Hot Corrosion Behavior of a Cr-Modified Aluminide Coatingon a Ni-Based Superalloy

Duoli Wu • Sumeng Jiang • Qixiang Fan • Jun Gong • Chao Sun

Received: 21 November 2013 / Revised: 9 January 2014 / Published online: 15 July 2014

� The Chinese Society for Metals and Springer-Verlag Berlin Heidelberg 2014

Abstract A Cr-modified aluminide coating is prepared on a Ni-based superalloy using arc ion plating and subsequent

pack cementation aluminizing. Hot corrosion behavior of the Cr-modified aluminide coating exposed to molten Na2SO4/

K2SO4 (3:1) or Na2SO4/NaCl (3:1) salts at 900 �C in static air are evaluated as well as the aluminide coating. The results

indicate that compared with the aluminide coating, the anti-corrosion properties of the Cr-modified aluminide coating in

the both salts are improved, which should be attributed to the beneficial effect of the Cr in the coating. The corrosion

mechanism of the Cr-modified aluminide coating, especially the role of Cr in the mixture salt corrosion, is discussed.

KEY WORDS: Arc ion plating; Pack cementation aluminizing; Superalloy; Cr-modified aluminide coating;

Hot corrosion

1 Introduction

Currently, Ni-based superalloys are widely applied to tur-

bine blades or other components of gas turbines for their

excellent high-temperature mechanical properties. These

components must exhibit a high level of resistance to the

oxidation and corrosion conditions generated by the com-

bustion environment and be resistant to any associated

erosion/corrosion conditions produced as a result of par-

ticulate ingestion or solids formed by incomplete com-

bustion [1]. The alloying requirements for these high

volume fraction gamma prime materials result in a reduc-

tion in corrosion resistance and therefore, surface coatings

are widely used. The general design philosophy is to select

a high-strength substrate alloy to withstand the stress and

apply a surface coating to give maximum protection

against the environment corrosion.

Ni-based superalloys turbine blades are commonly alu-

minized in a cementation pack for improving their oxida-

tion resistance. However, such aluminized coatings lack

adequate resistance to hot corrosion caused by deposits of

fused alkali sulfates. The corrosion resistance of aluminide

coatings intensively depends on their ability to form a

continuous, adherent and slow growing layer of a-Al2O3

[2]. Simple aluminide coating made of b-NiAl is brittle and

sensitive to sulfur which tends to segregate at grain

boundaries, weakening the oxide–metal interface, and thus

leading to exfoliation [3].

A Cr-modified aluminide coating would promise sub-

stantial improvements in the hot corrosion resistance of the

coating [4]. On one hand, the presence of Cr in the coating

can promote the establishment of an Al2O3 scale on the

coating with much lower Al concentration, which can be

called the third element effect. On the other hand, the

presence of Cr can prevent the martensitic phase transfor-

mation of the b-NiAl phase, which helps to reduce the

degradation of the coating.

Available online at http://link.springer.com/journal/40195

D. Wu � S. Jiang � Q. Fan � J. Gong � C. Sun (&)

State Key Laboratory of Corrosion and Protection, Institute of

Metal Research, Chinese Academy of Sciences,

Shenyang 110016, China

e-mail: csun@imr.ac.cn

123

Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634

DOI 10.1007/s40195-014-0108-5

In this paper, a novel method is investigated to prepare

the Cr-modified aluminide coating. The Cr-modified alu-

minide coating is prepared on a Ni-based superalloy by arc

ion plating (AIP) Cr layer and subsequent pack cementa-

tion aluminizing. Compared with co-deposition of Cr and

Al by powder packing, the compositions of the Cr-modified

aluminide coating prepared by this two-step method can be

well distributed and the contents of Cr and Al can be easily

controlled. The hot corrosion behavior of the Cr-modified

aluminide coating and the aluminized coating in molten

Na2SO4/K2SO4 (3:1, in weight) or Na2SO4/NaCl (3:1, in

weight) salts at 900 �C are studied. The hot corrosion

degradation mechanism of the Cr-modified aluminide

coating and the beneficial effect of the Cr element in the

coating are also discussed.

2 Experimental

A Ni-based superalloy DSM11 with nominal composition

(wt%) of Al 2.9, Ti 4.9, Cr 13.5, Co 9.5, W 3.7, Mo 1.5, Ta

2.8, C 0.1 and balanced of Ni was used as the substrate.

Specimens with dimensions of 15 mm in diameter and

2 mm in thickness were ground with 800-mesh SiC paper,

peened in a wet atmosphere (200-mesh glass ball), and then

ultrasonically cleaned with acetone, ethanol, and deionised

water successively. The Cr-modified aluminide (for short

as Al–Cr) coating was prepared by depositing the Cr layer

firstly and then aluminizing by pack cementation. Before

deposition, bombardment cleaning was carried out after

base pressure of the chamber was pumped below

6 9 10-3 Pa. The working pressure was maintained at

0.2 Pa by flowing argon into the chamber. Detailed depo-

sition parameters are given in Table 1. Pack powder mix-

ture for aluminizing consisted of 96 wt% FeAl powders as

source of aluminum and 4 wt% NH4Cl as activator. The

specimens were buried in the powder mixture. The furnace

chamber was pumped to at least 1 9 102 Pa, then heated to

900 �C. The soaking time was 5 h. The aluminized coating

was also prepared by pack cementation. The contents of the

two coatings are listed in Table 2.

Hot corrosion behaviors of the DSM11 alloy, the alu-

minized coating, and the Al–Cr coating specimens were

performed in a muffle furnace in static air. Before hot

corrosion test, specimens were placed on a hot plate and

then brushed with Na2SO4/K2SO4 (3:1) or Na2SO4/NaCl

(3:1) mixtures. The content of deposited salt on each

sample was about 1 mg/cm2. Afterwards, the samples were

exposed in the muffle furnace at 900 �C to carry out the hot

corrosion testing. At regular intervals of 20 h, the speci-

mens were taken out, cooled down to room temperature

and then washed in boiling distilled water to obliterate the

remained salt. After being dried, the specimens were

weighed by an electronic balance with sensitivity of

0.01 mg. Subsequently, a fresh salt coating was brushed

again to continue the hot corrosion experiment.

Phase identifications of the coating and corrosion scales

were carried out using an X-ray diffractometer (XRD).

Microstructure and morphology were characterized by a

scanning electronic microscope (SEM) equipped with dis-

persive X-ray spectrometer (EDS). Electroless Ni-plating

was plated on the surface of the cross-section samples to

prevent spallation of scales from the surface in sample

preparation process.

3 Results and Discussion

3.1 Microstructures of the Two Coatings

Figure 1 shows the XRD patterns of the aluminized coating

and the Al–Cr coating. Both coatings are composed of b-

NiAl phase and Ni2Al3 phase. The a-Cr phase and AlCr2

phase could be also detected in the Al–Cr coating. But the

diffraction peaks of the AlCr2 phase are weak compara-

tively. In the Al–Cr coating, a-Cr particles are precipitated

near the original substrate surface and also grow inwardly

from the gas phase at the coating surface. Figure 2a and b

show the cross-sectional BSE images of the Al–Cr coating

and the aluminized coating. The thickness of the Al–Cr

coating is about 40 lm, and the aluminized coating is a

little thicker than the former. These two coatings are dense

and adhere tightly to the substrate. The microstructures of

the Al–Cr coating and aluminized coating are characterized

in two distinguished areas: the outer layer and the inter-

diffusion zone. The interdiffusion zone is comprised of b-

Table 1 Deposition parameters of Cr layer by AIP

Operation stage Arc voltage (V) Arc current (A) Bias voltage (V) Bias duty cycle (%) Temperature(�C)

Bombardment cleaning 20–30 60–70 -800 30 50–100

Chromium target 20–30 60–70 -200 30 100–150

Table 2 The compositions of the two coatings (wt%)

Coating Al Ti Cr Co Ni

Al–Cr coating 31.9 0.4 13.2 6.4 Bal.

aluminized coating 25.6 0.5 5.3 6.8 Bal.

628 D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634

123

NiAl and Ni3Al, M23C6, MC-type carbides, and r-phase

[5, 6].

Figure 3a and b show the line scanning images on the

cross-sections of the aluminized coating and the Al–Cr

coating. According to the figures, the distribution and the

relative content of different elements can be seen in the two

coatings. In the aluminized coating, the content of Ni and

Al elements is the highest, and Al element in the outer

layer of the coating distributes uniformly. In the Al–Cr

coating, the content of Ni, Al, and Cr elements rise to the

top with a uniform distribution of Al and Cr element in the

coating. These figures also prove that the structures of the

two coatings are continuous and dense with the main ele-

ments evenly distributed.

3.2 Corrosion Kinetics

Figure 4a, b show the corrosion kinetic curves of the

DSM11 alloy, the aluminized coating and the Al–Cr

coating deposited with salts of Na2SO4/K2SO4 (3:1, and

Na2SO4/NaCl (3:1). The mass change consists of a mass

gain owing to the formation of the scales, and a mass loss

caused by scale spallation and dissolution. The net mass

change of the specimens in molten sulfate represents the

combined effects of these two processes.

The DSM11 alloy, the aluminized coating and the Al–Cr

coating show different corrosion behavior in the molten

Fig. 1 XRD patterns of the aluminized coating and the Al–Cr coating

Fig. 2 Cross-sectional BSE images of the two coatings: a Al–Cr coating; b aluminized coating

Fig. 3 Line scanning images on the cross-sections of the aluminized coating a, the Al–Cr coating b

D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634 629

123

sulfate. According to Fig. 4a, the corrosion of the DSM11

alloy is catastrophic with a sharp increase in the kinetic

curve in the first 20 h. In the first 20 h, the DSM11 alloy

reaches its maximum mass gain of 1.25 mg/cm2. After

40 h, the kinetic curve goes through a sharp decline. It can

be attributed to the low-melting-point phases at the crystal

boundaries as fast diffusion channels. The molten salt can

diffuse through these channels and reacts with the alloy

rapidly. While for the aluminized coating and the Al–Cr

coating, the mass gain is much less, and no rapid mass gain

is observed. It demonstrates that both coatings enhance the

alloy’s anti-corrosion ability greatly. For the aluminized

coating, the mass gain increases in the first 80 h, with a

maximum mass gain of 0.65 mg/cm2 and then starts to fall.

However, the mass gain of the Al–Cr coating is much

lower than that of the aluminized coating throughout the

process. According to Fig. 4a, after corrosion for 40 h, the

mass gain of the Al–Cr coating reaches to 0.35 mg/cm2.

The mass gain data of the Al–Cr coating almost remain

unchanged at the stage of 40 h to 100 h. The kinetic curve

comes into a ‘‘platform’’ stage. It shows that in this hot

corrosion interval, the growth rate of the oxidation film

equaled to its ‘‘dissolved or peeling off’’ rate.

According to Fig. 4b, the DSM11 alloy demonstrates

rather poor anti-corrosion ability in the salt of Na2SO4/

NaCl (3:1). It could be seen that NaCl was more corrosive

than K2SO4. Furthermore, when NaCl is added into

Na2SO4 salt deposit, the exfoliation of corrosion products

happens more easily than in Na2SO4 and K2SO4 salts

deposit. Some studies show that the presence of NaCl can

destroy the continuity of the surface oxidation film [7, 8].

The addition of NaCl will trigger a reaction at high tem-

perature [9]: NaCl (l) ? O2 (g) = 1/2 Na2O (s) ? Cl2(g) to generate Cl2. The Cl2 will react with Cr and Al

elements in the alloy to generate volatile elements such as

chloride. After that, the chloride will spread along the

cracks and voids of the alloy. When reach to the outer

surface of the oxide film will react with O2 again to gen-

erate Cl2. The newly generated Cl2 can diffuse into the

oxide film, and again react with the elements such as Cr

and Al. With the continuous reaction, the whole process of

hot corrosion accelerated. As a result, the mass loss

quantity increased. The corrosion kinetic curve shows a

significant increase of mass gain at initial stage up to 20 h,

and follows an obvious drop after 20 h. The severe mass

losses are caused by significant scale spallation and O2

fluxing, confirming it is necessary to apply protective

coatings on the alloy. For the aluminized coating, a rapid

increase of mass gain occurs in the first 20 h, with a

maximum mass gain of 1.75 mg/cm2 and then starts to fall.

The protective Al2O3 scale starts to crack, and spallation

also occurs. However, for the Al–Cr coating, the mass gain

maintains sustained increasing throughout the corrosion

process. The mass gain owing to the formation of the scales

is much bigger than the mass loss caused by scale spalla-

tion and dissolution. Obviously, the Al–Cr coating provides

much better protection to the alloy than the aluminized

coating in this kind of mixed salt.

3.3 Corrosion Scales

Figure 5a shows the XRD patterns of the surfaces of the

DSM11 alloys with the aluminized coating and the Al–Cr

coating after hot corrosion at 900 �C in the mixture salt of

Na2SO4/K2SO4 (3:1). Besides the Cr2O3 scale, a large

amount of un-protective corrosion products, identified as

Cr3S4 and TiO2 are formed on the surfaces of the samples.

The existence of S in the corrosion products proves that

S2- can diffuse from the surface of the oxidation film into

the substrate through the rapid diffusion channel, which is

formed by the low-melting phase. As for the aluminized

coating and the Al–Cr coating, the corrosion scales formed

on both coatings are mainly composed of a-Al2O3. A large

amount of b-NiAl phase is also detected, which

(a) (b)

Fig. 4 Corrosion kinetic curves of DSM11 substrate, aluminized coating, and Al–Cr coating in the mixed salts of Na2SO4/K2SO4 a, Na2SO4/

NaCl b at 900 �C

630 D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634

123

demonstrates that the Al reservoir in the two coatings is

still abundant. It can promise subsequent formation and

growth of Al2O3 in the later hot corrosion process.

Figure 5b shows the XRD patterns of the corrosion

scales formed on the surfaces of the DSM11 alloys with the

aluminized coating and the Al–Cr coating after hot corro-

sion at 900 �C in the mixture salt of Na2SO4/NaCl (3:1). A

mixed oxidation scale of NiO and Cr2O3 is formed on the

surface of the samples. Meanwhile, NiCr2O4 spinel forms

via solid-phase reaction between NiO and Cr2O3, resulting

in the formation of a mixture layer of NiO, Cr2O3, and

NiCr2O4 spinel. The major phases on the surfaces of the

aluminized coating and the Al–Cr coating are a-Al2O3 and

b-NiAl, while some peaks of Ni3Al are also identified in

the Al–Cr coating.

Figure 6 shows the SEM images of the DSM11 alloy,

the aluminized coating and the Al–Cr coating exposed in

salts of Na2SO4/K2SO4 (3:1) at 900 �C for 100 h, sep-

arately. As it can be seen in Fig. 6a, different structures

of oxide scales are formed on the surface of the DSM11

substrate. These oxidate scales were mixed uniformly

rather than hierarchical, with distinct characteristics of

the rod-like TiO2 oxide. According to Fig. 6b and c,

similar Al2O3 scales are formed on the surface of the

aluminized coating and the Al–Cr coating. Although the

morphologies of Al2O3 scales are similar, however, the

grains sizes of the oxide films are a little different. The

grains of the oxide film on the surface of the Al–Cr

coating are much smaller, which illustrates the growth

rate of the oxide is also much slower. Obviously, the Al–

Cr coating has much better corrosion resistance in this

kind of salt.

Figure 7 shows the cross-sectional BSE morphologies of

the DSM11 alloys with aluminized coating and the Al–Cr

coating exposed in salts of Na2SO4/K2SO4 (3:1) at 900 �C

for 100 h. According to Fig. 7a, the DSM11 alloy suffers

serious hot corrosion with cracks and spallation arise on the

surface. The DSM11 alloy does not form a protective

Al2O3 scale due to its low Al content. By the cross-sec-

tional BSE morphologies (Fig. 7b, c) of the aluminized

coating and the Al–Cr coating show the similar hot cor-

rosion behaviors in this kind of salt. It is found that a dense

and intact layer of a-Al2O3 is formed on the surface of both

coatings. Besides, the Al reservoir phase b-NiAl is extre-

mely abundant in both coatings, which agrees well with the

XRD patterns shown in Fig. 5a.

(b)(a)

Fig. 5 XRD patterns of the DSM11 substrate, the aluminized coating and the Al–Cr coating after corrosion at 900 �C for 100 h in mixed salts of

Na2SO4/K2SO4 a, Na2SO4/NaCl b

Fig. 6 SEM images showing the surfaces of DSM11 substrate a, aluminized coating b, Al–Cr coating c after corrosion at 900 �C for 100 h in the

salt of Na2SO4/K2SO4

D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634 631

123

Figure 8 shows the SEM images of the DSM11 alloy

with aluminized coating and the Al–Cr coating after cor-

rosion in the mixed salt of Na2SO4/NaCl (3:1) at 900 �C

for 100 h. According to Fig. 8, the mixture oxide scales

including NiO, Cr2O3, and NiCr2O4 spinel are formed on

the surfaces of the samples. The oxide formed on the sur-

face of the aluminized coating, and the Al–Cr coating is sill

a-Al2O3, while the density of the oxide declines.

Figure 9 shows the cross-sectional BSE morphologies of

the DSM11 alloys with aluminized coating and the Al–Cr

coating after corrosion in the mixed salt of Na2SO4/NaCl

(3:1) at 900 �C for 100 h. As seen in Fig. 9a, the DSM11

alloy suffers more severe hot corrosion than that in the salt of

Na2SO4/K2SO4 (3:1). When NaCl is introduced into the salts

film, the corrosion is more serious since internal oxide and

chromium sulfide emerged in the two coatings. The major

role of NaCl in hot corrosion induced by Na2SO4–NaCl

mixtures is to cause cracking of the protective oxide scale as

well as to produce the internal voids by means of oxychlo-

rination and chlorination/oxidation cyclic reactions,

Fig. 7 Cross-sectional BSE images of DSM11 substrate a, aluminized coating b, Al–Cr coating c after corrosion at 900 �C for 100 h in the salt

of Na2SO4/K2SO4

Fig. 8 SEM images showing the surfaces of DSM11 substrate a, aluminized coating b, Al–Cr coating c after corrosion at 900 �C for 100 h in the

salt of Na2SO4/NaCl

Fig. 9 Cross-sectional BSE images of DSM11 substrate a, aluminized coating b, Al–Cr coating c after corrosion at 900 �C for 100 h in the salt

of Na2SO4/NaCl

632 D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634

123

additionally resulting in the dissolution of the protective

oxide scale [10–13]. According to Fig. 9b, the protective a-

Al2O3 scale on the surface of the aluminized coating is

depleted and a continuous a-Al2O3 scale does not exist after

corrosion for 100 h. The interdiffusion zone disappears

completely and heavy internal oxidation and sulfidation are

found beneath the surface. The inward diffusion of S would

increase the basicity of the molten sulfates, leading to the

dissolution of Al2O3, since alumina can dissolve in basic

sulfate, but be stable in neutral salt. Obviously, the alumi-

nized coating almost loses its protective effect against this

kind of hot corrosion. As seen in Fig. 9c, a very dense Al2O3

scale also exists on the surface of the Al–Cr coating after

corrosion for 100 h, which is mainly related to the NiAl

phase and Cr element in the coating. The addition of Cr

element in the coating can promote to form the Al2O3 scale

on the coating with much lower Al concentration, which can

be called the third element effect [14]. Besides, slight

internal oxidation is found under the Al2O3 scale. According

to the above analyses, a certain amount of b-NiAl phase still

remains in the coating and the scale is also continuous and

compact, which testifies the Al–Cr coating can still protect

the substrate from corrosion effectively.

3.4 Hot Corrosion Mechanism of the Al–Cr Coating

According to Ref. [7] and Ref. [8], hot corrosion is an

accelerated form of oxidation that occurs when metals are

heated in the temperature range 700–900 �C. Hot corrosion

problems are a direct result of the combination of salt

contaminants such as Na2SO4, NaCl, and V2O5, which

produces low-melting point deposits that can dissolve the

protective surface oxides [15]. A number of fluxing

mechanisms has been proposed to account for the different

corrosion morphologies that are observed [16–18] and this

has resulted in the general classification of high tempera-

ture (type I, 800–950 �C) hot corrosion and low tempera-

ture (type II, 600–800 �C) hot corrosion. The degradation

sequence usually consists of an initiation stage, during

which the attack is virtually the same as for the alloy in the

absence of deposit, and a propagation stage, during which

the attack is substantially increased.

The Al–Cr coating presents excellent corrosion resis-

tance in the molten salt with low corrosion rate mainly due

to the beneficial effect of the Cr element in the out layer of

the coating. The molten Na2SO4, which deposits on the

surface of the superalloy, can cause hot corrosion of the

alloy and flux the normally protective oxides such as Cr2O3

and Al2O3 [19, 20]. In a liquid deposit, there are the fol-

lowing thermodynamic equilibriums:

Na2SO4 ¼ Na2Oþ SO3: ð1Þ

SO3 ¼1

2S2 þ

3

2O2: ð2Þ

The experimental results presented in this paper are

consistent with the observations proposed by other

researchers [21–24] on the beneficial effect of chromium

in inhibiting the hot corrosion of an aluminide coating.

According to Otsuka and Rapp [25], an oxide scale that

contains Cr2O3 will dissolve partially in a Na2SO4 melt by

a basic dissolution reaction to form chromate anions:

Cr2O3 þ3

2O2 þ 2SO2�

4 ¼ 2CrO2�4 þ 2SO3: ð3Þ

Similarly, if Al2O3 is present in the scale it can also

undergo basic dissolution:

Al2O3 þ SO2�4 ¼ 2AlO�2 þ SO3: ð4Þ

These two reactions result in the consumption of sulfate

ions. However, SO42- consumption by the reaction of

forming CrO42- is much more effective than by AlO2

-

formation. This is primarily a consequence of the greater

solubility of the chromate anion in the basic salt melt [26].

From Eq. (4), if the Cr2O3 reaction consumes a sufficient

amount of the sulfate ions, the basic dissolution of the

Al2O3 can be suppressed. As a result, the Al2O3-rich scale

is able to act as an effective protective barrier between the

coating and the salt medium.

Later, as the hot corrosion process proceeds, when most

of the Al reservoirs are sacrificed and the internal layer

enriching of Cr is exposed, the Cr will supersede the Al to

form a protective Cr2O3 scale [27]. Since the fluxing of

Cr2O3 primarily obeys the equation:

Cr2O3 þ O2� þ 3

2O2ðgÞ ! Cr2O2�

7 : ð5Þ

The dissolution of chromia needs the assistance of oxygen

[28]. Thus, the solubility of Cr2O3 at the salt/gas interface

is higher than that at the oxide/salt, and there will exist as a

‘‘positive solubility gradient’’ which is not self-sustainable

for Cr2O3 to be fluxed. Moreover, the Cr can sacrifice itself

to entrap the element sulfur to form a relatively stable

Cr2S3 [29] as long as the chromium sulfide is effectively

insulated from further oxidation.

As reported by Deb et al. [13], chlorides can cause the

formation of volatile species, which form voids and pits at

grain boundaries, thus forming an easy path for flowing

corrodents and oxygen. The deleterious feature of the

sodium chloride should be attributed to the extremely

small ion of Cl-. Since the melting point (801 �C) of

NaCl is lower than that of Na2SO4 (884 �C), the molten

mixed salt at the current exposure temperature establishes

an easier corrosive electrolyte where the Cl- can travel

freely.

D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634 633

123

As the protective feature of alumina scale has been

deteriorated, oxygen and sulfur could easily penetrate into

the coatings and cause the formation of internal oxide and

chromium sulfide [30, 31]. These internal oxides and

chromium sulfides would accelerate the consumption of

benefit element, and consequently speed up the corrosion

process. Compared with the aluminized coating, the Al–Cr

coating has sufficient aluminum to maintain a self-healing

ability and prevent the inward diffusion of oxygen and

sulfur, accompanying with inhibiting the invasion of NaCl.

In addition, it can be seen that the Cr-rich inner layer of the

Al–Cr coating is not attacked yet. So, it can be predicable

that the Al–Cr coating would resist the corrosion attack

much longer than the aluminized in its service life.

4 Conclusions

(1) The Cr-modified aluminide coating is prepared by a

novel method with two steps of deposition of AIP Cr

and then Al pack cementation. The coating is com-

prised essentially of an outward-grown b-NiAl matrix

with a-Cr precipitates.

(2) The Cr-modified aluminide coating and the alumi-

nized coating both show excellent corrosion resis-

tance in the mixture salt of Na2SO4/K2SO4 (3:1). The

Cr-modified aluminide coating presents much better

corrosion resistance than the aluminized coating.

(3) The Cr-modified aluminide coating still possesses

good corrosion resistance in the mixture salt of

Na2SO4/NaCl (3:1) due to the beneficial effect of Cr

in the coating. However, the aluminized coating has

been damaged seriously and lost its protective effect

after corrosion.

Acknowledgments This work was financially supported by the

National Natural Science Foundation of China (No. 51001106) and

National Basic Research Program of China (No. 2012CB625100).

References

[1] J.R. Nicholls, D.J. Stephenson, Met. Mater. 7, 156 (1991)

[2] J. Ma, S.M. Jiang, H.Q. Li, W.X. Wang, J. Gong, C. Sun,

Corros. Sci. 53, 1417 (2011)

[3] J. Smialek, Metall. Trans. A 18, 164 (1987)

[4] E. Godlewska, K. Godlewski, Oxid. Met. 22, 117 (1984)

[5] G.W. Goward, D.H. Boone, Oxid. Met. 3, 475 (1971)

[6] W.F. Gale, J.E. King, Metall. Trans. A 23, 2657 (1992)

[7] P. Hancock, Mater. Sci. Technol. 3, 536 (1987)

[8] N. Eliaz, G. Shemesh, R.M. Latanision, Eng. Fail. Anal. 9, 31

(2002)

[9] E.L. Simons, G.V. Browning, H.A. Liebhafsky, Corrosion 11,

505 (1955)

[10] D.W. McKee, D.A. Shore, K.L. Lurthra, J. Electrochem. Soc.

125, 411 (1978)

[11] I. Gurrappa, Oxid. Met. 51, 353 (1999)

[12] M.K. Hossain, S.R.J. Saunders, Oxid. Met. 12, 1 (1978)

[13] D. Deb, S.R. Iyer, V.M. Radhakrishnan, Mater. Lett. 29, 19

(1996)

[14] W.C. Hagel, Corrosion 21, 316 (1965)

[15] J.R. Nicholls, N.J. Simms, W.Y. Chan, H.E. Evans, Surf. Coat.

Technol. 149, 236 (2002)

[16] N. Birks, G.H. Meier, in Introduction to High Temperature

Oxidation of Metals, (Edward Arnold, London, 1983)

[17] P. Kofstad, in High Temperature Corrosion, (Elsevier Applied

Science Publishers Ltd., New York, 1988)

[18] C.S. Giggins, F.S. Pettit, France patent 11545 (1979)

[19] S.W. Yang, Oxid. Met. 15, 375 (1981)

[20] P.S. Liu, K.M. Liang, H.Y. Zhou, S.R. Gu, A.D. Lin, X.F. Sun,

H.R. Guan, T. Jin, K.N. Yang, Surf. Coat. Technol. 52, 497

(2001)

[21] R. Bianco, R.A. Rapp, J. Electrochem. Soc. 137, 211 (1990)

[22] R. Bianco, M.A. Harper, R.A. Rapp, J. Met. 43, 68 (1991)

[23] R. Sivakumar, Oxid. Met. 17, 27 (1982)

[24] K. Godlewski, E. Godlewska, Oxid. Met. 26, 125 (1986)

[25] N. Otsuka, R.A. Rapp, J. Electrochem. Soc. 137, 53 (1990)

[26] R.A. Rapp, Corrosion 42, 568 (1986)

[27] Z.B. Bao, Q.M. Wang, W.Z. Li, X. Liu, J. Gong, T.Y. Xiong, C.

Sun, Corros. Sci. 51, 860 (2009)

[28] R.A. Rapp, Corros. Sci. 44, 209 (2002)

[29] I. Gurrappa, Surf. Coat. Technol. 139, 272 (2001)

[30] S.M. Jiang, X. Peng, Z.B. Bao, S.C. Liu, Q.M. Wang, J. Gong,

C. Sun, Corros. Sci. 50, 3213 (2008)

[31] Q.M. Wang, Y.N. Wu, P.L. Ke, H.T. Cao, J. Gong, C. Sun, L.S.

Wen, Surf. Coat. Technol. 186, 389 (2004)

634 D. Wu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(4), 627–634

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