Corrosion behavior of Ni-based structural materials for electrolytic reduction in lithium molten salt

Journal of Nuclear Materials 412 (2011) 157–164

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Corrosion behavior of Ni-based structural materials for electrolytic reduction in lithium molten salt Soo Haeng Cho a,⇑, Sung Bin Park a, Jong Hyeon Lee b,⇑, Jin Mok Hur a, Han Soo Lee a a b

Korea Atomic Energy Research Institute, 1045 Daedeokdaero Yuseong-gu, Daejeon 305-353, Republic of Korea Graduate School of Green Energy Technology, Chungnam National University, 79 Daehak-ro, Yuseong-gu, Daejeon 305-764, Republic of Korea

a r t i c l e

i n f o

Article history: Received 21 January 2011 Accepted 17 February 2011 Available online 23 February 2011

a b s t r a c t In this study, the corrosion behavior of new Ni-based structural materials was studied for electrolytic reduction after exposure to LiCl–Li2O molten salt at 650 °C for 24–216 h under an oxidizing atmosphere. The new alloys with Ni, Cr, Al, Si, and Nb as the major components were melted at 1700 °C under an inert atmosphere. The melt was poured into a preheated metallic mold to prepare an as-cast alloy. The corrosion products and fine structures of the corroded specimens were characterized by scanning electron microscope (SEM), Energy Dispersive X-ray Spectroscope (EDS), and X-ray diffraction (XRD). The corrosion products of as cast and heat treated low Si/high Ti alloys were Cr2O3, NiCr2O4, Ni, NiO, and (Al,Nb,Ti)O2; those of as cast and heat treated high Si/low Ti alloys were Cr2O3, NiCr2O4, Ni, and NiO. The corrosion layers of as cast and heat treated low Si/high Ti alloys were continuous and dense. However, those of as cast and heat treated high Si/low Ti alloys were discontinuous and cracked. Heat treated low Si/ high Ti alloy showed the highest corrosion resistance among the examined alloys. The superior corrosion resistance of the heat treated low Si/high Ti alloy was attributed to the addition of an appropriate amount of Si, and the metallurgical evaluations were performed systematically. Ó 2011 Elsevier B.V. All rights reserved.

1. Introduction In recent research and development studies, processes such as electrorefining, lithium reduction, and pyrochemical separation have gained more popularity as alternatives for the treatment of spent oxide nuclear fuels than hydrometallurgical processes. Argonne National Laboratory (ANL) has been developing pyrochemical technologies based on a molten-salt system [1,2]. In addition, the Advanced Spent Fuel Conditioning Process (ACP) has been under development at the Korea Atomic Energy Research Institute (KAERI) since 1997. The process aims to convert the spent oxide fuel into a metallic form by electrolytic reduction [3,4]. The electrolytic reduction process for the spent oxide nuclear fuel is carried out in LiCl–Li2O molten salt at 650 °C. The liberation of oxygen at the anode and the high-temperature molten salts used in the electrolytic reduction result in a chemically aggressive environment that is too corrosive for various components of the major electrolysis equipment [1]. Hence, studies on the corrosion of nickel-based superalloys used to fabricate the equipment employed for handling high-temperature molten salts have also been conducted [5–13]. Based on an earlier KAERI study on molten salt corrosion of commercially available nickel-based superalloys [14], Inconel

⇑ Corresponding authors. Tel.: +82 42 868 2584; fax: +82 42 868 2024 (S.H. Cho), tel.: +82 42 821 6596; fax: +82 42 822 5850 (J.H. Lee). E-mail addresses: [email protected] (S.H. Cho), [email protected] (J.H. Lee). 0022-3115/$ - see front matter Ó 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2011.02.047

713LC was one of the superalloys that exhibited exceptional corrosion resistance. However, the corrosion resistance of Inconel 713LC did not meet the requirement of long-term operation. Hence, we have been trying to improve the corrosion resistance of structural materials for the electrolytic reduction processes further, and we developed two new customized superalloys, N101 and N102, on the basis of a modified composition of Inconel 713LC. In this study, the hot corrosion behaviors of N101, N101H (N101 with heat treatment), N102, and N102H (N102 with heat treatment) in the presence of lithium molten salt were investigated under simulated electrolytic reduction conditions. 2. Experimental 2.1. Specimen preparation The new alloy system comprises Ni, Cr, Nb, Si, Ti, and Al as the major components. It has a composition similar to that of the commercial Inconel 713LC, but the new alloy has reinforced Si in its matrix and a reduced Mo content. A batch of 50 kg of alloy was melted at 1700 °C under an inert atmosphere. The melt was poured into a preheated metallic mold to prepare an as-cast alloy. Among the alloying elements, the iron and molybdenum contents were decreased to the minimum level compared to Inconel 713LC, because they were found to adversely affect the corrosive behavior in a previous study [14]. The composition of the alloy formed is

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Table 1 Chemical compositions of tested alloys (wt%).

b) N101H

Alloy

Ni

Cr

Fe

Co

C

Al

Ti

Nb

Si

N101 N102

Bal Bal

12.0 12.5

0.11 0.15

0.06 0.06

0.06 0.04

5.80 6.10

1.05 0.53

1.96 2.00

2.04 4.67

P2 P1

c) N102

d) N102H

P1

Ar gas outlet

Ar gas inlet

P2

P1

Flow meter

P2 P3

MixedGas (Ar-O2 )

50µm

Container Thermocouple Protective crucible Specimen Test crucible

Mixed-gas supply system

Elements

Al

Si

Nb

Ti

Cr

Ni

P1

5.4

5.2

4.9

1.5

9.9

Bal

P2

7.0

1.9

2.2

1.2

11.5

Bal

P1

5.6

2.3

3.3

1.5

9.7

Bal

Alloys Baffle plate

N101 a)

Specimen hanger Mixed-gas nozzle Molten salt

Heating furnace

Spectrum

Cooling jacket

N101H b)

N102 c)

Fig. 1. Schematic diagram of apparatus for corrosion test.

N102H d)

almost in agreement with the design goal, which aims to observe the effect of silicon on corrosion resistance. The chemical compositions of the alloys used are listed in Table 1. Specimens of size 70  15  2 mm3 were ground along with SiC, polished using a diamond paste, and cleaned in acetone. 2.2. Hot corrosion testing The experimental apparatus used is shown in Fig. 1. The LiCl– Li2O molten salt was introduced into a high-density MgO crucible and then heated at 300 °C for 3 h in an argon atmosphere to remove any possible moisture pickup. After reaching the set conditions, the specimens and alumina tube were immersed in the molten salt, and mixed gas (Ar–10% O2) was supplied through an alumina tube. The corrosion tests were carried out at 650 °C and for 24–216 h. The Li2O concentration in LiCl was 3 wt%. Following the corrosion test, the specimens were withdrawn from the salt and kept in argon gas while the furnace was cooled to room temperature. 2.3. Characterization The vessel was opened, and the specimens were removed for visual examination before cleaning. The specimens were then cleaned ultrasonically in deionized water, measured, and weighed. The corrosion products and fine structures of the corroded samples were analyzed by XRD (X-ray diffraction, Rigaku, DMAX/1200), SEM (Scanning Electron Microscope, JEOL, JSM-6300), and EDS (Energy Dispersive X-ray Spectroscope, JEOL, JSM-6300). 3. Results and discussion 3.1. Microstructure of the as-cast alloys Fig. 2 shows the microstructures and EDS analysis results of the N101 and N102 alloys before and after the heat treatment. The ascast alloys predominantly have dendritic microstructures that

P2

5.4

2.1

3.2

1.3

11.5

Bal

P1

10.9

5.3

0.8

0.6

8.9

Bal

P2

7.9

4.2

2.4

0.6

9.9

Bal

P1

4.6

4.9

1.4

0.4

19.2

Bal

P2

3.3

10.8

16.0

1.6

5.8

Bal

P3

8.6

4.8

3.1

0.9

6.1

Bal

Fig. 2. Cross-sectional microstructures and EDS analysis results of as-cast N101 and heat-treated N101H.

contain segregated solutes between the dendritic arms. Hence, it is believed that solutes having a relatively lower solubility were ejected out of the solidifying dendrite, resulting in microsegregation. The EDS analysis result clearly indicated that the interdendritic segregation comprises Si- and Nb-rich Ni phases, as shown in P1 (Fig. 2a). Among the alloying elements, Si and Nb have a relatively lower solubility in a Ni matrix (c-phase), so excess quantities of these elements were rejected by the solidified dendrite, and they were deposited into the interdendritic channels. Hence, the concentration of Cr, which has the highest solubility in the Ni matrix among the alloying elements, is relatively lower in the segregation area. In other words, Cr is preferentially dissolved in the Ni matrix than ejected out of the matrix. In the case of N101 heat-treated at 1100 °C for 3 h and then quenched, the dendritic microstructure and microsegregation were completely absent, but the second phase (P1) having a size of 20 lm was present, as shown in Fig. 2b. The EDS analysis result of the matrix P2 in Fig. 2b shows that the alloying elements are relatively evenly distributed, while P1 has slightly higher Nb and lower Cr concentrations than the c-phase. The result suggests that the heat treatment at 1100 °C for 3 h solutionized the solute into the c-phase without severe microsegregation. In the case of the N102 alloy, a completely different microstructure having three distinguished regions was observed after heat treatment, as shown in Fig. 2d. P1 believed to be the c-phase has a very high Cr concentration, whereas other analyzed alloying elements were found to have relatively low concentrations. Contrary to the elemental distribution of P1, P2 has very high Si and Nb concentrations, and all elements except for Cr are more or less evenly distributed at P3, where Cr has a low concentration. It should be noted that a slight increase in Si in the N102 alloy caused a drastic change in the microstructure as compared to that of N101. A reasonable explanation for this phenomenon can be the formation of intermetallic compounds, such

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24

N101 N101H N102 N102H Inconel 713LC ( f (reference) )

We eightt los ss mg/cm

2

20 16 12 8 4 0

0

50

100

150

200

250

Time (h) Fig. 3. Weight losses of alloys corroded at 650 °C as a function of time.

as Cr3Si, Nb5Si3, and NbSi2 caused by excess free Si, which was removed from the c-phase. The heterogeneity of the microstructure is vulnerable to all sorts of corrosion, especially pitting, even after the heat treatment. Therefore, the corrosion rate of N102 is expected to be higher than that of N101 under the electrolytic reduction environment. 3.2. Corrosion rate The weight loss of the specimens after the corrosion tests in the LiCl–Li2O molten salt as a function of time is shown in Fig. 3. The

(a)

(c)

80μ μm

80μ m

corrosion rates were in the following order: N101H < N101 < N102 < N102H. N101H showed the highest corrosion resistance among the examined alloys. Fig. 4 shows representative crosssectional microstructures of N101 and N101H after the corrosion test for 72 h and 168 h. It is clearly seen that a dense oxide layer is formed on the surface of the as cast N101 specimen with 80 lm of thickness after 72 h of corrosion test (Fig. 4a). The oxide layer, however, is detached from the specimen after 168 h of corrosion test (Fig. 4b). The thickness of the oxide layer of the heat treated N101H specimen is below 55 lm after the corrosion test for 72 h and the thickness increased to 65 lm after the corrosion test for 168 h as shown in Fig. 4c and d. It should be noted that intact interface is found in the heat treated N101H specimen even after 168 h of corrosion test. Hence, this microstructural observation indirectly confirms the enhancement of interfacial strength between the oxide layer and the base metal after the heat treatment of the specimen. The weight loss of the specimens with time was attributed to the cracking of the protective layer, leading to spallation from the base metal surface as well as dissolution of the scales and other corroded products [15]. Therefore, it was thought that the formation of the corrosion layer and its adherence to the base metal significantly influenced the weight variation of the specimens. 3.3. Corrosion products Fig. 5 shows XRD patterns of the corrosion products after the corrosion tests in a LiCl–3% Li2O molten salt at 650 °C for 72– 216 h. For superalloys N101 and N101H corroded for 72 h, peaks attributed to the corrosion products Cr2O3, NiCr2O4, and NiO were observed (Fig. 5a and b). NiCr2O4 (spinel-type) is believed to be

(b)

(d)

80μm

80μm

Fig. 4. Cross-sectional microstructures of N101 ((a) 72 h, (b) 168 h) and N101H ((c) 72 h, (d) 168 h).

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: NiO : Cr C 2O3

(a)

: Ni : NiCr2O4 : (Al,Nb,Ti)O (Al Nb Ti)O2

: NiO : Cr C 2O3

(b)

: Ni : NiCr2O4 : (Al,Nb,Ti)O (Al Nb Ti)O2 216 h

Intens sity

Intens sity

216 h

72 h

20

30

40

50

60

70

80

90

72 h

100

20

30

40

50

60

2θ

70

80

90

100

2θ : NiO : Cr2O3

: Ni

: NiCr2O4

: NiO : Cr2O3

(d)

216 h

: Ni

: NiCr2O4 216 h

In ntensity

In ntensity

(c)

72 h

20

30

40

50

60

70

80

90

72 h

20

100

30

40

50

60

70

80

90

100

2θ

2θ

Fig. 5. XRD patterns of corrosion products of (a) N101, (b) N101H, (c) N102, and (d) N102H corroded at 650 °C for 72–216 h.

90

E Eleme ent portion (wt.%) ( )

80

substrate

corrosion layer

70 60

O Al Si Ti Cr Ni Nb

50 40 30 20 10

30μm

0

S1

S2

S3

S4

S5

S6

Point (No.) Fig. 6. Cross-sectional microstructures and EDS analysis results of N101 corroded at 650 °C for 168 h.

formed by a solid-state reaction with the corresponding oxide (NiO and Cr2O3) [16]. As the corrosion progressed, (Al,Nb,Ti)O2 was formed; the formation was attributed to the external diffusion of oxygen active elements, such as Al, Nb, and Ti [17]. The corrosion products of superalloys N102 and N102H were Cr2O3, NiO, and NiCr2O4 for 72 and 216 h (Fig. 5c and d). From the thermodynamic point of view, Cr2O3 is the most stable oxide in Ni–Cr-based alloys [18]; therefore, it was formed initially as a corrosion product.

3.4. Corrosion behavior Fig. 6 shows the cross-sectional microstructure and EDS analysis results of N101 after the corrosion test at 650 °C for 168 h in the LiCl–3% Li2O molten salt. A distinct dense oxide layer was formed at the surface of the specimen, and the concentration of Al and Si increased at the interface between the oxide layer and the metallic part marked as S5 in Fig. 6. However, an element depletion zone was observed beneath the oxide layer (S6 in Fig. 6). This is because

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90

Element portio on (wtt.%)

substrate

corrosion layer

80 70 60

O Al Si Ti Cr Ni Nb

50 40 30 20 10

30μm

0

S1

S2

S3

S4

S5

Point (No.) Fig. 7. Cross-sectional microstructures and EDS analysis results of N101H corroded at 650 °C for 168 h.

90

Element portiion (w wt.%)

substrate

corrosion layer

80 70 60

O Al Si Ti Cr Ni Nb

50 40 30 20 10

200 200μm

0

S1

S2

S3

S4

S5

S6

S7

Point (No.) Fig. 8. Cross-sectional microstructures and EDS analysis results of N102 corroded at 650 °C for 168 h.

the elements having a higher oxygen affinity migrated to the surface during the formation of the oxide layer at the alloy surface under the corrosion environment, causing local concentration differences in the alloying elements. Fig. 7 shows the cross-sectional microstructure and EDS analysis results of the heat-treated N101H specimen after the corrosion test at 650 °C for 168 h in the LiCl–3% Li2O molten salt. The concentration deviation of the alloying elements was significantly mitigated. However, slight heterogeneity was retained. The microsegregation in the casting structure formed during solidification was eliminated by solutionizing heat treatment, enhancing the homogeneity of the alloying elements not only in the metallic matrix but also in the interface layer. Fig. 8 shows the cross-sectional microstructure and EDS analysis results of the N102 specimen after the corrosion test at 650 °C for 168 h in the LiCl–3% Li2O molten salt. An increase of 4.67% in the Si content seemed to be excessive because the segregation of the alloying elements was prevalent throughout the matrix. The segregation also affected the oxide layer; this resulted in a microstructure more irregular than that of the N101 specimen. In the Ni–Si binary system, the solubility of Si in Ni is about 5 wt% at 700 °C, but the actual solubility is expected to decrease with the addition of another alloying element. Hence, severely segregated phases are observed in N102, as shown in Fig. 9, even after the heat treatment, owing to the decreased solubility of the alloying

elements. It is believed that the segregated phases accelerated the local corrosion under the corrosion environment. This resulted in an irregular oxide layer as well as an irregular interface between the oxide layer and the metallic matrix, as shown in Fig. 9. As shown in Figs. 6 and 7, all the corrosion layers are intact without any spallation. In particular, N101 and N101H have a continuous, dense external corrosion layer, while N102 and N102H have occasional cracks (Figs. 8 and 9). In addition, the adherence of the corrosion layer of N101H was higher than that of N101. This result is also closely related to weight loss, as shown in Fig. 3. Fig. 10 shows the cross-sectional microstructure and EDS mapping results of the N101 specimen after the corrosion test at 650 °C for 72 h in the LiCl–3% Li2O molten salt. The oxide layer is continuous and dense, containing mainly Cr and Ni oxides, as shown in Fig. 10. The external oxide layer is thought to be of the NiCr2O4 spinel type, formed by the solid-state reaction between Cr2O3 and NiO [16]. An aluminum oxide layer was formed beneath the outer oxide layer, and it is believed that the oxygen active element of Al was preferentially oxidized among the alloying elements. It was reported that addition of approximately 4–5 wt% Al in the Ni-based alloy contributed to the high temperature oxidation resistance due to the formation of Al2O3 [19]. Fig. 11 shows the cross-sectional microstructure and EDS mapping results of the heat-treated N101H specimen after the

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90

Element porttion (wt.%)

substrate b t t

corrosion layer

80 70

O Al Si Ti C Cr Ni Nb

60 50 40 30 20 10 0

μ 200μm

S1

S2

S3

S4

S5

Point (No.) Fig. 9. Cross-sectional microstructures and EDS analysis results of N102H corroded at 650 °C for 168 h.

O

Al

Si

Cr

Nb

Ti

μm 80μ Ni

Fig. 10. Cross-sectional microstructures and maps of Cr, Ni, O, Si, Nb, Ti, and Al for N101 corroded at 650 °C for 72 h.

O

Al

Si

Cr

Nb

Ti

80μ 80 μm Ni

Fig. 11. Cross-sectional microstructures and maps of Cr, Ni, O, Si, Nb, Ti, and Al for N101H corroded at 650 °C for 72 h.

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O

Al

Si

Cr

Nb

Ti

80μ μm Ni

Fig. 12. Cross-sectional microstructures and maps of Cr, Ni, O, Si, Nb, Ti, and Al for N102 corroded at 650 °C for 72 h.

O

Al

Si

Cr

Nb

Ti

80μ 80 μm Ni

Fig. 13. Cross-sectional microstructures and maps of Cr, Ni, O, Si, Nb, Ti, and Al for N102H corroded at 650 °C for 72 h.

corrosion test at 650 °C for 72 h in the LiCl–3% Li2O molten salt. Like in the as-cast N101 specimen, the outer corrosion layer is continuous and dense, containing mainly Cr and Ni oxide, which are probably present in the form of a NiCr2O4 spinel. A Ni-rich area was formed on the outer oxide layer, and the Ni-rich phase retarded the internal diffusion of oxygen ions [20]. Hence, the corrosion resistance of the specimens is attributed to Ni itself in the high-temperature oxidative molten salt environment. In addition, an aluminum oxide layer like in the case of the N101 specimen was formed beneath the outer oxide layer. The dense Ni-rich layer responsible for the retarded internal diffusion of oxygen ions is closely related to the corrosion rate measurements, as shown in Fig. 3. Fig. 12 shows the cross-sectional microstructure and EDS mapping results of the as-cast N102 specimen after the corrosion test at 650 °C for 72 h in the LiCl–3% Li2O molten salt. The morphology of the outer corrosion layer of the as-cast N102 specimen was completely different, representing an intermittent and irregular structure containing Cr and Ni oxide present in the form of a NiCr2O4 spinel. The contribution of the Al oxide layer to the adhesive strength of the protective oxide layer seems to be negligible in spite of the formation of the former beneath the outer corrosion

surface. Because the intermittent and irregular outer corrosion layer is porous, the internal diffusion of oxygen ions cannot be effectively prohibited, accelerating the spallation of the corrosion layer. This microstructural evaluation result agrees well with the corrosion rate measurements as shown in Fig. 3. Fig. 13 shows the cross-sectional microstructure and EDS mapping results of the heat-treated N102H specimen after the corrosion test at 650 °C for 72 h in the LiCl–3% Li2O molten salt. The outer oxide layer of the heat-treated N102H exhibits even more intermittent and irregular cracks than the as-cast N102. In addition, the outer oxide layer showed partial spallation. The outer oxide layer is composed of Cr and Ni oxides, and a NiCr2O4 spinel too is believed to be formed by the abovementioned solid-state reaction between Cr2O3 and NiO. The porous oxide corrosion layer of the heat-treated N102H specimen accelerates the internal diffusion of oxygen. Hence, the corrosion rate of this specimen is expected to be higher than that of any other specimens owing to spallation of the corrosion layer. The decrease in the solubility caused by adding excess Si is responsible for these irregularities, and the microstructural observations agree closely with the corrosion rate measurements as shown in Fig. 3.

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4. Conclusion

References

The corrosion rates of the customized Ni-based superalloys in the high-temperature lithium chloride molten salt are in the order of N101H < N101 < N102 < N102H. The heat-treated N101H specimen exhibited superior corrosion resistance owing to the formation of the most continuous, dense, and adherent protective oxide layer. The corrosion product of the as-cast N101 and heat-treated N101H is composed of Cr2O3, NiCr2O4, Ni, NiO, and (Al,Nb,Ti)O2. The main constituents of the corrosion product of the as-cast N102 and heat-treated N102H are Cr2O3, NiO, Ni, and NiCr2O4. The superior corrosion resistance of N101H is attributed to the formation of the oxide layer and its sustenance owing to a higher adhesive strength between the metallic matrix and the protective oxide layer than any other specimen. The preferential oxidation of the oxygen active elements of Al, Ti, and Nb that contributed to the retardation of internal diffusion of oxygen with the Ni-rich layer and the addition of Si should not exceed 2 wt% owing to the solubility limit.

[1] J.E. Indacochea, J.L. Smith, K.R. Litko, E.J. Karell, A.G. Raraz, Oxid. Met. 55 (2001) 1. [2] E.J. Karell, K.V. Gourishankar, J.L. Smith, L.S. Chow, L. Redey, Nucl. Technol. 136 (2001) 342. [3] S.B. Park, B.H. Park, S.M. Jeong, J.M. Hur, C.S. Seo, S.H. Choi, S.W. Park, J. Radioanal. Nucl. Chem. 268 (2006) 489. [4] S.M. Jeong, H.S. Shin, S.H. Cho, J.M. Hur, H.S. Lee, Electrochim. Acta 54 (2009) 6335. [5] M.A. Uusitalo, P.M.J. Vuoristo, T.A. Mantyla, Corros. Sci. 46 (2004) 1311. [6] J.G. Gonzalez, S. Haro, A. Martinez-Villafane, V.M. Salinas-Bravo, J. PorcayoCalderon, Mat. Sci. Eng. A 435–436 (2006) 258. [7] B.P. Mohanty, D.A. Shores, Corros. Sci. 46 (2004) 2893. [8] A. Ruh, M. Spiegel, Corros. Sci. 48 (2006) 679. [9] Tz. Tzvetkoff, J. Kolchakov, Mater. Chem. Phys. 87 (2004) 201. [10] S. Mitsushima, N. Kamiya, K.I. Ota, J. Electrochem. Soc. 137 (1990) 2713. [11] R.A. Rapp, Corros. Sci. 44 (2006) 209. [12] B. Zhu, G. Lindbergh, Electrochim. Acta 46 (2001) 2593. [13] M. Spiegel, P. Biedenkipf, H.J. Grabke, Corros. Sci. 39 (1997) 1193. [14] S.H. Cho, J.M. Hur, C.S. Seo, S.W. Park, J. Alloys Compd. 452 (2008) 11. [15] M.H. Guo, Q.M. Wang, P.L. Ke, J. Gong, C. Sun, R.F. Huang, L.S. Wen, Surf. Coat. Technol. 200 (2006) 3942. [16] G.C. Wood, F.H. Stott, Mater. Sci. Technol. 3 (1987) 519. [17] E.T. Turkdogan, Physical Chemistry of High Temperature Technology, Academic Press, New York, 1980. [18] H. Izuta, Y. Komura, J. Jpn. Inst. Met. 58 (1994) 1196. [19] Y. Harada, Jpn. Therm. Spray. Soc. 33 (1996) 128. [20] G.C. Wood, B. Chattopadhyay, Corros. Sci. 10 (1970) 471.

Acknowledgment This work was funded by the National Mid- and Long-term Atomic Energy R&D Program supported by the Ministry of Education, Science and Technology of Korea.