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Low-temperature formation and oxidation resistance of ultrafine aluminide coatings on Ni-base superalloy
Surface & Coatings Technology 203 (2009) 2337–2342
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s u r f c o a t
Low-temperature formation and oxidation resistance of ultrafine aluminide coatings on Ni-base superalloy Zhaolin Zhan a,⁎, Yedong He b, Li Li a, Hongxi Liu a, Yongnian Dai a a b
National Engineering Laboratory of Vacuum Metallurgy, Kunming University of Science and Technology, Kunming 650093, China Beijing Key Laboratory for Corrosion, Erosion and Surface Technology, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 28 April 2008 Accepted in revised form 25 February 2009 Available online 9 March 2009 Keywords: Aluminide coatings Low-temperature aluminizing, refining grains High-temperature oxidation Ni-based superalloy
a b s t r a c t Ultrafine aluminide coatings were successfully produced on Ni–18Fe–17Cr superalloy at 540–600 °C in a modified pack-aluminizing process. Repeated ball-impacts accelerated the formation of the aluminide coatings by a surface refining process, resulting in atomic diffusion occurring at a relatively low temperature. The effects of the operation temperature and the treatment duration on the formation of the coatings have been investigated. The coatings possessed a two-layer structure. The top layer, approximately 5 µm in thickness, exhibited equiaxial coarse grains and was dominated by NiAl3, with small amounts of Fe2Al5 and CrAl5. The bottom layer showed high density, homogeneous, ultrafine grains with diameters approximately 30–50 nm. High-temperature oxidation tests were carried out at 1000 °C. The oxidation kinetics and microstructure of the oxide scale were studied. The experimental results indicated that the coatings greatly enhanced the high-temperature oxidation resistance of Ni–18Fe–17Cr superalloy. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Aluminide coatings are frequently applied on to nickel-base superalloys to protect them against high-temperature oxidation, relying on the formation of a stable, dense, adherent and continuous α-Al2O3(alumina) layer on the surface of the nickel-base superalloy components. The pack cementation process is a widely applied technique for the production of aluminide coatings on nickel-base superalloys. These aluminide coatings are formed on the superalloys either by inward diffusion of Al or outward diffusion of Ni. Most of previous studies have concentrated on the chemical reactions of the pack powder and the products of the reactions [1–5], the microstructure of the aluminide coatings [6–12], the effects of elements on coatings [13–15] and the high-temperature properties of the coatings [16–19]. However, there are certain disadvantages to pack aluminizing, such as the sintering tendency of pack powders at high temperature, the formation of an inhomogeneous multilayer structure, coarse grains and inclusions in the coatings. It was reported [20–22] that surface micro- or nano-crystallization could effectively promote atomic diffusion and chemical reaction, due to the large number of grain boundaries acting as fast atomic diffusion channels. This provided a new option for the production of diffusing coatings. For example, pure iron has been nitrided in flowing ammonia gas (NH3) at a relatively low temperature (300 °C) after surface nanostructuring treatment [22]. However, as the grain growth of micro-
⁎ Corresponding author. Tel./fax: +86 871 5192665. E-mail address: [email protected] (Z. Zhan). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.02.127
or nano-crystallization occurred at a relatively high temperature, the atomic diffusion along grain boundaries would transform to crystal lattice diffusion. Therefore, it was difficult to utilize the advantages of surface micro- or nano-crystallization in the pack-aluminizing processes. Recently, the authors presented a new method for the production of aluminide coatings on carbon steel by combining the pack-aluminizing process with a surface micro- or nano-structuring process [23]. Ultrafine aluminide coatings were successfully produced on carbon steel at low temperatures (440–600 °C) and with a short treatment (0.25–2 h). In this paper, we have applied this technique to produce aluminide coatings on Ni–18Fe–17Cr superalloy and investigated the effects of the operational temperature and the treatment duration on the formation of the coatings, their features and the high-temperature oxidation resistance of the coatings. The further studied of the coatings features and its oxidation resistance will be carried out in future. 2. Experimental Ni–18Fe–17Cr superalloy was used as the substrate material in the experiments. The nominal composition of the alloy was 63.1Ni– 18.3Fe–16.6Cr–1.22Nb–0.78Ti (wt.%). Samples of the alloy were cut into rectangular specimens with dimensions 12 × 12 × 2 mm3. The surfaces of the specimens were polished with abrasive papers to 600grit and washed with acetone in an ultrasonic bath. The experimental apparatus has been described elsewhere in detail [23]. A retort (Φ50 × 80 mm3) was filled with the treatment agent, consisting of pack-aluminizing powder (40 g in total, pure Al powder: 10–30 wt.%, 60 µm, α-Al2O3 filler: 60–80 wt.%, 100 µm and
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NH4Cl activator: 10 wt.%), Ni-base alloy specimens and FeCrAl alloy balls (4 mm in diameter). The retort was heated to 540, 560, 580 and 600 °C in a furnace and vibrated by a mechanical vibrator at a frequency of 25 Hz. The pack-aluminizing process took place as a result of the repeated impact of the alloy balls. Higher Al and activator contents were employed in the process. The major precursor species was found to be AlCl3 (g) [24–26]. A higher AlCl3 pressure was achieved and NH3 would drive the air out off the retort to avoid oxidation of specimens. After deposition of the modified pack-aluminizing coatings, the effects of a subsequent heat treatment on the evolution of coating thickness and the resulting phases were also studied. The heat treatment was carried out at 850 °C/4.5 h followed by natural cooling to room temperature under pure Ar gas. High-temperature oxidation tests were carried out in a horizontal tube furnace at 1000 °C with flowing air. The specimens were treated in the modified pack-aluminizing process at 560 °C for 3 h and the thickness of the Al-coatings was approximately 20 µm. Two Al-coated specimens were employed for each oxidation time. The mass changes and the spallations were measured to evaluate the performance of the high-temperature oxidation resistance of the coatings. Surface and cross-sectional observation of the specimens was preformed on a scanning electron microscope (SEM) and an atomic force microscopy (AFM) with contact mode. The composition was analyzed by X-ray energy dispersive spectroscopy (EDS). X-ray diffraction (XRD) analysis of the surface layer in the specimens was carried out with 12 kW, CuKα radiation of λ = 0.15405 nm, and small steps of 2θ = 0.02°. 3. Results 3.1. Effects on the formation of the aluminide coatings
Fig. 2. Variation of the thickness of the coatings with treatment duration (treated at 560 °C).
The variation of the thickness of the coatings with the treatment duration at 560 °C is shown in Fig. 2. Aluminide coatings could be rapidly produced on the Ni-base superalloy. A 7 µm aluminide coating was formed at 560 °C with treatment for only 1 h. The thickness of the coating reached 20 µm as the duration of treatment was extended to 2 h, but remained unchanged when further increasing the duration of treatment to 4 h. 3.2. Characterization of the aluminide coatings Fig. 3 displays the cross-sectional microstructure and the composition distribution of the coated specimen treated at 560 °C for 3 h. The aluminide coatings had a two-layer structure. The top layer was
Fig. 1 illustrates the variation of the thickness of the coatings on the Ni-base superalloy as a function of the operational temperature. 5–30 µm thick aluminide coatings could be formed at relatively low temperatures in the range from 540 °C to 600 °C. It is particularly important to note that the temperature required to form the aluminide coatings in our experiments is apparently lower than that for conventional pack aluminizing on a Ni-base superalloy, which are generally above 760 °C. The thicknesses of the aluminide coatings increased with enhancement of the operation temperature. 5–9 µm thick aluminide coatings were produced with treatment at 540 °C for 3 h. When the operational temperature was raised to 560 °C, the coating was about 19 µm in thickness. The thicknesses of the coatings were about 23 µm at 580 °C and 28 µm at 600 °C, respectively.
Fig. 1. Variation of the thickness of the coatings with operational temperature (treated for 3 h).
Fig. 3. (a) Cross-sectional microstructure and (b) composition distribution of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
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Fig. 4. AFM micrographs of the cross-section in the bottom layer of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
approximately 5 µm thick and exhibits coarse grains with high nickel and aluminium contents (∼47 wt.% and ∼50 wt.% respectively). The contents of iron and chromium in the top layer were ∼ 3 wt.% and ∼ 1 wt.% respectively (see Fig. 3b). It can be deduced that the top layer of the coating was dominated by Al-rich phases such as NiAl3, with a small amount of Al–Fe and Al–Cr intermetallics. The bottom layer of the coating was about 20 µm in thickness. It had high density, being homogeneous and free of coarse-grain, columnar-crystal and precipitates, which were usually observed in the aluminide coatings on Ni-base superalloy from conventional packaluminizing processes. In order to analyze the microstructure of the coating in detail, atomic force microscopy (AFM) was applied to observe the coating. The AFM micrograph is displayed in Fig. 4. The grains in the bottom layer appeared to be of equiaxial shape, with random orientations and proportional distribution. The grain size was
Fig. 6. (a) SEM micrograph and (b) XRD pattern from the top of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
approximately 30–50 nm in diameter, indicating that the grain growth of the coating was greatly hindered in the treatment process even at high temperature.
Fig. 5. Elements distribution of EDS analysis of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
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The elemental content changed sharply across the interface between the top layer and the bottom layer. In the bottom layer, the aluminium content close to the top layer was up to ∼57 wt.% and then gradually decreased across the layer, to ∼46 wt.% near the substrate. In contrast, the nickel content increased from ∼18 wt.% to ∼ 32 wt.% across the bottom layer. However, the iron and chromium contents remained constant, being ∼ 10.5 wt.% and ∼ 9.5 wt.% respectively. The accurate EDS analysis demonstrated that the matrix of the bottom layer consisted of Al-rich phases, namely NiAl3. Fe2Al5 and CrAl5 phases were also formed, as the contents of iron and chromium in the bottom layer were higher than their solubilities in the matrix. Fig. 5 displays X-ray maps of selected elements from the EDS analysis of the coating, showing uniform distribution of Al, Ni, Fe, Cr and Nb. Fig. 6 is the top micrograph and the corresponding XRD pattern of the aluminide coating on the Ni-base superalloy. The top layer showed an ultrafine grain microstructure and the dent induced by the ball impact was apparent (the bright zone in Fig. 6a). XRD pattern identified that the coating consisted mainly of NiAl3, with small amounts of Fe2Al5 and CrAl5 (Fig. 6b), supported by the EDS results. 3.3. Heat treatment of the aluminide coatings The coated specimens were heated to 850 °C and held at this temperature for 4.5 h. The heat treatment brought about transport of the different elements, which in turn gave rise to thicker coatings and a change of phases. A typical coating morphology and the composition distribution across the coating are shown in Fig. 7. The coating possessed a two-layer structure. The top grey layer was smooth in morphology. Across top layer, steady Al, Ni, and Fe contents were identified by EDS analysis [see Fig. 7b], being ∼44 wt.%, ∼ 42 wt.% and ∼ 10 wt.% respectively. The top layer might grow by countercurrent Fig. 8. (a) Oxidation kinetics and (b) spallation of specimens exposed to air at 1000 °C for 200 h.
diffusion of aluminium and iron [26]. Underneath the top layer, the bottom layer was composed of a matrix of two phases (NiAl and FeAl) containing different precipitates. EDS analysis (see Fig. 7b) indicated that all elements in the substrate alloy were present throughout the coating in relative amounts corresponding to their concentrations in the alloy, and showed a gradient distribution across the coating. Absolute concentrations of these elements were lower than the substrate concentrations in proportion to dilution by aluminium in the coating. According to XRD and EDS analysis, the precipitated particles were typically constituted by Cr5Al8 and α-Cr. Chromium is known to be insoluble in Ni–Al phases [27] and thus, upon formation of NiAl it precipitated as α-Cr. The chromium showed a high concentration in the substrate beneath the coating. The most likely origin for this was that chromium had been pushed into the interface because of the low solubility of chromium in the matrix (NiAl) after heat treatment and formed a CrNi barrier layer against aluminium diffusion into the substrate. 3.4. High-temperature oxidation resistance
Fig. 7. (a) Cross-sectional microstructure and (b) EDS analysis of the aluminide coating on Ni-base superalloy followed by heat treatment at 850 °C for 4.5 h.
The oxidation kinetic curves and the spallation of specimens, exposed in air at 1000 °C for 200 h, are illustrated in Fig. 8. Ni-base superalloy showed rapid oxidation at 1000 °C. The mass gain and the spallation of the Ni-base superalloy were up to ∼20.7 g/m2 and ∼62.4 g/m2 respectively, after exposure to air at 1000 °C for 200 h. In the case of oxidation at 1000 °C, Cr2O3 was formed on the surface of Ni-base superalloy and transformed to CrO3 with further oxidation. CrO3 would volatile at temperature of above 950 °C, causing Cr loss from the outer-oxide layer [28,29], showing severe oxidation.
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medium) formed by inward diffusion of aluminium through the surface of the substrate, causing a phase transformation to brittle δ(Ni2Al3). Subsequent heat treatment caused Ni to diffuse outwards forming a three-zone β(NiAl) coating. The other is low activity coating (using alloyed-aluminium powder in the coating medium) formed by outward diffusion of nickel from the substrate, producing a two-zone β(NiAl) coating. In our experiments, pure aluminium powder was used in the packaluminizing process on Ni-base superalloy (high activity aluminizing). However, the aluminide coatings were two-layer, exhibiting the same features as the low activity coating. Furthermore, the coatings obtained in our experiments are in sharp contrast to those formed by the conventional pack-aluminizing method in terms of their microstructure, elemental concentration, phases of the coating and the operational temperature. These results indicated that the formation mechanism of the coating was different to that of the conventional pack aluminization of Ni-base superalloy. The formation of the ultrafine aluminide coatings can be attributed to the repeated ball-impact. Combined with the formation mechanism of aluminide coatings on carbon steel [23], the formation processes of aluminide coatings on Ni-base superalloy are described as follows:
Fig. 9. Cross-section of oxide scale exposed to air at 1000 °C for 200 h (a) the uncoated specimen and (b) the coated specimen.
The coated specimens possessed excellent oxidation resistance. The mass gain of the coated specimen was 5.6 g/m2, approximately 1/ 4 of the uncoated specimen. The coated specimens showed a slight spallation of 5.1 g/m2, being approximate 8% of the uncoated specimen. This indicated that Al2O3 scale was formed on the surface of coated specimens and protected alloy further oxidized at 1000 °C [30,31]. Fig. 9 displays the cross-section of the oxide scale of specimens after oxidation in air at 1000 °C for 200 h. In order to protect the oxide scale, Ni was plated on to the surface of the specimens. As can be seen, a discontinuous and porous oxide scale was formed on the uncoated specimen (see Fig. 9a). Beneath the oxide scale, holes and pores were apparently observed, evidence of oxidation in the substrate. It was due to the fact that the coarse grains of the substrate and oxygen atoms went through the grain boundaries easily, causing intergranular internal oxidation. However, a protective α-Al2O3 scale was formed on the coated specimen after exposure to air at 1000 °C for 200 h (see Fig. 9b). The oxide scale was dense and continuous. Also the substrate was uniform, dense and free of pores. The results showed that the substrate had been protected well from further oxidation in a high-temperature environment. The top views of the uncoated and coated specimens are displayed in Fig. 10, after exposure to air at 1000 °C for 200 h. The surface of oxide scale formed on the uncoated specimen consisted of equiaxial grains with a typical size of approximately 2–3 µm, and which were rough and porous (see Fig. 10a). The coated specimen showed dense and homogeneous alumina scale morphology with a slight spallation (see Fig. 10b).
1) Ultrafine grains were produced on the top layer of the metals by repeated ball-impact [32–34]. In the modified aluminizing process, the repeated ball-impact caused the surface of the substrate to undergo micro- and nano-structuring and pinned the pure aluminium particles into the coating medium adhering on to the surface of the substrate, producing a pure aluminium layer. 2) Atomic diffusion occurred at 540–600 °C as the large amount of grain boundaries and defects induced by the repeated ball-impact were acting as the fast atomic diffusion channel, producing an initial alloy layer. 3) Subsequent ball-impact refined the initial alloy layer and caused the homogenization and densification of the coatings. At the same
4. Discussion Generally, two types of aluminide coatings are formed on Ni-base superalloys using the pack aluminide processes [12]. One is referred to as high activity coating (using aluminium powder in the coating
Fig. 10. SEM micrograph of (a) the uncoated specimen and (b) the coated specimen exposed to air at 1000 °C for 200 h.
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time, another pure aluminium layer was produced on the surface layer. 4) As the coating grew, the outward diffusion rates of the substrate elements gradually decreased. The surface layer of the coating was broken by the ball impact, hindering further atomic diffusion. Therefore, the growth of the coating was suspended even though the treatment continued, resulting in a two-layer coating. The oxidation resistance of the coated specimens was greatly enhanced due to the formation of α-Al2O3 scale instead of Cr2O3 scale. In high-temperature environment, Cr2O3 was firstly produced on Nibase superalloy and then transformed to CrO3 with further oxidation. CrO3 would be volatile at a temperature of above 950 °C, losing prevention for oxygen diffusion and causing further oxidation. αAl2O3 scale, however, showed excellent oxidation resistance at a temperature of above 1000 °C. And ultrafine Fe–Al coatings would induce formation of uniform and dense α-Al2O3 scale, exhibiting strong prevention for oxygen diffusing into the substrate. 5. Conclusions a) Ultrafine aluminide coatings were produced on Ni–18Fe–17Cr superalloy at low temperatures (540–600 °C) with treatment of short duration (1–4 h) in a modified pack-aluminizing process. Repeat ball-impact accelerated the formation of the aluminide coatings by surface micro- and nano-structuring, resulting in atomic diffusion occurring at a relatively low temperature. b) The coatings possessed a two-layer structure. The top layer showed equiaxial coarse grains with high nickel and aluminium contents, consisting of NiAl3 and small amounts of Fe2Al5 and CrAl5. The bottom layer was of high density and homogeneity, with ultrafine equiaxial grains 30–50 nm in diameter, and consisted of Al-rich phases. c) A subsequent heat treatment up to 850 °C thickens the coatings and brings about the formation of two-layer coating. In the outer layer NiAl and FeAl predominates whereas in the bottom layer a twophase matrix is present containing various chromium-containing precipitates. d) The aluminide coatings apparently improved the high-temperature oxidation resistance of Ni–18Fe–17Cr superalloy. After oxidization in air at 1000 °C for 200 h, a full, dense and continuous α-Al2O3 scale was formed on the surface of the coated specimen.
The mass gain and the spallation of the coated specimen were around 25% and 8%, respectively, of the uncoated specimen. Acknowledgments The authors gratefully acknowledge the support of the Chinese National Science Foundation (Grant 50671045) and Yunnan Science Foundation (Grant 2007PY01-05). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]
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Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s u r f c o a t
Low-temperature formation and oxidation resistance of ultrafine aluminide coatings on Ni-base superalloy Zhaolin Zhan a,⁎, Yedong He b, Li Li a, Hongxi Liu a, Yongnian Dai a a b
National Engineering Laboratory of Vacuum Metallurgy, Kunming University of Science and Technology, Kunming 650093, China Beijing Key Laboratory for Corrosion, Erosion and Surface Technology, University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 28 April 2008 Accepted in revised form 25 February 2009 Available online 9 March 2009 Keywords: Aluminide coatings Low-temperature aluminizing, refining grains High-temperature oxidation Ni-based superalloy
a b s t r a c t Ultrafine aluminide coatings were successfully produced on Ni–18Fe–17Cr superalloy at 540–600 °C in a modified pack-aluminizing process. Repeated ball-impacts accelerated the formation of the aluminide coatings by a surface refining process, resulting in atomic diffusion occurring at a relatively low temperature. The effects of the operation temperature and the treatment duration on the formation of the coatings have been investigated. The coatings possessed a two-layer structure. The top layer, approximately 5 µm in thickness, exhibited equiaxial coarse grains and was dominated by NiAl3, with small amounts of Fe2Al5 and CrAl5. The bottom layer showed high density, homogeneous, ultrafine grains with diameters approximately 30–50 nm. High-temperature oxidation tests were carried out at 1000 °C. The oxidation kinetics and microstructure of the oxide scale were studied. The experimental results indicated that the coatings greatly enhanced the high-temperature oxidation resistance of Ni–18Fe–17Cr superalloy. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Aluminide coatings are frequently applied on to nickel-base superalloys to protect them against high-temperature oxidation, relying on the formation of a stable, dense, adherent and continuous α-Al2O3(alumina) layer on the surface of the nickel-base superalloy components. The pack cementation process is a widely applied technique for the production of aluminide coatings on nickel-base superalloys. These aluminide coatings are formed on the superalloys either by inward diffusion of Al or outward diffusion of Ni. Most of previous studies have concentrated on the chemical reactions of the pack powder and the products of the reactions [1–5], the microstructure of the aluminide coatings [6–12], the effects of elements on coatings [13–15] and the high-temperature properties of the coatings [16–19]. However, there are certain disadvantages to pack aluminizing, such as the sintering tendency of pack powders at high temperature, the formation of an inhomogeneous multilayer structure, coarse grains and inclusions in the coatings. It was reported [20–22] that surface micro- or nano-crystallization could effectively promote atomic diffusion and chemical reaction, due to the large number of grain boundaries acting as fast atomic diffusion channels. This provided a new option for the production of diffusing coatings. For example, pure iron has been nitrided in flowing ammonia gas (NH3) at a relatively low temperature (300 °C) after surface nanostructuring treatment [22]. However, as the grain growth of micro-
⁎ Corresponding author. Tel./fax: +86 871 5192665. E-mail address: [email protected] (Z. Zhan). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.02.127
or nano-crystallization occurred at a relatively high temperature, the atomic diffusion along grain boundaries would transform to crystal lattice diffusion. Therefore, it was difficult to utilize the advantages of surface micro- or nano-crystallization in the pack-aluminizing processes. Recently, the authors presented a new method for the production of aluminide coatings on carbon steel by combining the pack-aluminizing process with a surface micro- or nano-structuring process [23]. Ultrafine aluminide coatings were successfully produced on carbon steel at low temperatures (440–600 °C) and with a short treatment (0.25–2 h). In this paper, we have applied this technique to produce aluminide coatings on Ni–18Fe–17Cr superalloy and investigated the effects of the operational temperature and the treatment duration on the formation of the coatings, their features and the high-temperature oxidation resistance of the coatings. The further studied of the coatings features and its oxidation resistance will be carried out in future. 2. Experimental Ni–18Fe–17Cr superalloy was used as the substrate material in the experiments. The nominal composition of the alloy was 63.1Ni– 18.3Fe–16.6Cr–1.22Nb–0.78Ti (wt.%). Samples of the alloy were cut into rectangular specimens with dimensions 12 × 12 × 2 mm3. The surfaces of the specimens were polished with abrasive papers to 600grit and washed with acetone in an ultrasonic bath. The experimental apparatus has been described elsewhere in detail [23]. A retort (Φ50 × 80 mm3) was filled with the treatment agent, consisting of pack-aluminizing powder (40 g in total, pure Al powder: 10–30 wt.%, 60 µm, α-Al2O3 filler: 60–80 wt.%, 100 µm and
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NH4Cl activator: 10 wt.%), Ni-base alloy specimens and FeCrAl alloy balls (4 mm in diameter). The retort was heated to 540, 560, 580 and 600 °C in a furnace and vibrated by a mechanical vibrator at a frequency of 25 Hz. The pack-aluminizing process took place as a result of the repeated impact of the alloy balls. Higher Al and activator contents were employed in the process. The major precursor species was found to be AlCl3 (g) [24–26]. A higher AlCl3 pressure was achieved and NH3 would drive the air out off the retort to avoid oxidation of specimens. After deposition of the modified pack-aluminizing coatings, the effects of a subsequent heat treatment on the evolution of coating thickness and the resulting phases were also studied. The heat treatment was carried out at 850 °C/4.5 h followed by natural cooling to room temperature under pure Ar gas. High-temperature oxidation tests were carried out in a horizontal tube furnace at 1000 °C with flowing air. The specimens were treated in the modified pack-aluminizing process at 560 °C for 3 h and the thickness of the Al-coatings was approximately 20 µm. Two Al-coated specimens were employed for each oxidation time. The mass changes and the spallations were measured to evaluate the performance of the high-temperature oxidation resistance of the coatings. Surface and cross-sectional observation of the specimens was preformed on a scanning electron microscope (SEM) and an atomic force microscopy (AFM) with contact mode. The composition was analyzed by X-ray energy dispersive spectroscopy (EDS). X-ray diffraction (XRD) analysis of the surface layer in the specimens was carried out with 12 kW, CuKα radiation of λ = 0.15405 nm, and small steps of 2θ = 0.02°. 3. Results 3.1. Effects on the formation of the aluminide coatings
Fig. 2. Variation of the thickness of the coatings with treatment duration (treated at 560 °C).
The variation of the thickness of the coatings with the treatment duration at 560 °C is shown in Fig. 2. Aluminide coatings could be rapidly produced on the Ni-base superalloy. A 7 µm aluminide coating was formed at 560 °C with treatment for only 1 h. The thickness of the coating reached 20 µm as the duration of treatment was extended to 2 h, but remained unchanged when further increasing the duration of treatment to 4 h. 3.2. Characterization of the aluminide coatings Fig. 3 displays the cross-sectional microstructure and the composition distribution of the coated specimen treated at 560 °C for 3 h. The aluminide coatings had a two-layer structure. The top layer was
Fig. 1 illustrates the variation of the thickness of the coatings on the Ni-base superalloy as a function of the operational temperature. 5–30 µm thick aluminide coatings could be formed at relatively low temperatures in the range from 540 °C to 600 °C. It is particularly important to note that the temperature required to form the aluminide coatings in our experiments is apparently lower than that for conventional pack aluminizing on a Ni-base superalloy, which are generally above 760 °C. The thicknesses of the aluminide coatings increased with enhancement of the operation temperature. 5–9 µm thick aluminide coatings were produced with treatment at 540 °C for 3 h. When the operational temperature was raised to 560 °C, the coating was about 19 µm in thickness. The thicknesses of the coatings were about 23 µm at 580 °C and 28 µm at 600 °C, respectively.
Fig. 1. Variation of the thickness of the coatings with operational temperature (treated for 3 h).
Fig. 3. (a) Cross-sectional microstructure and (b) composition distribution of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
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Fig. 4. AFM micrographs of the cross-section in the bottom layer of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
approximately 5 µm thick and exhibits coarse grains with high nickel and aluminium contents (∼47 wt.% and ∼50 wt.% respectively). The contents of iron and chromium in the top layer were ∼ 3 wt.% and ∼ 1 wt.% respectively (see Fig. 3b). It can be deduced that the top layer of the coating was dominated by Al-rich phases such as NiAl3, with a small amount of Al–Fe and Al–Cr intermetallics. The bottom layer of the coating was about 20 µm in thickness. It had high density, being homogeneous and free of coarse-grain, columnar-crystal and precipitates, which were usually observed in the aluminide coatings on Ni-base superalloy from conventional packaluminizing processes. In order to analyze the microstructure of the coating in detail, atomic force microscopy (AFM) was applied to observe the coating. The AFM micrograph is displayed in Fig. 4. The grains in the bottom layer appeared to be of equiaxial shape, with random orientations and proportional distribution. The grain size was
Fig. 6. (a) SEM micrograph and (b) XRD pattern from the top of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
approximately 30–50 nm in diameter, indicating that the grain growth of the coating was greatly hindered in the treatment process even at high temperature.
Fig. 5. Elements distribution of EDS analysis of the aluminide coating on the Ni-base superalloy (treated at 560 °C for 3 h).
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The elemental content changed sharply across the interface between the top layer and the bottom layer. In the bottom layer, the aluminium content close to the top layer was up to ∼57 wt.% and then gradually decreased across the layer, to ∼46 wt.% near the substrate. In contrast, the nickel content increased from ∼18 wt.% to ∼ 32 wt.% across the bottom layer. However, the iron and chromium contents remained constant, being ∼ 10.5 wt.% and ∼ 9.5 wt.% respectively. The accurate EDS analysis demonstrated that the matrix of the bottom layer consisted of Al-rich phases, namely NiAl3. Fe2Al5 and CrAl5 phases were also formed, as the contents of iron and chromium in the bottom layer were higher than their solubilities in the matrix. Fig. 5 displays X-ray maps of selected elements from the EDS analysis of the coating, showing uniform distribution of Al, Ni, Fe, Cr and Nb. Fig. 6 is the top micrograph and the corresponding XRD pattern of the aluminide coating on the Ni-base superalloy. The top layer showed an ultrafine grain microstructure and the dent induced by the ball impact was apparent (the bright zone in Fig. 6a). XRD pattern identified that the coating consisted mainly of NiAl3, with small amounts of Fe2Al5 and CrAl5 (Fig. 6b), supported by the EDS results. 3.3. Heat treatment of the aluminide coatings The coated specimens were heated to 850 °C and held at this temperature for 4.5 h. The heat treatment brought about transport of the different elements, which in turn gave rise to thicker coatings and a change of phases. A typical coating morphology and the composition distribution across the coating are shown in Fig. 7. The coating possessed a two-layer structure. The top grey layer was smooth in morphology. Across top layer, steady Al, Ni, and Fe contents were identified by EDS analysis [see Fig. 7b], being ∼44 wt.%, ∼ 42 wt.% and ∼ 10 wt.% respectively. The top layer might grow by countercurrent Fig. 8. (a) Oxidation kinetics and (b) spallation of specimens exposed to air at 1000 °C for 200 h.
diffusion of aluminium and iron [26]. Underneath the top layer, the bottom layer was composed of a matrix of two phases (NiAl and FeAl) containing different precipitates. EDS analysis (see Fig. 7b) indicated that all elements in the substrate alloy were present throughout the coating in relative amounts corresponding to their concentrations in the alloy, and showed a gradient distribution across the coating. Absolute concentrations of these elements were lower than the substrate concentrations in proportion to dilution by aluminium in the coating. According to XRD and EDS analysis, the precipitated particles were typically constituted by Cr5Al8 and α-Cr. Chromium is known to be insoluble in Ni–Al phases [27] and thus, upon formation of NiAl it precipitated as α-Cr. The chromium showed a high concentration in the substrate beneath the coating. The most likely origin for this was that chromium had been pushed into the interface because of the low solubility of chromium in the matrix (NiAl) after heat treatment and formed a CrNi barrier layer against aluminium diffusion into the substrate. 3.4. High-temperature oxidation resistance
Fig. 7. (a) Cross-sectional microstructure and (b) EDS analysis of the aluminide coating on Ni-base superalloy followed by heat treatment at 850 °C for 4.5 h.
The oxidation kinetic curves and the spallation of specimens, exposed in air at 1000 °C for 200 h, are illustrated in Fig. 8. Ni-base superalloy showed rapid oxidation at 1000 °C. The mass gain and the spallation of the Ni-base superalloy were up to ∼20.7 g/m2 and ∼62.4 g/m2 respectively, after exposure to air at 1000 °C for 200 h. In the case of oxidation at 1000 °C, Cr2O3 was formed on the surface of Ni-base superalloy and transformed to CrO3 with further oxidation. CrO3 would volatile at temperature of above 950 °C, causing Cr loss from the outer-oxide layer [28,29], showing severe oxidation.
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medium) formed by inward diffusion of aluminium through the surface of the substrate, causing a phase transformation to brittle δ(Ni2Al3). Subsequent heat treatment caused Ni to diffuse outwards forming a three-zone β(NiAl) coating. The other is low activity coating (using alloyed-aluminium powder in the coating medium) formed by outward diffusion of nickel from the substrate, producing a two-zone β(NiAl) coating. In our experiments, pure aluminium powder was used in the packaluminizing process on Ni-base superalloy (high activity aluminizing). However, the aluminide coatings were two-layer, exhibiting the same features as the low activity coating. Furthermore, the coatings obtained in our experiments are in sharp contrast to those formed by the conventional pack-aluminizing method in terms of their microstructure, elemental concentration, phases of the coating and the operational temperature. These results indicated that the formation mechanism of the coating was different to that of the conventional pack aluminization of Ni-base superalloy. The formation of the ultrafine aluminide coatings can be attributed to the repeated ball-impact. Combined with the formation mechanism of aluminide coatings on carbon steel [23], the formation processes of aluminide coatings on Ni-base superalloy are described as follows:
Fig. 9. Cross-section of oxide scale exposed to air at 1000 °C for 200 h (a) the uncoated specimen and (b) the coated specimen.
The coated specimens possessed excellent oxidation resistance. The mass gain of the coated specimen was 5.6 g/m2, approximately 1/ 4 of the uncoated specimen. The coated specimens showed a slight spallation of 5.1 g/m2, being approximate 8% of the uncoated specimen. This indicated that Al2O3 scale was formed on the surface of coated specimens and protected alloy further oxidized at 1000 °C [30,31]. Fig. 9 displays the cross-section of the oxide scale of specimens after oxidation in air at 1000 °C for 200 h. In order to protect the oxide scale, Ni was plated on to the surface of the specimens. As can be seen, a discontinuous and porous oxide scale was formed on the uncoated specimen (see Fig. 9a). Beneath the oxide scale, holes and pores were apparently observed, evidence of oxidation in the substrate. It was due to the fact that the coarse grains of the substrate and oxygen atoms went through the grain boundaries easily, causing intergranular internal oxidation. However, a protective α-Al2O3 scale was formed on the coated specimen after exposure to air at 1000 °C for 200 h (see Fig. 9b). The oxide scale was dense and continuous. Also the substrate was uniform, dense and free of pores. The results showed that the substrate had been protected well from further oxidation in a high-temperature environment. The top views of the uncoated and coated specimens are displayed in Fig. 10, after exposure to air at 1000 °C for 200 h. The surface of oxide scale formed on the uncoated specimen consisted of equiaxial grains with a typical size of approximately 2–3 µm, and which were rough and porous (see Fig. 10a). The coated specimen showed dense and homogeneous alumina scale morphology with a slight spallation (see Fig. 10b).
1) Ultrafine grains were produced on the top layer of the metals by repeated ball-impact [32–34]. In the modified aluminizing process, the repeated ball-impact caused the surface of the substrate to undergo micro- and nano-structuring and pinned the pure aluminium particles into the coating medium adhering on to the surface of the substrate, producing a pure aluminium layer. 2) Atomic diffusion occurred at 540–600 °C as the large amount of grain boundaries and defects induced by the repeated ball-impact were acting as the fast atomic diffusion channel, producing an initial alloy layer. 3) Subsequent ball-impact refined the initial alloy layer and caused the homogenization and densification of the coatings. At the same
4. Discussion Generally, two types of aluminide coatings are formed on Ni-base superalloys using the pack aluminide processes [12]. One is referred to as high activity coating (using aluminium powder in the coating
Fig. 10. SEM micrograph of (a) the uncoated specimen and (b) the coated specimen exposed to air at 1000 °C for 200 h.
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time, another pure aluminium layer was produced on the surface layer. 4) As the coating grew, the outward diffusion rates of the substrate elements gradually decreased. The surface layer of the coating was broken by the ball impact, hindering further atomic diffusion. Therefore, the growth of the coating was suspended even though the treatment continued, resulting in a two-layer coating. The oxidation resistance of the coated specimens was greatly enhanced due to the formation of α-Al2O3 scale instead of Cr2O3 scale. In high-temperature environment, Cr2O3 was firstly produced on Nibase superalloy and then transformed to CrO3 with further oxidation. CrO3 would be volatile at a temperature of above 950 °C, losing prevention for oxygen diffusion and causing further oxidation. αAl2O3 scale, however, showed excellent oxidation resistance at a temperature of above 1000 °C. And ultrafine Fe–Al coatings would induce formation of uniform and dense α-Al2O3 scale, exhibiting strong prevention for oxygen diffusing into the substrate. 5. Conclusions a) Ultrafine aluminide coatings were produced on Ni–18Fe–17Cr superalloy at low temperatures (540–600 °C) with treatment of short duration (1–4 h) in a modified pack-aluminizing process. Repeat ball-impact accelerated the formation of the aluminide coatings by surface micro- and nano-structuring, resulting in atomic diffusion occurring at a relatively low temperature. b) The coatings possessed a two-layer structure. The top layer showed equiaxial coarse grains with high nickel and aluminium contents, consisting of NiAl3 and small amounts of Fe2Al5 and CrAl5. The bottom layer was of high density and homogeneity, with ultrafine equiaxial grains 30–50 nm in diameter, and consisted of Al-rich phases. c) A subsequent heat treatment up to 850 °C thickens the coatings and brings about the formation of two-layer coating. In the outer layer NiAl and FeAl predominates whereas in the bottom layer a twophase matrix is present containing various chromium-containing precipitates. d) The aluminide coatings apparently improved the high-temperature oxidation resistance of Ni–18Fe–17Cr superalloy. After oxidization in air at 1000 °C for 200 h, a full, dense and continuous α-Al2O3 scale was formed on the surface of the coated specimen.
The mass gain and the spallation of the coated specimen were around 25% and 8%, respectively, of the uncoated specimen. Acknowledgments The authors gratefully acknowledge the support of the Chinese National Science Foundation (Grant 50671045) and Yunnan Science Foundation (Grant 2007PY01-05). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]
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