A duplex nanocrystalline coating for high-temperature applications on single-crystal superalloy

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A duplex nanocrystalline coating for high-temperature applications on single-crystal superalloy Lanlan Yang, Minghui Chen ∗ , Jinlong Wang, Shenglong Zhu, Fuhui Wang Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China

a r t i c l e

i n f o

Article history: Received 13 July 2015 Received in revised form 24 September 2015 Accepted 25 September 2015 Available online xxx

a b s t r a c t A new duplex nanocrystalline coating is designed for high temperature oxidation and hot corrosion protection. This coating combines the advantages of traditional NiCrAlY and nanocrystalline coatings, i.e., providing high resistance to oxidation and hot corrosion simultaneously, while avoids any disadvantages that the traditional coatings have suffered from, such as scale spallation, element interdiffusion (along with the formation of harmful TCP phases). It gives a good choice as the bond coating of a TBC system. © 2015 Elsevier Ltd. All rights reserved.

Keywords: A. superalloys A. metal coating C. hot corrosion C. oxidation

1. Introduction The quest for improving engine efficiency and reducing pollutants has led to increasingly higher operating temperatures and extending the operating lives of gas-turbine superalloys. Considering operation surroundings of gas turbine, only composite materials are able to meet the harsh requirements: the necessary mechanical properties is offered by the base material such as singlecrystal superalloy for it can provide higher temperature capacity as compared to its polycrystalline counterpart; and the protection against oxidation or corrosion is furnished by coating in the case of thermal barrier coatings (TBCs) [1–3]. The TBCs coupled with modern cooling systems can provide up to an almost 150 ◦ C temperature reduction at the superalloy blade surface. A TBC usually consists of a thermal-insulating ceramic top coating and an oxidation-resistant metallic bond coating [4–7]. As we all know, all the metallic coatings including the traditional NiCrAlY coating and nanocrystalline coating rely on the formation of slow-growing, stable, and adherent surface oxides such as Al2 O3 or Cr2 O3 to protect them from oxidation and hot corrosion. However, as the service environment has become increasingly harsh, the current coatings feel more and more difficult to meet all the applying requirements. NiCrAlY coating [8,9] that has been applied universally as stand along protective coating or as bond coat of TBC even suffers from many threatens, such as cracking or scale

∗ Corresponding author. Fax: +86 24 23893624. E-mail address: [email protected] (M. Chen).

spallation under thermal cyclic environments and continuing fast consumption of Al or Cr. Cracking or spallation of the oxide scale TGO (thermal grown oxide) plays the most important role in the final invalidation of TBC [10,11]. While the Al diffusion from coating into the underlying alloy substrate leads to the formation of SRZ (second reaction zone) with abundant harmful TCP (topologically close-packed) phases that destroy the mechanical properties of the superalloy component [12,13]. Accordingly, many novel metallic coatings [14–17] or interdiffusion barriers [18–23] were invented. However, none of these invents can completely solve all the problems simultaneously. Nanocrystalline coating invented by Lou and Wang twenty years ago [24], with the same composition to the alloy substrate, is one of such designs to avoid the problems of element interdiffusion and scale spallation. It still has a shortcoming of low corrosion resistance in sulfate environment. Corrosive media such as S and Cl- will diffuse easily along boundaries of the nano-grains as well to the superalloy substrate front, leading to hot corrosion [25,26]. In this work, we designed and prepared a duplex nanocrystalline coating which was expected to integrate the advantages of both the conventional NiCrAlY coating and the late-model sputtering nanocrystalline coating while avoids their shortcomings. 2. Experimental Procedure Cylindrical specimens of  15 × 2.0 mm were machined from a second generation single-crystal superalloy N5 bar. They were ground consecutively with 240#, 400#, 1000# and 2000# SiC papers, and then degreased by an ultrasonic cleaner in acetone

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Fig. 1. (a) Fracture section and (b) AFM surface morphologies of the as-sputtered duplex nanocrystalline coating.

before magnetron sputtering. The target used for sputtering the inner nanocrystalline layer was a 382 × 128×8 mm sheet with the same composition of the superalloy substrate. Sputtering parameters were deposited as follows: argon pressure was 0.2 Pa; sputtering current was 3.5 A; and substrate temperature was 200 ◦ C. The samples were rotated in the target during sputtering to ensure better uniformity. Thereafter, NiCrAlY outer layer was

deposited upon the inner nanocrystalline layer using arc ion plating (AIP) in argon pressure of 0.2 Pa to obtain a two-layer structured duplex nanocrystalline coating. Compositions of the two layers are listed in Table 1. Isothermal and cyclic oxidation tests were adopted to assess the interdiffusion, oxidation and scale spallation behaviors of the duplex nanocrystalline coating. Isothermal oxidation test was con-

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Fig. 2. (a) SEM cross-sectional microstructure of the duplex nanocrystalline coating, (b) EDS scanning along the yellow line in (a).

ducted in static atmosphere at 1050 ◦ C for 1000 h. Specimens with or without coatings were placed in crucibles which were all heattreated in furnace at 1200 ◦ C for enough time to ensure that their

weight would not change in the following isothermal oxidation test. After a certain period of time at elevated temperature, samples were removed out from the muffle furnace and cooled down

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Table 1 Nominal composition of the single-crystal superalloy N5 and the NiCrAlY coating (wt%) Elements

Ni

Co

Cr

Mo

W

Ta

Al

Re

N5 Nominal composition NiCrAlY composition

Bal. Bal.

7.5

7.0 27

1.5

5.0

6.5

6.2 11

3.0

Y 0.5

for 3 h in dry oven to room temperature. Then the samples along with the crucibles were weighted. Thermal cyclic oxidation tests were taken at 1050 ◦ C for 400 cycles. In each cycle the specimens were exposed to the elevated temperature for 1 h and then taken out and cooled in air for 15 min. The samples were weighted after cooling to room temperature for certain numbers of cycle. Hot corrosion was conducted at 850 ◦ C for specimens with 2.5–3.0 mg/cm2 sulfate (75 wt% Na2 SO4 + 25 wt% K2 SO4 ) sprayed on their surfaces. Samples with the sulfate were put into muffle furnace for a certain time, and then taken out and cooled down to room temperature. The samples were weighted after cleaning the surface sulfate with distilled boiling water for three times and drying. For comparison, specimens of the single-crystal superalloy substrates, sputtering nanocrystalline coatings, NiCrAlY coatings were tested as well. By means of atomic force microscope (AFM, N9401S, NAVITAR, UK), scanning electron microscope (SEM, Inspect F 50, FEI CO., Hillsboro, Oregon) with energy dispersive spectrometer (EDS, X-Max, Oxford instruments Co, UK) and transmission electron microscope (TEM, JEM-2100F, JEOL, Tpkyo, Japan), surface and microstructure behavior were studied. And during oxidation and hot corrosion, the weight gain of the specimens was quantified discontinuously by electron balance with a sensitivity of 0.01 mg. 3. Results 3.1. As-sputtered coating Fig. 1 shows fracture section and the AFM surface morphologies of the as-sputtered duplex nanocrystalline coating. It exhibits a double-layered microstructure as we designed, of which the inner layer displays a columnar microstructure with a thickness of 19 um and the outer one with a thickness of 9 um just as shown in Fig. 1(a). Interfaces of inner layer/superalloy and inner layer/outer layer are well bonded physically. The outer layer of the coating presents a rough surface with many bulges randomly dispersed. Fig. 1(b) presents more details of the outer layer. It is no longer of a splaton-splat microstructure like a traditional NiCrAlY coating. Instead, it shows a cluster structure, of which each cluster is composed of abundant nano-sized granules. The more detailed information of the duplex nanocrystalline coating can be found in Fig. 2. Fig. 2(a) shows close combination at two interfaces, i.e., superalloy substrate/inner sputtering layer/outer NiCrAlY layer. The duplex nanocrystalline coating is about 28 ␮m in thick, of which the inner layer is about two times of the outer one. EDS analysis (Fig. 2(b)) reveals that there is no notable fluctuation in elements content across interface I (the interface of superalloy substrate/inner sputtering layer, which is denoted by vertical yellow lines in Fig. 2(b)). While at interface II, it shows big differences in chemical composition between the two layers. To be exactly, the percentages of Al and Cr are much higher in the out NiCrAlY layer than in the inner nanocrystalline one. Fig. 3 exhibits TEM photographs of the inner nanocrystalline layer of the duplex coating near the two interfaces I and II. The selected area electron diffraction (SAED) pattern of the inner nanocrystalline layer near interface I presents rings as shown in Fig. 3(a), indicating the grain size there is less than the diameter of selected area of SAED, i.e., less than 100 nm. However, near interface II the bright field TEM image shows that the grain size of the

Fig. 3. TEM images of the inter nanocrystalline layer of the duplex nanocrystalline coating: (a) selected area electron diffraction (SAED) pattern in the vicinity of interface I, (b) bright field image near interface II.

columnar grains is about 120 nm as shown in Fig 3(b). It seems that the grain size of the inner nanocrystalline layer increases from the interface I to interface II. Besides, abundant defects, such as stacking faults and dislocations, are clearly observed in the columnar grains near interface II. 3.2. Isothermal oxidation at 1050 ◦ C Fig. 4 shows isothermal oxidation kinetics of the N5 substrate, the NiCrAlY and the duplex nanocrystalline coatings at 1050 ◦ C. As shown in Fig. 2(a), the N5 substrate shows high oxidation resistance for the initial 200 h, however, its long time oxidation resistance is low due to severe spallation. The oxidation rate of the duplex nanocrystalline coating is a little higher than that of the NiCrAlY one. In order to make clear the oxidation regulation of the N5 superalloy, the NiCrAlY coating and the duplex nanocrystalline coating, a function relationship between log (y) and log (t) is plotted in Fig. 4(b), where y is weight gain and t is oxidation time. Just as shown in Fig. 4(b), the slopes of the two curves, log (y) vs. log (t), of the superalloy substrate and the NiCrAlY coating are of almost the same value. When fitting these curves, the first point corresponding to 5 h oxidation is discarded to avoid the influence of initial oxidation. However, the slope was not equal to 0.5, indicating that the oxidation did not obey exactly parabolic. Indeed, the slope equaled closely to 0.25, i.e., the oxidation followed the empirical law: yn = kt

(1)

where k is rate constant, n is a empirical exponent = 4. For many nanocrystalline alloy or coatings, it has been reported that the

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Fig. 4. Isothermal oxidation kinetics of the N5 superalloy, the NiCrAlY coating and the duplex nanocrystalline coating at 1050 ◦ C in air: (a) y vs. t and (b) log (y) vs. log (t).

exponent might deviate from 2 for their nano-microstructures [24,25,27–32]. But for the single crystal superalloy and the NiCrAlY coating their abnormal oxidation mechanisms might be associated with the Ta and Y effect. In fact, it has been previously reported that Ta and the reactive element Y can affect the mode of alumina scale, i.e., from equiaxed alumina grains to columnar grains. Similar oxidation regulations of single crystal superalloy N5 and NiCrAlY coating are observed at 1000 ◦ C in ref [26]. For the duplex nanocrystalline coating, its oxidation kinetic does not follow the empirical law Eq. (1) like the NiCrAlY coating does, though its outer layer has the same composition with the NiCrAlY coating. The reasons lie in the different microstructures of the two coatings. As shown in Fig. 1(b), the outer NiCrAlY layer of the duplex coating is of a nanostructure, which should be affected by the underlying inner nanocrystalline layer. Abundant grain boundaries exist in the duplex coating and give more fast diffusion way for ions than the traditional NiCrAlY coating. So it is not surprising that its oxidation

behavior deviates from the empirical law. In addition, the nanosized grains of the duplex coating are not stable during oxidation, i.e., they must grow. This unstable microstructure of the duplex coating leads to the fact that its oxidation law varies with oxidation time. Anyway, it is difficult to describe exactly the underlying relationships between the microstructure and the oxidation law of the duplex coating here and should be investigated in detail in our further study. Fig. 5 shows the microstructure of the NiCrAlY coating after oxidation at 1050 ◦ C for 1000 h. It presents the typical cauliflower-like morphologies. Spallation and cracking are universal at the coating surface. As shown in Fig. 5(a), the sizes of a certain spallation spot reaches to 54 × 67 ␮m, and the length of the crack is 26 ␮m. From the cross-sectional view (Fig. 5(c)), the oxide scale is not continuous, which is coincident with its surface morphology. After oxidation, the NiCrAlY coating has transformed almost fully from   +  phases to a single  phase due to the Al depletion [33–35].

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Fig. 5. Microstructures of the NiCrAlY coating after oxidation at 1050 ◦ C for 1000 h in air: (a) and (b) surface, (c) cross-section.

Its phase transformation from the original   +  phases to  phase leads to volume change of the NiCrAlY coating, which should contribute partly to the scale spallation. Beneath the NiCrAlY coating, the interdiffusion zone (IDZ) and second reaction zone (SRZ), containing large amounts of TCP precipitates with high content of Cr, Re, W and low Al, have developed to an average size of 24 ␮m into the N5 single crystal superalloy.Fig. 6 shows the surface and cross-sectional microstructures of the duplex nanocrystalline coat-

Fig. 6. Microstructures of the duplex nanocrystalline coating after oxidation at 1050 ◦ C for 1000 h in air: (a) surface, (b) cross-section and (c) EDS scanning along the yellow line in (b).

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Fig. 7. Cyclic oxidation kinetics of the N5 superalloy, the NiCrAlY coating and the duplex nanocrystalline coating at 1050 ◦ C in air.

ing after oxidation at 1050 ◦ C for 1000 h. It exhibits as well a typical cauliflower-like morphology at surface just like most AIP NiCrAlY coatings. However, neither spallation nor cracking takes place at this case. From the cross-sectional view (Fig. 6(b)), the oxide scale is continuous and flat without any obvious undulation. It combines well with the duplex nanocrystalline coating and oxide pegs are observed at some places. The average thickness of the oxide scale is about 6.8 ␮m, a little thicker than it formed on the single NiCrAlY coating. This is agreed with oxidation kinetic curves in Fig. 4(a), showing a small higher weight gain of the duplex nanocrystalline coating than the NiCrAlY one. It seems that the two layers of the duplex coating have been completely integrated. Correspondingly, alloy elements in the two layers have been completely averaged as analyzed by EDS (Fig. 6(c)). In vicinity to the original interface I, no TCP phases are found in the N5 single crystal superalloy, and EDS results indicate that no notable fluctuation in elements content occurs as the scanning line cross the interface. After a careful observation of the oxygen distribution curve in Fig. 6(c), it is found that O does not enter the interior of the duplex nanocrystalline coating. It only exists in the Al2 O3 scale. 3.3. Cyclic oxidation at 1050 ◦ C Cyclic oxidation kinetic curves of the N5 alloy substrate and its coatings at 1050 ◦ C are shown in Fig. 7. It is evident that the oxide scale formed on the superalloy substrate is very susceptible to cyclic oxidation. Oxide film spalls off severely only after 5 cycles. Scale spallation leads to a dramatic drop in weight of the N5 superalloy. The NiCrAlY coating shows a good spallation resistance for the initial 220 cycles. But weight loss happens after 240 cycles at 1050 ◦ C and the trend continues with oxidation cycles as indicated by arrow. However, no weight loss is observed for the duplex nanocrystalline coating during the whole 400 cycles at 1050 ◦ C. The oxidation rate of the duplex nanocrystalline coating was a little higher than the NiCrAlY coating for the 240 cycles which is consistent with the isothermal oxidation kinetics as shown in Fig. 4(a). Fig. 8 shows the typical surface morphologies of the NiCrAlY coating after 300 h cyclic oxidation at 1050 ◦ C. It exhibits the similar surface morphologies with the NiCrAlY coating after isothermal oxidation at 1050 ◦ C for 1000 h (Fig. 5) except that more spallation

spots and cracks are observed. The size of the two spallation spots in the visual field are about 64 × 45 ␮m and 36 × 29 ␮m. The longest length of cracks has propagated to 90 ␮m. By contrary, the surface of the duplex nanocrystalline coating is intact with no cracking or spallation formed after 400 h cyclic oxidation at 1050 ◦ C, as shown in Fig. 9(a). Anyway, it presents as well the cauliflower-like surface features. 3.4. Hot corrosion at 850 ◦ C Fig. 10 shows corrosion kinetics of the alloy and its coatings in molten 75% Na2 SO4 + 25% K2 SO4 at 850 ◦ C. As shown in the corrosion kinetics, a significant mass gain as a consequence of hot corrosion up to 40 h, followed by a linear rapid increase is observed for the alloy substrate. It means that unstable hot corrosion takes place. Though the sputtered nanocrystalline coating provides a good protection in molten 75% Na2 SO4 + 25% K2 SO4 at 850 ◦ C for the initial 80 h, it experiences a sudden decrease in mass gain thereafter as indicated by arrow. This is a common phenomenon since it is observed on all the six repeated samples. For NiCrAlY coating, it loosed weight even for the first 5 h hot corrosion, and followed this trend among the whole 100 h corrosion time. As compared to the single crystal superalloy substrate, the weight change of the samples coated with the NiCrAlY coating or nanocrystalline coating is much lower after hot corrosion for 100 h. Obviously these two coatings provide moderate corrosion protection for the alloy substrate from corrosion. However, their protection resistance is still low. The duplex nanocrystalline coating shows an even lower mass change during hot corrosion and no mass loss is observed, indicating the oxide scale is still protective. After 100 h hot corrosion in sulfate, there is only a little increase in weight, 0.115 mg/cm2 . Fig. 11 shows the surface morphologies of the nanocrystalline coating after hot corrosion in sulfate at 850 ◦ C for 100 h. Some spallation spots with so many pores at the central are observed from the surface view. It is clear that these pores can give a rapid path for the inward diffusion of corrosive media, such as S and O. EDS analysis indicates that Mo and W are enriched there. For the NiCrAlY coating, spallation spots of large scale have been seen at surface, as shown in Fig. 12. The size of spallation in the visual field is about 175×81 ␮m. EDS analysis (Fig. 12(c)) indicate that at the spalling

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Fig. 8. Surface morphology of the NiCrAlY coating after cyclic oxidation at 1050 ◦ C for 300 cycles in air (a) low magnification and (b) high magnification.

Fig. 9. Surface morphology of the duplex nanocrystalline coating after cyclic oxidatin at 1050 ◦ C in air for 400 cycles (a) low magnification and (b) high magnification.

Fig. 10. Hot corrosion kinetics of the N5 superalloy, the duplex nanocrystalline coating, the NiCrAlY coating and the nanocrystalline coating at 850 ◦ C for 100 h in sulfate salt.

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Fig. 11. Surface morphologies of the nanocrystalline coating after hot corrosion at 850 ◦ C for 100 h in sulfate salt: (a) low magnification and (b) high magnification.

spots there were so many Cr-rich oxides in addition to alumina. While for the duplex nanocrystalline coating, a perfect intact surface being covered full of oxides is observed. XRD analysis indicates that the oxide is ˛- and -Al2 O3 . As shown in Fig. 13, no spallation or crack is detected. This is agreed with its corrosion kinetic curve. The needle-like oxides in Fig. 13(b) are -Al2 O3 .

Fig. 12. Surface morphologies of the NiCrAlY coating after hot corrosion at 850 ◦ C for 100 h in sulfate salt: (a) low magnification, (b) high magnification, (c) EDS point analysis at point “a” in (b).

4. Discussion Just as we designed and prepared, the duplex coating contains a thin NiCrAlY outer layer and a thick nanocrystalline inner layer. The special coating has already reached to our expectation. Isothermal and cyclic oxidation at 1050 ◦ C reveals high resistance to oxidation and scale spallation, and hot corrosion at 850 ◦ C shows high corrosion resistance of the duplex nanocrystalline coating. Next we will discuss its protective mechanisms in detail. 4.1. Element interdiffusion Generally, NiCrAlY coating contains high content of Al and Cr for the convenience of fast growing of protective Al2 O3 scale at surface. However, as the service temperature has reached 1000 ◦ C

or higher, element diffusion between the Ni-based superalloy substrate and the NiCrAlY coating becomes especially serious, e.g., Cr and Al in NiCrAlY coating diffuse into the substrate superalloy. As observed in Fig. 5(c), the Al diffusion into the superalloy substrate promotes phase transformation from  to   and precipitation of TCP phases. With increasing oxidation time, most TCP phases forming in substrate will not disappear but grow. In this respect, it is necessary to control difference in chemical compositions between the metallic coating and the alloy substrate. As long as the difference is small enough, element interdiffusion could be reduced or even eliminated. For the novel designed duplex nanocrystalline coating, results show that no TCP phases precipitate out after isothermal oxidation at 1050 ◦ C for 1000 h. From this perspective, it could be basically

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superalloy substrate. This is the right reason that the outer NiCrAlY layer is designed very thin. 4.2. Scale spallation In addition to avoiding element interdiffusion, the duplex nanocrystalline coating exhibits higher spallation resistance than the NiCrAlY coating under cyclic circumstances or in sulfate environments, as shown in Figs. 7 and 10. During service, internal stresses generate in oxide scale owing to the oxide growing, CTE (Coefficient of Thermal Expansion) mismatch between oxide scale and coating, or phase transformation of the metallic coating. There are three possible ways for releasing such stress: (1) oxide scale plastic deformation; (2) alloy/coating plastic deformation; (3) oxide scale cracking and spallation. At high temperatures, plastic deformation of oxide scale is achieved mainly by diffusion creep. It has been reported that the .

creep rate, G, is closely associated with the grain size, d [36]: .

=

 B2 ıDb ) (B1 DV + d d2 KT

(2)

where B1 and B2 are constants,  is internal stress,  is atomic volume, DV and Db are lattice and grain boundary diffusion coefficient, ı is thickness of the grain boundary, K is the Boltzmann constant, and T is the Kelvin temperature. When the grain boundary diffusion is dominant, Eq. (2) can be simplified to .

=

Fig. 13. Surface morphologies of the duplex nanocrystalline coating after hot corrosion at 850 ◦ C for 100 h in sulfate salt: (a) the low magnification, (b) the high magnification.

expected that the high temperature mechanical properties of the superalloy substrate are not ruined during oxidation. However, this conclusion seems to be contradictory with the fact that the chemical composition of the outer NiCrAlY layer of the duplex nanocrystalline coating differs largely to the underlying superalloy substrate. At such a case, element interdiffusion between the duplex nanocrystalline coating and the single crystal superalloy substrate should occur. In fact, the key to avoid element interdiffusion lies in the two-layer structure of the duplex nanocrystalline coating. The thick inner nanocrystalline layer possesses the same chemical composition with the N5 single crystal superalloy, as shown in Figs. 1 and 2. In this way, element interdiffusion between the inner layer and the substrate should be avoided in the absence of disturbing of the outer NiCrAlY layer. Though the outer NiCrAlY layer has much higher content of Al and Cr than in the alloy substrate, its thickness is so thin that, most of the Al content will be consumed in forming an oxide scale. The others of composition differences are diluted by the inner thick nanocrystalline layer after a long period of oxidation. As shown in Fig. 6, the two-layer structured duplex nanocrystalline coating has integrated to be one, whose chemical composition deviates little to the N5 single crystal superalloy substrate. As a result, precipitation of TCP phases caused by element interdiffusion has been successfully prohibited in the

B˝ıDb d3 KT

(3)

The creep rate is inversely proportional to the third power of grain size. This means that fine grains of oxide scale facilitate creep deformation. There are so many defects and grain boundaries in the duplex nanocrystalline coating, which provide cores for oxides nucleation. The grain size of the oxide is thus very small. Comparing to the NiCrAlY coating, oxide scale formed on the duplex nanocrystalline coating can withstand a higher thermal stress in cyclic or sulfate environments. At the same time, the inner nanocrystalline layer of the duplex nanocrystalline coating is of a nano-sized columnar microstructure, which can release stress easily by their plastic deformation. After plastic deformation by both the underlying duplex nanocrystalline coating and the oxide scale, most of the internal stresses have been released and the residual stress dose not surpass the critical value for scale cracking or spalling off. This is the right reason why the duplex nanocrystalline coating shows higher resistance to scale cracking or spalling off than the traditional NiCrAlY coating does. 4.3. Internal sulfidation As we all know, refractory elements Mo, W and V in superalloys have a stronger affinity to O2− than Ni and Al and would form volatile MoO3 , WO3 and V2 O5 after a very short stable corrosion stage. The initial short stable corrosion belongs to alkaline corrosion. Some of these oxides react again with the oxygen ion in the molten sulfate following reactions: MoO3 + O2− = MoO2− 4

(4)

WO3 + O2− = WO2− 4

(5)

2−

V2 O5 + O

=

2VO− 3

(6) O2−

at the interface between the molten These reactions consume salt and the alloy/coating, making the molten sulfate to be acidic there. Thereafter, the initially formed thin oxide film of Al2 O3 or NiO on the surface decomposes to sustain the concentration of O2− : NiO = Ni2+ + O2−

(7)

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Al2 O3 = 2Al3+ + 3O2−

(8)

The ions Ni2+ , Al3+ , MoO4 2− , WO4 2− and VO3− produced by the above reactions then diffuse to the interface of molten salt/gas. On arriving at the out surface of molten salt, MoO4 2− , WO4 2− , VO3 − volatilize in the form of their oxides after decomposition following the inverse direction of reactions (4), (5) and (6). Accompanied with the decomposition of MoO4 2− , WO4 2− , VO3 − , O2− is released out at surface. With increasing the activity of O2− , reactions (7) and (8) take place toward the left hand, resulting in the precipitation of NiO and Al2 O3 at surface. Generally, the finally formed oxide scale composed of NiO and Al2 O3 is porous and less-protective. This hot corrosion process is named acidic corrosion and unstable [37–40].For the N5 single crystal superalloy and its sputtered nanocrystalline coating, acidic corrosion takes place by virtue of the presence of 5 wt% W and 1.5 wt% Mo in their compositions as shown in Table 1. The hot corrosion rate of the superalloy substrate is very high, especially after 40 h corrosion. Comparing with the superalloy substrate, a protective alumina scale is more easily to form on the nanocrystalline coating at the initial stable corrosion stage, which is revealed in the slow corrosion rate before 80 h as shown in Fig. 10. After hot corrosion at 850 ◦ C for 100 h, abundant pores form on scale surface of the nanocrystalline coating (Fig. 11), indicating the generation of volatile products. Combing with the EDS results showing enrichment of Mo and W around the pores, the acidic corrosion should have happened during corrosion of the nanocrystalline coating. In addition to the formed pores, .its columnar structure provides as well rapid diffusion channels for O or S, leading to internal sulfidation or oxidation. While for the NiCrAlY coating, its hot corrosion is alkaline corrosion since no Mo, W and V is included in the coating. Metallic elements, M (M =Ni, Cr and/or Al), react with the sulfate salt to form a thin oxide film following [41–45]: 2− 4M + SO2− 4 = MS + MO + O

(9)

Due to the above reaction, the local basicity, i.e., O2− activity, increases at the coating/molten salt interface. Then the following reaction occurs: MO + O2− = MO2− 2

(10) 2−

The generated MO2 is dissolved in sulfate salt and diffuse form the coating/molten salt interface to the out surface of molten salt. On arriving at the out surface, MO2 2− is decomposed following the inverse direction of reaction (10). In addition, the formed oxide scale on NiCrAlY coating has a high tendency to spalling off under the impact of deposited sulfate salt, as shown in Fig. 12. So corrosion is a competition between formation and consumption of oxide scale MO. The consumption of MO includes its dissolution in salt and spallation. This is the real corrosion mechanisms of NiCrAlY coating. To be exactly, the formation of Al2 O3 (the oxide scale is composed mostly of Al2 O3 ) on NiCrAlY coating cannot compensate its dissolution and spallation in sulfate salt at this case. For the duplex nanocrystalline coating, however, it combines the advantages of the nanocrystalline coating and the NiCrAlY one. The outer NiCrAlY layer obstructs the oxidation of refractory elements W, Mo and V, and ensures that only alkaline corrosion takes place at the surface of the duplex nanocrystalline coating during hot corrosion. From this point of view, the hot corrosion resistance of the duplex coating is higher than the nanocrystalline one. As compared to the NiCrAlY coating, the thick inner nanocrystalline layer of the duplex nanocrytalline coating can release most of the internal stresses by plastic deformation as discussed in Section 4.2, which ensures that the protective oxide scale formed on the duplex coating has a higher resistance to spalling off. Without spallation, the oxide scale is consumed only by dissolution in salt, whose rate is lower than its formation rate. So the duplex nanocrystalline coat-

11

ing can be protected by its oxide scale during the whole corrosion stage of 100 h. This is agreed with the corrosion kinetic curve. As shown in Fig. 10, its corrosion rate is extremely low. 5. Conclusions From the above study, the following conclusions can be drawn: (1) A duplex nanocrystalline coating is prepared by magnetron sputtering and arc ion plating. It is composed of a thick inner nanocrystalline layer and a thin outer NiCrAlY layer. (2) The duplex nanocrystalline coating combines the advantages of the traditional nanocrystalline and the NiCrAlY coatings. It possesses high resistance to oxidation, scale spallation and hot corrosion simultaneously. (3) The duplex nanocrystalline coating overcomes the shortcomings of the traditional nanocrystalline and the NiCrAlY coatings. It avoids element interdiffusion that has disturbed the NiCrAlY coating, and obstructs inward diffusion of corrosive media, such as S and O, along grain boundaries, which has limited the application of the nanocrystalline coating. Acknowledgements This project is financially supported by the National High Technology Research and Development Program of China (863 Program, No. 2012AA03A512), and by the National Key Basic Research Program of China (973 Program, No. 2012CB625100) References [1] T.N. Rhys-Jones, Coatings for blade and vane applications in gas turbines, Corros. Sci. 29 (1989) 623–646. [2] S. Nath, I. Manna, J.D. Majumdar, Kinetics and mechanism of isothermal oxidation of compositionally graded yttria stabilized zirconia (YSZ) based thermal barrier coating, Corros. Sci. 88 (2014) 10–22. [3] X. Chen, Y. Zhao, L. Gu, B. Zou, Y. Wang, X. Cao, Hot corrosion behaviour of plasma sprayed YSZ/LaMgAl11 O19 composite coatings in molten sulfate–vanadate salt, Corros.Sci. 53 (2011) 2335–2343. [4] M. Habibi, L. Wang, J. Liang, S. Guo, An investigation on hot corrosion behavior of YSZ-Ta2 O5 in Na2SO4+ V2O5 salt at 1100 ◦ C, Corros.Sci. 75 (2013) 409–414. [5] C. Amaya, W. Aperador, J.C. Caicedo, F.J. Espinoza-Beltran, J. Munoz-Saldana, G. Zambrano, P. Prieto, Corrosion study of alumina/yttria-stabilized zirconia (Al2O3 /YSZ) nanostructured thermal barrier coatings (TBC) exposed to high temperature treatment, Corros. Sci. 51 (2009) 2994–2999. [6] K. Shirvani, S. Mastali, A. Rashidghamat, H. Abdollahpour, The effect of siliconon thermal shock performance of aluminide-thermal barrier coatings, Corros. Sci. 75 (2013) 142–147. [7] F.H. Yuan, Z.X. Chen, Z.W. Huang, Z.G. Wang, S.J. Zhu, Oxidation behavior of thermal barrier coating with HVOF and detonation-sprayed NiCrAlY bondcoats, Corros. Sci. 50 (2008) 1608–1617. [8] N.P. Padture, M. Gell, E.H. Jordan, Thermal barrier coatings for gas-turbine engine applications, Science 296 (2002) 280–284. [9] T.N. Rhys-Jones, Coatings for blade and vane applications in gas turbines, Corros. Sci. 29 (1989) 623–646. [10] A. Ajaya, V. Raja, G. Sivakumar, S. Joshi, Hot corrosion behavior of solution precursor and atmospheric plasma sprayed thermal barrier coatings, Corros. Sci. 98 (2015) 271–279. [11] Z. Zhou, H. Guo, J. Wang, M. Abbas, S. Gong, Microstructure of oxides in thermal barrier coatings grown under dry/humid atmosphere, Corros. Sci. 53 (2011) 2630–2635. [12] C. Guo, W. Wang, Y. Cheng, S. Zhu, F. Wang, Yttria partially stabilized zirconia as diffusion barrier between NiCrAlY and Ni-base single crystal Rene N5 superalloy, Corros. Sci. 94 (2015) 122–128. [13] X. S.M. Jiang, Z.B. Peng, S.C. Bao, Q.M. Liu, J. Wang, C. Sun Gong, Preparation and hot corrosion behaviour of a MCrAlY + AlSiY composite coating, Corros. Sci. 50 (2008) 3213–3220. [14] J.R. Nicholls, N.J. Simms, W.Y. Chan, H.E. Evans, Smart overlay coatings-concept and practice, Surf. Coat. Technol. 149 (2002) 236–244. [15] K. Kawagishi, A. Sato, H. Harada, A concept for the EQ coating system for nickel-based superalloys, JOM 60 (2008) 31–35. [16] A. Sato, H. Harada, K. Kawagishi, Development of a new bond coat EQ Coating system, Metall. Mater. Trans. 37 (2006) 789–790. [17] B. Gleeson, Thermal barrier coatings for aero engine applications, J. Propul. Power 22 (2006) 375–383.

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Please cite this article in press as: L. Yang, et al., A duplex nanocrystalline coating for high-temperature applications on single-crystal superalloy, Corros. Sci. (2015), http://dx.doi.org/10.1016/j.corsci.2015.09.020