 20
6Ta, and (c) 8Ta alloys. Fig. 19. Microstructure of 8Ta alloys after exposure at 1100 oC for (a) 100 h, (b) is magnified morphology in zone marked A in (a), (c) 500 h and (d) 1000 h. Fig. 20. Schematics of perpendicular γ' rafting under the influence of the interdiffusion in the 6Ta alloy after thermal exposure.
21
Highlights
The influence of Ta on the microstructure degradation of NiCrAlYSi coated Ni-base single crystal superalloys at 1000 oC and 1100
o
C has been
investigated.
At 1000 oC, Ta encourages more inter-diffusion zone (IDZ) and secondary reaction zone (SRZ) formation.
At 1100 oC, Ta suppresses the formation of IDZ, and no SRZ is formed.
Ta has a potent influence on the γ' rafting perpendicular to the coating/substrate interface in the substrate beneath IDZ.
Declaration of interest statement The authors declare no conflict of interest.
Credit author statement We declare that this manuscript entitled “Effect of Ta on microstructural evolution of NiCrAlYSi coated Ni-base single crystal superalloys” is original and integrated, and has not been published before, and is not currently being considered for publication elsewhere. We confirm that it will not be published elsewhere in the same form, in English and in any other language, including electronically without the written consent of the copyright holder. We confirm that the manuscript has been read and approved by all authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us.