The mechanism for self-formation of a CeO2 diffusion barrier layer in an aluminide coating at high temperature

Surface & Coatings Technology 224 (2013) 62–70

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The mechanism for self-formation of a CeO2 diffusion barrier layer in an aluminide coating at high temperature X. Tan, X. Peng ⁎, F. Wang State Key laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 62 Wencui Road, Shenyang 110016, China

a r t i c l e

i n f o

Article history: Received 31 December 2012 Accepted in revised form 4 March 2013 Available online 14 March 2013 Keywords: Aluminide coating Annealing Diffusion barrier Electroplating Interdiffusion

a b s t r a c t A CeO2 dispersed δ-Ni2Al3 was formed by partially aluminizing an electrodeposited Ni film containing CeO2. The aluminide/Ni–CeO2 coating system itself quickly formed a CeO2-rich diffusion barrier between aluminide and Ni during annealing in vacuum at 1000 °C. A model for the formation of the diffusion barrier was proposed, based on the characterization of the evolution with time of the phase compositions of the aluminide at the interface. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Diffusion aluminide coatings which were introduced into service around 1950 [1,2] are the most widely used in industry as the first generation of environmental-protection coatings [3,4], because they thermally grow a protective scale of alumina. Works on the high temperature oxidation behavior of the diffusion coatings [5,6] and the optimization of the aluminizing processes (e.g., adding Pt [7–9] and rare earth elements [10–13], decreasing aluminizing temperature [14–16]) to further improve the oxidation resistance have been extensively reported. However, in a practical high temperature atmosphere, the aluminide coating was degraded, not only by oxidation but also severely by the interdiffusion between the aluminide and the underlying metallic substrate [17–19]. For example, the coatings of β-NiAl would degrade into the γ′-Ni3Al by its interdiffusion with the Ni-based superalloys [19–21]. Moreover, the diffusion of Al into the alloy would increase volume fraction of the Al-rich intermetallic compounds (normally in β and γ′ phases [22]), which, in turn, facilitates the precipitation of the topologically close packed (TCP) phases containing alloy components (such as W, Re, Mo, Cr) because of their lower solubility in Al-rich intermetallics than in γ-Ni [23]. For a similar reason, the penetration of the above alloy components into the β-NiAl coatings by the interdiffusion also precipitates numerous TCPs [21,22,24]. It is well known that the TCPs are brittle and their precipitation is always harmful to the mechanical properties of both alloys and coatings [25,26]. In addition, the oxidation resistance of the aluminide coatings was decreased by the diffusion of some elements (such as W, Ti, and Mo) from the superalloy substrate [27–29]. Single ⁎ Corresponding author. Tel.: +86 24 23893753; fax: +86 24 23893624. E-mail address: [email protected] (X. Peng). 0257-8972/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.surfcoat.2013.03.003

crystal (SC) Ni-base super alloys have been extensively applied due to the greatly increased mechanical properties in comparison with normal polycrystalline alloys. However, the penetration of Al from the aluminide coating into the high-Re SC alloy (such as those of third- [22,30] and fourth-generation [31–33]) substrates due to the interdiffusion destabilizes the γ/γ′ microstructure in a certain depth zone of the alloy and finally converts the phases into a γ′ matrix with the γ and TCP needle-like precipitates in equilibrium state. The phase-transformed zone is the so-called second reaction zone with poor mechanical properties [34,35]. Therefore, to develop a layer for blocking the interdiffusion between aluminide coatings and metal substrates is currently a concern to extend the coating service life. A good diffusion barrier always has a low diffusion coefficient of Al or elements from the metal substrates, a high thermodynamical stability at high temperatures and a desired compatibility with both coating and alloy [36]. So far, the diffusion barriers developed for the aluminide coating/metallic substrate systems are usually metallic based coatings. The typical of them are Re–Cr–Ni [36–38], Ni–W [29], Co–Ru [31] and Hf [39]. These metallic diffusion barrier coatings are normally pre-deposited onto the metal substrates for aluminizing, by using different technologies such as sputtering [31,39], electroplating [29] and electroplating plus chromizing [36–38], etc. These metallic coatings work as diffusion barriers, either because they intrinsically have low diffusivities of both Al and some refractory components from the substrates [36–38] or because they can react with some components from the metal substrates or coatings to form new phases (e.g., W-rich precipitates [29], Ni3Hf [39], Ru–Al compound [31]) which resist the interdiffusion between the aluminide coatings and the metal substrates. However, these metallic coatings just retard rather than fully block the interdiffusion between aluminide and metal substrate. They will, accordingly, lose

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resistance. The result also showed that compared with the CeO2-free δ-Ni2Al3 coating, the CeO2-dispersed aluminide coating was profoundly more resistant to the interdiffusion between the aluminide coating and Ni base substrate during 100 h cyclic oxidation in air at 1000 °C. The reason for this, however, remains unclear. We assume that the enhanced ability against the interdiffusion was correlated with the role of the dispersed CeO2, which would act as a diffusion barrier through a mechanism which is not currently understood. To clarify this, the time-dependent evolution of the distribution of the dispersed CeO2, together with its effect on the interdiffusion-correlated degradation of the aluminide in the deliberately-designed coating systems of δ-Ni2Al3/Ni containing or not containing CeO2, has been investigated in the present work. The result will be helpful to develop novel aluminide coatings with an excellent resistance to the aluminide degradation not only by the surface oxidation but also by the interdiffusion with the metal substrates.

Ni-CeO2 film

50 µm Fig. 1. The cross-sectional morphology of the Ni sample with electrodeposition of a ~ 35 μm-thick Ni–CeO2 film.

their effectiveness after long-term operation due to a gradual loss of some components from them [36–38] and an increased instability of the formed phases [29,31,39]. Nonetheless, compared to the metallic coatings, ceramic coatings appear to be better as diffusion barriers, because they generally have unique properties such as exceptional resistance to the diffusion of metal atoms and extra-high thermal stability. Unfortunately, the ceramic diffusion barrier coatings (e.g., TiN [40,41], Al–O–N [27,42], Cr–O–N [42,43], α-Al2O3 [44,45]) are currently used in the system of the metal substrate covered with a M–Cr–Al–Y coating rather than a diffusion aluminide coating. Hence, the development of ceramic diffusion barriers for preventing the aluminide coating degradation by interdiffusion at high temperatures should be another interesting route to be explored. Recently, Peng and co-workers [46] developed a novel ultrafinegrained and CeO2-dispersed δ-Ni2Al3 coating with excellent oxidation

2. Experimental Samples with dimensions of 15 × 10 × 2 mm were cut from pure Ni plates. After being abraded to a final 800 grit SiC paper and then ultrasonically cleaned in acetone, the samples were electrodeposited with a Ni film from a nickel sulfate bath (150 g/l NiSO4 · 6H2O, 120 g/l C6H5Na3O7 · 2H2O, 12 g/l NaCl, 35 g/l H3BO3) loaded with CeO2 particles (commercial products from Alfa Aesar company) with a size of 15–30 nm. The deposited Ni films with a similar CeO2 content but different thicknesses were available through controlling the deposition time. Subsequently, the deposited samples were aluminized at 620 °C using a conventional halide activate pack-cementation in a powder mixture of Al (particles size: ~75 μm) + 55 wt.% Al2O3 (~75 μm) + 5 wt.% NH4Cl in an Ar (purity: 99.99%) atmosphere. The aluminized coating consisted of the aluminide in its δ phase and it grew inward at the temperature as demonstrated in our previous work [46]. The literature also shows that the degradation of the aluminide coating by the interdiffusion with the Ni base substrate was greatly mitigated if the electrodeposited film contained CeO2. Accordingly, two different CeO2-containing aluminide coating systems were specially prepared. One was a similar δ-Ni2Al3/Ni coating system

Al

Ce

30 µm Ni2Al3 Ni-CeO2 film Ni substrate

Ni

O

Fig. 2. The backscattered electron image (BEI) and the corresponding elemental X-ray mappings of cross-section of the Ni2Al3/Ni coating system containing CeO2, by partially aluminizing a ~35 μm-thick Ni–CeO2 film.

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Fig. 3. The Ni–Al binary phase diagram [47].

containing CeO2 as reported in our previous work [46] by partially aluminizing a ~35 μm-thick Ni–CeO2 film, and the other is a δ-Ni2Al3–CeO2/Ni coating system by fully aluminizing a thinner (~10 μm-thick) Ni–CeO2 film for 5 h pack cementation. In addition, another CeO2-free aluminide coating system of δ-Ni2Al3/Ni was also prepared by 5 h aluminizing a Ni film with the thickness close to that of the ~35 μm-thick CeO2-containing film, electrodeposited from a similar bath above but without loading CeO2. Afterwards, the aluminized samples were annealed at 1000 °C in a dynamic vacuum system through the following three subsequent steps. First, the samples were placed into the quartz tube chamber and then the chamber was filled with Ar gas (purity: 99.99%). Second, the chamber was mechanically pumped to a vacuum level of 10−1 Pa. After alternately repeating the two steps several times, the chamber was further evacuated to a value of 1 × 10 −5 Pa by a molecular pump in the last step. The vacuum level remained unchanged during the entire time of annealing. Scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS), electron probe microanalysis (EPMA) and X-ray diffraction (XRD) were used to investigate the structure and phase of the composition of the aluminized coating annealed for different lengths of time. Typical evolution of the structures of the aluminized samples obtained for 10 min and 60 min annealing is reported in the work. It clearly shows the effect of the added CeO2 on the interdiffusion between the aluminide phase of the coating and the underlying Ni, as revealed in detail below.

shows the backscattered electron image (BEI) and the corresponding elemental X-ray mappings of the ~35 μm-thick Ni–CeO2 film after 5 h aluminizing. In the condition, only the outer ~15 μm thickness of the

3. Results The electrodeposited Ni–CeO2 film contains ~3.5 wt.% CeO2, calculated based on the EDS analysis. Fig. 1 shows the cross-sectional morphology of the Ni plate with a ~35 μm-thick Ni–CeO2 film (after etching in 4 vol. % HNO3 + C2H5OH). The black holes correspond to the locations of the co-deposited CeO2 particles, which were exfoliated during etching. The CeO2 particles, because of the agglomeration in the bath, appear to be much larger than the as-received size. The particle clusters are randomly dispersed in the Ni film. During aluminizing at 620 °C, the Ni film can convert into a coating of aluminide in its δ phase as reported in our previous work [46]. Fig. 2

Fig. 4. XRD patterns of the Ni2Al3/Ni coating system containing CeO2 (a) as-aluminized, (b) as-annealed after 10 min at 1000 °C.

X. Tan et al. / Surface & Coatings Technology 224 (2013) 62–70

composite film has been transformed into the δ-Ni2Al3 coating (see the XRD characterization in Fig. 4(a) below), while the inner part of the film remains unaluminized. The aluminide coating is ~43 μm-thick, much thicker than the ~15 μm-thick Ni film that has been aluminized, because a significant molar volume expansion of the conversion of the Ni film into the δ-Ni2Al3. The volume expansion is also the reason that the dispersed CeO2 in the δ-Ni2Al3 is not as dense as that in the underlying Ni film that was not aluminized, as clearly demonstrated by the Ce X-ray mapping. In addition, the as-aluminized coating is very dense. No pores were seen in it. The interdiffusion between the δ-Ni2Al3 coating and the Ni film occurs at high temperatures. By such interdiffusion δ-Ni2Al3 is expected to be degraded into the aluminides with the lower atomic ratio of Al/Ni such as β-NiAl and γ′-Ni3Al, as indicated by the Ni–Al binary phase diagram given in Fig. 3 [47]. The XRD characterization indicates that only after 10 min annealing, the δ phase of the aluminide coating in the probed surface layer has been fully degraded into Ni0.9Al1.1 (JPSD44-1187), an Al-rich β-NiAl phase, as seen in Fig. 4(b). Similar XRD pattern was also acquired after 60 min. The δ phase degradation starts from the Ni2Al3/Ni film interface. Fig. 5 shows the evolution in phase composition of the δ-Ni2Al3 coating with time during annealing at 1000 °C. The EDS line-scan of Al is calibrated with the actual concentrations of Al at the three areas (one close to the surface and the other two close to the interface), based on the six to seven point measurements on various spots but in similar depths). After 10 min annealing, the δ-phase aluminide of the coating, although it was mostly transformed into an Al-rich β-NiAl, formed a thin (the maximum thickness of ~2 μm) Ni-rich β-NiAl band close to the Ni film (Fig. 5(a)). The Ni-rich β-NiAl has a different color contrast from the nonstoichiometric intermetallics rich in Al, as more clearly indicated in the inset, a magnified image of the A-framed zone. According to the morphological characteristics, the degraded aluminide coating can be divided into two zones: a thicker outer zone marked with “I” with the formation of pores and a thinner inner pore-free zone marked with “II”. In addition, the Al-rich β-NiAl of zone I was as thick as the original δ-Ni2Al3 coating. The thickness of the zone I remained almost unchanged after 60 min (Fig. 5(b)), indicating that the outer Al-rich β-NiAl was not varied in thickness with the annealing time. The morphologies of the two zones also seemed not be significantly affected by the annealing time, only the pores in zone I got larger in size and the thickness of zone II was slightly increased by ~3 μm from ~7 μm for 10 min annealing to ~10 μm for 60 min annealing (including the thickness of the inside band of the Ni-rich β-NiAl increased from ~2 μm to ~4 μm). Apparently, the degradation of the β-NiAl was not significant for 60 min annealing with respect to for 10 min. From the EDS line scans in Fig. 5(b), the depth profiles of Al and Ni for 60 min annealing were as sharp as those for 10 min at the degradation front of the aluminide (close to the interface between zone II and the Ni film, which indicates that the interdiffusion across the interface was not substantial). To clarify this, the cross-sectioned samples after annealing for the two different time periods were further investigated using EPMA and the results are shown in Figs. 6 and 7, respectively. From the corresponding Ce X-ray mapping, CeO2 strikingly did not appear in the zone II, but was enriched at the interface of the zone II and the Ni–CeO2 film after 10 min (Fig. 6). In addition, the CeO2 enrichment at the interface became more significant for 60 min (Fig. 7). To understand the role of CeO2 in the aluminide degradation of the coating/film system, the CeO2-free δ-Ni2Al3/Ni coating system was also prepared for comparison by aluminizing ~ 35 μm-thick Ni film in the same condition. XRD characterization indicates that the δ phase in the probed surface zone of the CeO2-free coating system has been transformed to the Al-rich β-NiAl (Ni0.9Al1.1 as presented in Fig. 4). However, heavier degradation was observed at the δ-Ni2Al3/Ni film interface, as revealed in Fig. 8. The annealed aluminide coating also exhibited two zones: zone I with the formation of pores and the pore-free zone II. Zone I was an Al-rich β-NiAl on a basis of the EDS

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analyses and almost remained unchanged in thickness for 60 min annealing (Fig. 8(c)) with respect to for 10 min (Fig. 8(b)). Although these results are similar to the evolution of zone I in the annealed δ-Ni2Al3/Ni coating system containing CeO2 (Fig. 5), the CeO2-free coating system exhibited a much broadened zone II. The zone had gentler slopes of the EDS lines of Al and Ni (particular for 60 min annealing), demonstrating more significant interdiffusion between the aluminide and the Ni film. The increased interface interdiffusion caused zone II to be double-layered (visible from the different color contrast in Fig. 8), forming an outer layer of Ni-rich β-NiAl (42 at.% by EDS measurement) and an inner layer of Ni3Al (26.2 at.% Al). The thickness of the zone was increased from 9 μm for 10 min annealing to ~ 18 μm for 60 min (including the increase of the thickness of the inner Ni3Al layer from ~ 3 μm to ~ 8 μm). The thickness increase was approximately three times as that of the zone II formed in the same time period in the CeO2-dispersed coating system. The aggravated interdiffusion led to the δ phase aluminide to be degraded into γ′ phase at the area close to the interface in the CeO2-free coating system. On

Al-rich -NiAl Ni-rich -NiAl Ni-CeO2 film

(a)

I

II

Ni substrate

A

Ni-rich -NiAl

5 μm

A 55 at.%

Al

53.1at.%

45 at.%

Ni

50 m

(b)

I

54.5 at.% Al

II

51.8 at.%

43 at.%

Ni

50 m Al-rich -NiAl Ni-rich -NiAl Ni-CeO2 film

Ni substrate

Fig. 5. SEM cross-sectional morphologies of the Ni2Al3/Ni coating system containing CeO2, with the EDS line scans of Al and Ni and local quantitative measurements of Al after (a) 10 min and (b) 60 min annealing at 1000 °C. The inset on the upper right is a magnified image of the A-framed zone in (a).

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2 µm

I

Al

Ce

Ni

O

II

Ni rich β-NiAl

30 µm Al rich β-NiAl Ni-CeO2 film Ni substrate

Fig. 6. BEI and the corresponding elemental X-ray mappings of cross-section of the Ni2Al3/Ni coating system containing CeO2 after 10 min vacuum annealing at 1000 °C.

the basis of the comparison result presented here, one can assume that the formation of a CeO2-rich layer at the aluminide/Ni interface would be the reason for the mitigated degradation of the aluminide coating during annealing. This is confirmed and interpreted below.

4. Discussion There are two important observations for the annealing of the δ-Ni2Al3/Ni coating system containing CeO2. First, the oxide accumulated

Fig. 7. BEI and the corresponding elemental X-ray mappings of cross-section of the Ni2Al3/Ni coating system containing CeO2 after 60 min vacuum annealing at 1000 °C.

X. Tan et al. / Surface & Coatings Technology 224 (2013) 62–70

(a)

50 m Ni2Al3

Ni film

Ni substrate

(b)

I

II

54.7 at.%Al 52.1 at.%

42 at.%

Ni

26.2 at.%

50 m Al rich -NiAl Ni rich -NiAl Ni3Al

Ni

(c)

I

II

53.5 at.%Al 51 at.%

Ni

41 at.% 26 at.%

at the degradation front of the aluminide. Second, the formed CeO2-rich layer appeared as a diffusion barrier suppressing the degradation of β-NiAl transformed from the δ-Ni2Al3. Based on the characterization of the phase compositions of the aluminides transformed during the annealing, the process for the formation of the CeO2-rich diffusion barrier, together with the structural evolution of the coating, can be schematically presented using a model in Fig. 9. The process is divided into three stages as demonstrated below. Stage I — Formation of a thin interdiffusion zone at the interface. The interdiffusion between the δ-Ni2Al3 coating (zone I) and the Ni film occurs at the onset of annealing. It causes the coating of the δ phase, which has a very narrow domain at 1000 °C in the Ni–Al binary phase diagram (Fig. 3), to be easily transformed into β phase (zone II), starting from the Ni2Al3/Ni interface. This accordingly forms two new interfaces of δ-Ni2Al3/β-NiAl (indicated by S1) and β-NiAl/γ-Ni (indicated by S2). From Fig. 5, the β-NiAl phase is richer in Al close to the S1 interface and richer in Ni close to S2 interface. The result is also consistent with the prediction by the Ni–Al binary phase diagram. Stage II — Formation of a CeO2-rich layer with the growth of the interdiffusion zone. The δ-β phase transformation in zone I, which is quickly and has been finished before 10 min (Fig. 4), is a result of the interdiffusion of Al and Ni across the interface S1. S1 with respect to the original surface of the as-aluminized sample is almost an immobile interface, because of an insignificant molar volume change of the δ-β phase transformation (only ~2% if the β phase is Ni0.9Al1.1 as identified by XRD above). Much higher diffusion rate of Al relative to the counterdiffusion rate of Ni across the Al-rich β-NiAl phase [48,49] at the interface S1 leads to the diffusion of Kirkendall vacancies into the zone I, condensing and then forming voids at high energy areas such as grain boundary triple junctions and the interface between CeO2 and aluminide. The continuous condensation of the vacancies leads to an increase of the voids in size with time (Fig. 5). Besides, the different thermal expansion coefficient of the ceramic particles relative to the metallic matrix may also contribute to the formation of the voids around the particles. In the same time, the interdiffusion at S2 is dominated by the outward diffusion of Ni from Ni–CeO2 film to the zone II, due to a relatively higher diffusion rate of the element in the Ni-rich side of the new β-NiAl layer to Al in the Ni film [50,51]. In the case, the Kirkendall vacancies diffuse with Al toward the Ni film. Thus, no pores occur in zone II. The interdiffusion of Al and Ni across S2, together with its induced about 200% molar volume expansion of the γ(Ni) to β transformation, causes the expansion of zone II through inward movement of S2. Thus, the interface will sweep the CeO2 particles in the Ni film. Afterward, the particles can then be dragged to move with the interface (this is the reason that no CeO2 was seen in the “II” zone as shown in Figs. 6 and 7) and then gradually enriched at S2, forming the CeO2-rich layer. It has been reported [52] that in the case of the particle movement by volume diffusion (Dv) of matrix atoms in a single phase matrix, the combined movement of a grain boundary–particle complex highly depends on the attraction force (F) between the boundary and the particle, the mobility of the particle (a function in reverse proportion to the radius (r) of the particle assumed spherical in shape and in proportion to Dv). The velocity (v) of the boundary–particle complex movement in the Ni film can be expressed by [53] v¼

50 m Al-rich -NiAl Ni-rich -NiAl Ni3Al

Ni

Fig. 8. SEM cross-sectional morphologies of the Ni2Al3/Ni coating system not containing CeO2, after annealing for (a) 0 min, (b) 10 min and (c) 60 min in vacuum at 1000 °C, with the indication of the EDS scanning lines and local quantitative measurements of Al in (b) and (c).

67

ΩDv F πr 3 kT

ð1Þ

where k is the Boltzmann’s constant, T is the absolute temperature, Ω is the Ni atomic volume of the film. If the grain boundary–particle complex is replaced using the interface–particle complex and accordingly the Dv is substituted by D, the relative diffusion coefficient of Ni to Al in the present case, we would interpret the particle movement together with the interface to form an enrichment layer. If there is an appropriate level of F between the CeO2 particles and the film Ni matrix, such a high level of D [50,51] and a small value of r (in nanosize regime) may offer a

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S1

S2

JNi JAl

S2

Al-rich

Ni-rich

-NiAl

-NiAl

Ni-CeO2 film Substrate

Ni2Al3

Stage I

S1

void CeO2

JNi

S2

JAl JVacancy J Ni JAl

JAl

JVacancy

JNi Ni-rich -NiAl Ni2Al3

Ni-CeO2 film

Al rich -NiAl

Substrate

S2

S1

Ni-CeO2 film Substrate

Stage II

Ni-rich -NiAl Al rich -NiAl

Ni-CeO2 film Substrate

Stage III Fig. 9. Model schematically showing the formation of a CeO2-rich diffusion barrier at the aluminide degradation front in the Ni2Al3/Ni coating system containing CeO2 in three stages: stage Ι in which a thin interdiffusion zone (zone II) forms between Ni2Al3 coating and Ni–CeO2 film, stage II in which a CeO2-rich layer forms with the growth of zone II toward the film, and stage III in which the CeO2-rich layer acts as an interdiffusion barrier.

driving force high enough to cause the oxide particles to be dragged by the moving S2. Stage III — Formation of a CeO2-rich diffusion barrier layer. The movement of the interface gets difficult, with the increase of the number of the CeO2 attached to S2 (the reason can be attributed to the increase in the value of r from Eq. (1)). Accordingly, the interface movement gradually slows down and finally stops. In the case, the CeO2-rich layer fully acts as a barrier for the interface interdiffusion. In contrast, there is no force for inhibiting the movement of the S2 interface in the δ-Ni2Al3/Ni coating system without CeO2 during annealing. With the incessant movement of S2, the aluminide more severely degrades. This explains the occurrence of a thicker zone II in the CeO2-free coating system with an inner layer of γ′-Ni3Al, which was formed from the further degradation of the β-NiAl (Fig. 8). From the above model, the interdiffusion between the β-NiAl phase and Ni–CeO2 film is proposed as the prerequisite for the formation of the CeO2-rich layer. In other words, the layer would not be seen in

the absence of the interdiffusion in case where the electrodeposited film has been entirely aluminized. Moreover, the CeO2-rich layer, once after its formation, is assumed to act as a diffusion barrier for the degradation of the β-NiAl. To substantiate the two assumptions, we prepared another δ-Ni2Al3 coating in the same aluminizing condition, whose growth was completely through the thickness of the Ni–CeO2 film (~10 μm-thick film used here as demonstrated in the Experimental section). The EPMA result of the aluminized sample is presented in Fig. 10. The outer CeO2-dispersed δ-Ni2Al3 layer (~25 μm-thick) was formed on the Ni–CeO2 film and the inner CeO2-free δ-Ni2Al3 layer (~23 μm-thick) on the pure Ni substrate. In the case, the interdiffusion between aluminide and Ni (rather than the CeO2-containing Ni) dominates the aluminide degradation. Fig. 11 shows the BEI and corresponding elemental X-ray maps of the aluminide coating after 60 min annealing. Clearly, no CeO2-rich layer occurred. The degradation front zone (i.e., zone II as indicated above) was ~17 μm-thick and double-layered. The outer was the Ni-rich β-NiAl layer and the inner was ~6 μm-thick

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Al

Ce

Ni

O

CeO2-free Ni2Al3

30 µm CeO2-dispersed Ni2Al3

Ni substrate

Fig. 10. BEI and the corresponding elemental X-ray mappings of cross-section of the coating system of Ni2Al3–CeO2/Ni.

γ′-Ni3Al layer. The morphological characteristics are very similar to that of the zone II in the annealed CeO2-free coating system of δ-Ni2Al3/Ni (Fig. 8). The result strongly confirms the above interpretation on the formation of the CeO2-rich layer and its role in blocking the aluminide degradation interdiffusion.

5. Conclusions 1) By partially aluminizing an electrodeposited Ni–CeO2 film using a conventional pack cementation method, a coating system of δ-Ni2Al3/Ni containing CeO2 was developed. It itself formed a layer of a CeO2-rich

Al

I

Ni rich β-NiAl

Ce

II

Ni3Al

30 µm Al rich β-NiAl

Ni substrate

Ni

O

Fig. 11. BEI and the corresponding elemental X-ray mappings of cross-section of the coating system of Ni2Al3–CeO2/Ni after 60 min annealing at 1000 °C.

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barrier to the aluminide degradation by the interdiffusion of Al and Ni during annealing at 1000 °C. 2) The CeO2-rich barrier layer was quickly formed and its formation is intrinsically correlated with the interdiffusion-induced interface movement of the aluminide/Ni toward the Ni film. The CeO2 particles in the film were swept and then dragged by the moving interface. Accordingly, the oxides accumulation at the interface finally retarded the interface movement by the interdiffusion.

[25] [26] [27] [28] [29]

Acknowledgement

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This research is sponsored by National Natural Science Foundation of China (NSFC, project Grant No. 51071162). 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]

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