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Corrosion Science xxx (xxxx) xxx–xxx
Contents lists available at ScienceDirect
Corrosion Science
journal homepage: www.elsevier.com/locate/corsci
Effects of alloying elements on the corrosion behavior of Ni-based alloys in
molten NaCl-KCl-MgCl2 salt at different temperatures
⁎
Hua Sun , Peng Zhang, Jianqiang Wang
Shanghai Institute of Applied Physics, Chinese Academy of Sciences, 2019 Jialuo Road, 201800 Shanghai, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords:
Ni-based alloy
Molten chloride salt
Corrosion
Alloying element
The effects of alloying elements on corrosion of Ni-based alloys in NaCl-KCl-MgCl2 under Ar in the temperature
range of 600∼800 °C were investigated by immersion test, SEM/EDS, EPMA and XRD techniques. A thermodynamic potential-pO2− diagram for Cr, Ni at 700 °C was constructed. The alloy corrosion is attributed to the
dissolution of Cr to form Cr chlorides into salt. The elements Cr, Mo and W can affect the alloy corrosion. The
corrosion is accelerated and the intergranular corrosion becomes more pronounced with temperature.
Furthermore, the effects of the alloying elements on corrosion have a strong temperature dependence. The
related mechanisms are discussed.
1. Introduction
Molten chloride salts are promising candidates of heat transfer fluid
(HTF) and thermal energy storage (TES) media for next generation
concentrating solar power (CSP) plants because of their beneficial
characteristics such as low cost, low melting point, high boiling point
and good heat transfer property [1–4]. However, a main challenge
using molten chloride salts in next generation CSP plants is the corrosion problem of structural materials at the aim operating temperatures
(> 600 °C) [1,4,5]. Therefore, the material degradation caused by
corrosion is one of key issues related to next generation CSP plants.
Ni-based alloys are considered as candidate structural materials for
next generation CSP plants due to combined good mechanical properties and corrosion resistance [4,6]. Much work was denoted to investigate the corrosion behavior of various Ni-based alloys in molten
chloride salts in order to evaluate their corrosion resistance [7–15].
Vignarooban et al. [7] investigated that the corrosion behavior of
Hastelloy N, C276, C22 in molten NaCl-KCl-ZnCl2 at 250 °C and 500 °C
and found that Hastelloy N showed the highest corrosion rate due to its
low Cr content. However, Ambrosek [8] evaluated the corrosion resistance of various Ni-based alloys including Hastelloy N, Inconel 617,
Inconel 625, Inconel 718 and Haynes 230 etc. in molten KCl-MgCl2 at
850 °C and pointed out that Hastelloy N exhibited the least attack depth
because of its low Cr content. Alkhamis [9] compared the corrosion
performance of Hastelloy C276 and Haynes 230 in molten MgCl2-KCl
and NaCl-KCl-ZnCl2 at 800 °C and found that Hastelloy C276 with high
Mo showed better corrosion resistance than Haynes 230 with high W.
⁎
On contrary, Oryshich and Kostyrko [10] demonstrated that Mo and W
were similar in improving the corrosion resistance of Ni-based alloys in
molten chloride salts. It can be seen that the alloying elements such as
Cr, Mo and W can affect the resistance of Ni-based alloys against molten
chloride salts. Nevertheless, no systematic studies have been performed
to clarify the relationship between the alloying elements and the corrosion behavior of Ni-based alloys in molten chloride salt. Since the CSP
plants will use mature commercial alloys with various complex alloying
elements, understanding how the alloying elements affect the corrosion
behavior of commercial alloys is crucial.
In addition, it is generally recognized that the corrosion of the alloys
in molten chlorides salt is attributed to the selective dissolution and/or
oxidation of the element Cr depending on the different experimental
conditions. The corrosion products covered on the alloy surfaces are
usually characterized as Cr-rich oxides such as Cr2O3, MgCr2O4,
NiCr2O4 or LiCrO2 according to the different molten salt systems
[11–13]. Nevertheless, the major existing form of Cr dissolved into the
molten salt is unknown due to the difficulties associated with characterizing the species with low concentration in molten salt. Therfore,
the corrosion mechanism of the alloys in molten chloride salts, especially in the molten salts under inert gas is not well understood. In fact,
it is convenient to predict the corrosion products of the alloys in molten
salt by constructing the potential-oxoacidity (E-pO2−) relationship
[16–19].
The objective of this work is to investigate the effects of the alloying
elements Cr, Mo and W on the corrosion behavior of commercial Nibased alloys in molten chloride salt under Ar atmosphere in the
Corresponding author.
E-mail address: sunhua@sinap.ac.cn (H. Sun).
https://doi.org/10.1016/j.corsci.2018.08.021
Received 19 November 2017; Received in revised form 31 July 2018; Accepted 8 August 2018
0010-938X/ © 2018 Elsevier Ltd. All rights reserved.
Please cite this article as: Sun, h., Corrosion Science (2018), https://doi.org/10.1016/j.corsci.2018.08.021
Corrosion Science xxx (xxxx) xxx–xxx
H. Sun et al.
Table 1
Concentration of main impurities in as-received salts (ppm, by weight).
Salt
Cr
Fe
Ni
Zn
Co
Mo
SO42−
PO43−
NO3−
NaCl
KCl
MgCl2
1.01
1.12
0.32
1.64
1.39
1.88
0.52
< 0.005
/
0.50
0.48
0.49
1.0
1.0
1.06
1.49
1.48
1.23
< 25
< 25
< 25
< 25
< 25
< 25
< 25
< 25
< 25
temperature range of 600–800 °C by means of weight change measurement and microstructure characterization techniques. A thermodynamical E-pO2− diagram of Cr, Ni is constructed. The related mechanisms are also discussed.
2. Experimental methods
Fig. 1. Schematic diagram of the experimental setup.
2.1. Preparation of salt mixture and specimens
Afterwards, the Al2O3 crucibles were divided into three groups according to the experimental temperatures, and subsequently the Al2O3
crucibles of the same group were encapsulated into the same 316SS
crucible. Ar arc welds were used to seal the 316SS crucibles. Notably,
all of these procedures were performed inside the glove box full of highpurity Ar (99.999%), where H2O and O2 contents were 2 ppm and
0.5 ppm (mole fraction), respectively and the system pressure was
about 1 atm. Correspondingly, H2O and O2 partial pressures in the experimental system were calculated to be 2 × 10−6 and 5 × 10-7 atm,
respectively. After the crucibles were sealed, the corrosion tests were
performed in the high temperature resistance furnaces outside the glove
box. The furnaces were heated to 600 °C, 700 °C and 800 °C, respectively and held at each temperature for 400 h.
After the tests, the furnaces were cooled to room temperature. The
crucibles were removed and cut using a mechanical cutting machine.
The specimens were retrieved and cleaned ultrasonically in deionized
water and alcohol to remove the residual salt. The specimens were then
dried and weighed to obtain the weight changes with a accuracy of
± 0.01 mg. The surface and cross-sectional images of the specimens
were examined using scanning electron microscopy (SEM, Merlin
Compact). The elemental distributions of the specimens were analyzed
by electron probe micro-analysis (EPMA, Shimadzu EPMA-1720). The
precipitated phases on the specimens were identified by energy dispersive X-ray spectroscopy (EDS, LEO 1530 V P) and X-ray diffraction
(XRD, D8 Advance). It should be noted that the specimens for crosssectional analyses were embedded into epoxy resin and again abraded
and polished. Before measurement, the prepared cross-sections were
sprayed Au or Pt to enhance the conductance of the specimens.
The ternary eutectic NaCl-KCl-MgCl2 (33-21.6–45.4 mol%) was selected as the experimental molten salt. Before tests, the salt mixture was
prepared using commercial NaCl, KCl and MgCl2 (greater than 99.5 wt
% purity), supplied by Sinopharm Chemical Reagent Corporation
(Shanghai, China). The concentration of main impurities in as-received
salts is listed in Table 1, which was analyzed by inductively coupled
plasma-optical emission spectroscopy (ICP-OES) and ion chromatography (IC). The salts were weighed and mixed according to the above
ratio, and then placed in Al2O3 crucibles. The salts were oven dried at
300 °C for 24 h. After that, the salts were transferred rapidly into a glove
box full of high-purity Ar (99.999%) to minimize moisture absorption.
Four commercial Ni-based alloys including three Ni-Mo-Cr alloys
(Hastelloy N, C276 and C22) and one Ni-W-Cr alloy (Haynes 230) were
selected as the alloys tested according to their chemical compositions,
as seen in Table 2. The Mo content (16.8 wt.%) of Hastelloy N is similar
to that (15.6 wt.%) of C276, while the Cr content (7.01 wt.%) of Hastelloy N is obviously lower than that (16.2 wt.%) of C276. The Cr
content (21.6 wt.%) of Hastelloy C22 is similar to that (21.6 wt.%) of
Haynes 230, while the Mo content (13.1 wt.%) of C22 is replaced by the
W content (13.4 wt.%) of Haynes 230. The specimens with a size of
10 mm × 10 mm × 2 mm were abraded with emergy paper to 1500 grit
and polished with 0.05 μm Al2O3 power, followed by cleaning with
deionized water and alcohol and then drying.
2.2. Experimental setup and procedure
The corrosion tests were performed in an experimental setup with
double layer crucibles, as shown in Fig. 1. Al2O3 crucibles were chosed
as the inner layer crucibles because of their chemical inertness in
molten chloride salts [5,20]. 316 stainless steel (316SS) crucibles with
lids were chosed as the outer layer crucibles in order to avoid the effects
of H2O and O2 from the external environment on the experimental results. Before corrosion tests, Al2O3 crucibles, 316SS crucibles and lids
were cleaned with deionized water and alcohol and then dried at 120 °C
for 24 h to remove residual moisture.
Three specimens of each alloy were placed horizontally in each
Al2O3 crucible in order to ensure the same contact surfaces between the
specimens and the molten salt. Solid NaCl-KCl-MgCl2 (3321.6–45.4 mol%) salt of 40 g was then added into each Al2O3 crucible.
3. Results
3.1. Weight change of the alloys
Fig. 2 shows the weight changes of the Ni-based alloys exposed in
molten NaCl-KCl-MgCl2 under Ar atmosphere in the temperature range
of 600∼800 °C for 400 h. All alloys underwent weight loss, except that
Haynes 230 showed weight gain at 800 °C. The weight loss of the alloys
at 700 °C was obviously larger than that at 600 °C, indicating a degraded corrosion resistance. In addition, the weight loss of Haynes 230
was obviously larger than that of the other Ni-based alloys at 600 °C and
Table 2
Chemical compositions of the alloys tested (wt.%).
Alloy
Ni
Mo
Cr
Fe
Co
Mn
Al
Ti
Si
C
W
HastelloyN
C276
C22
Haynes230
70.56
58
Bal
Bal
16.8
15.6
13.1
1.2
7.01
16.2
21.6
21.6
4.16
5.9
3.7
1.17
0.002
0.17
0.56
0.24
0.52
0.4
0.26
0.47
0.15
0.002
0.36
0.015
0.055
0.004
< 0.001
0.1
3.2
2.9
13.4
2
0.39
< 0.01
0.37
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H. Sun et al.
Fig. 4 shows the cross-sectional EPMA mappings of the alloys exposed in molten NaCl-KCl-MgCl2 under Ar atmosphere at 600 °C for
400 h. Hastelloy N, C276 and C22 exhibited little Cr depletion with a
depth less than 5 μm, while the Cr depletion depth of three Ni-Mo-Cr
alloys seemed to increase in the following order: Hastelloy N <
C276 < C22. Haynes 230 showed non-uniform Cr depletion and its
maximum Cr depletion depth was up to 30 μm. No significant Mo and W
depletion was detected on the Ni-Mo-Cr and Ni-W-Cr alloys, suggesting
that Mo and W exhibited the excellent resistance against molten
chlorides. Hosoya et al. [21] also pointed out that Mo and W can not be
attacked in molten chloride salts. It should be noted that Mo was enriched in the Cr depletion regions of C276 and C22, which could be
attributed to the diffusion of Mo from the alloy base to the surface
because C276 and C22 suffered more severe corrosion than Hastelloy N.
In addition, some Mo-rich and W-rich phases were distributed randomly on Hastelloy N and Haynes 230 matrices, respectively, which
were likely related to Mo-rich and W-rich carbides.
Fig. 5 shows the cross-sectional EPMA mappings of the alloys exposed in molten NaCl-KCl-MgCl2 under Ar atmosphere at 800 °C for
400 h. The corrosion of the alloys was still attributed to the Cr depletion. Compared with 600 °C, the Cr depletion depths of all tested alloys
inceased and the Cr depletion along the grain boundaries became more
marked. The results reveal that the corrosion of the alloys is accelerated
and shifted gradually from general corrosion to intergranular corrosion
with increasing temperature in accordance with the above SEM results.
The extent of Cr depletion of four alloys was very different at 800 °C.
The Cr depletion depth of Hastelloy N was about 30 μm, whereas that of
C276 was increased to about 50 μm. Although there was no significant
difference in the Cr depletion depths of C22 and C276, more Cr was
dissolved in the Cr depletion layer of C22. The Cr depletion layer of
Haynes 230 was almost throughout the whole specimen, and Cr depletion primarily occurred along the grain boundaries in addition to the
uniform Cr depletion near the surface of the alloy. The EPMA results
were also good consistent with the above SEM results, that is, the alloying elements Cr, Mo and W can affect the corrosion of Ni-based alloys, resulting that the corrosion resistance of the alloys degraded in the
following order: Hastelloy N > C276 > C22 > Haynes 230.
Furthermore, the effects of the alloying elements on corrosion have a
strong temperature dependence.
Compared with 600 °C, a number of Mo-rich precipitated phases
were observed along the grain boundaries of the Ni-Mo-Cr alloys at
800 °C. It should be noted that the Mo-rich phases near the Hastelloy N
surface reduced versus the alloy base, whereas the Mo-rich phases may
be concentrated on the surfaces of C276 and C22. In contrast, no significant W-rich phases were detected at the grain boundaries of Haynes
230. Furthermore, the W-rich phases in the near-surface region of the
alloy obviously reduced at 800 °C, suggesting that the W-rich phases
could be dissolved during corrosion. Otherwise, the Ni and W depletion
layers with a depth of about 20 μm were also detected in the nearsurface region of Haynes 230 at 800 °C, indicating that the elements Ni
and W could suffer slight dissolution.
Besides, the significant amounts of O element were observed in the
Cr depletion regions of the alloys exposed in molten NaCl-KCl-MgCl2
under Ar atmosphere at 800 °C for 400 h, as shown in Fig. 6. Therefore,
it may be assumed that the enriched O mainly came from the insoluble
species in molten salt instead of the corrosion products of the alloys,
which caused the weight gains of the alloys. Furthermore, the amount
of enriched O increased with the accelerated corrosion, resulting that
the alloys exhibited smaller weight loss and even weight gain even
though they suffered more severe corrosion as evidenced in Fig. 2. The
concrete reasons for the O enrichment will be further discussed.
As we known, the composition and structure of the precipitated
phases at grain boundaries are closely related to the corrosion behavior
of the alloys. Therefore, the cross-sections of various alloys were etched
with metallographic etchant to further study the differences of the
precipitated phases using SEM/EDS analysis, and the corresponding
Fig. 2. Weight changes of Ni-based alloys exposed in molten NaCl-KCl-MgCl2
under Ar in the temperature range of 600∼800 °C for 400 h.
700 °C, suggesting that the Ni-W-Cr alloy exhibited worse corrosion
resistance than the Ni-Mo-Cr alloys. However, the weight changes of
four Ni-based alloys followed the different tendency as temperature
increased from 700 °C to 800 °C. The weight loss of Hastelloy N increased, and the weight of Hastelloy C276 and C22 showed little
change, while Haynes 230 shifted from weight loss to weight gain. As a
result, it is not possible to accurately assess the corrosion resistance of
the alloys only according to the result of weight changes. The further
microstructure analyses are necessary.
3.2. Cross-section analyses
Fig.3 shows the cross-sectional SEM images of the alloys exposed in
molten NaCl-KCl-MgCl2 under Ar atmosphere in the temperature range
of 600∼800 °C for 400 h. At 600 °C, Hastelloy N showed slight corrosion attack in the form of shallow subsurface voids with a depth of
about 4 μm. At 700 °C, the alloy exhibited obvious corrosion with a
depth of about 15 μm. The voids were distributed uniformly in the nearsurface region of the alloy, while particularly pronounced along the
grain boundaries close to the alloy base. At 800 °C, the alloy corrosion
was further accelerated. Apart from small amounts of surface voids,
Hastelloy N exhibited intergranular corrosion and the whole corrosion
depth was about 36 μm. These results reveal that the corrosion of
Hastelloy N is accelerated and the intergranular corrosion becomes
pronounced with increasing temperature.
The similar phenomena were also observed on C276 and C22.
However, the corrosion grow rates of three Ni-Mo-Cr alloys were different with temperature. At 600 °C, the corosion depths of C276, C22
and Hastelloy N were very similar. However, the corrosion depths of
C276 were increased to 25 μm at 700 °C and more than 50 μm at 800 °C,
respectively, which were much higher than that of Hastelloy N. The
corrosion depths of C22 were the same as that of C276 at 700 °C and
800 °C, whereas its void density and the general corrosion depth were
significantly higher than that of C276.
Haynes 230 underwent more severe corrosion than the above NiMo-Cr alloys. Its local maximum corrosion depth increased sharply
from 25 to 100 μm as temperature increased from 600 °C to 700 °C,
while the corrosion voids were almost throughout the whole specimen
at 800 °C. It is concluded that the corrosion resistance of the Ni-Mo-Cr
alloys in molten chloride salt increases with an increase of Mo content
and a decrease of Cr content. Moreover, three Ni-Mo-Cr alloys show
better corrosion resistance than the Ni-W-Cr alloy. Furthermore, the
differences in the corrosion resistance of the alloys become more significant with temperature.
3
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H. Sun et al.
Fig. 3. Cross-sectional SEM images of Hastelloy N(a,b,c), C276(d,e,f), C22(g,h,i) and Haynes 230(j,k,l) exposed in molten NaCl-KCl-MgCl2 under Ar for 400 h at
600 °C(a,d,g,j), 700 °C(b,e,h,k) and 800 °C (c,f,i,l).
observed on the surfaces of the Ni-Mo-Cr alloys. As a replacement,
numerous voids were distributed uniformly on the alloy surfaces at
600 °C (Fig. 8a, c and e), while the voids were mainly presented at the
grain boundaries of the alloys at 800 °C (Fig. 8b, d and f). The results
were consistent with the cross-sectional SEM and EPMA results.
Otherwise, many Mo-rich phases were segregated around the voids on
the surfaces of the Ni-Mo-Cr alloys based on the surface SEM/EDS results. The XRD analysis showed these Mo-rich phases primarily presented as MoNi3 and small amounts of Mo2C, which were different from
those distributed on the alloy bases (Fig. 7). Combining with the surface
and cross-section results, therefore, it can be assumed that the element
Mo can diffuse from the alloy base to the surface and subsequently form
MoNi3 and Mo2C around the voids during the corrosion process of the
Ni-Mo-Cr alloys. Luo et al. [22] also found that Mo can transport onto
the surface of Inconel 617 exposed in molten LiF-MgF2-KF, which can
be attributed to the vacancy difference between the alloy base and the
Cr depletion layer. Similarly, Yin et al. [23] also found the Mo-rich
phases within the Cr depleted regions of Hastelloy N after corrosion in
FLiNaK salt at 700 °C. Moreover, the amount of Mo-rich phases in
Hastelloy N increased with the acceleration of corrosion. The Mo-rich
phases around the voids are taken as a reason for better resistance of the
Ni-Mo-Cr alloys. It is worth mentioning that the Ni enrichment was also
found in the general corrosion layers of C276 and C22 (Fig. 4b, c and 5
results are shown in Fig. 7. It can be seen that the significant continuous
precipitated phases were found at the grain boundaries of four Ni-based
alloys, even though they were not observed at the grain boundaries of
Haynes 230 by the above EPMA analysis. According to the EDS results,
the precipitated phases on the alloys consisted of carbides, whereas
their chemical compositions were different. The carbides on the Ni-MoCr alloys were composed of Ni, Mo, Cr and C, and the Cr content of
carbides increased with increasing the Cr content of the alloys. In
general, however, the carbides on the Ni-Mo-Cr alloys were enriched in
Mo and Ni, and the maximum Cr content of the carbides was not above
20 wt.%. In contrast, the carbides on the Ni-W-Cr alloy consisted of Ni,
W, Cr and C. Furthermore, the carbides were enriched in Cr (above
45 wt.%) which was far more than that of the carbides at the grain
boundaries of the Ni-Mo-Cr alloys. The differences in the chemical
compositions of the carbides presented at the grain boundaries were
probably responsible for the different resistance of various alloys
against the molten chloride salt.
3.3. Surface analyses
Figs. 8 and 9 show the surface SEM/EDS results and XRD patterns of
various alloys exposed in NaCl-KCl-MgCl2 under Ar atmosphere for
400 h at 600 °C and 800 °C, respectively. No corrosion products were
4
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H. Sun et al.
Fig. 4. Cross-sectional EPMA mappings of various alloys exposed in molten NaCl-KCl-MgCl2 under Ar at 600 °C for 400 h: (a)Hastelloy N; (b) C276; (c) C22; (d)
Haynes 230.
and XRD results, while the Mo-rich phases decreased near the surface of
Hastelloy N at 800 °C based on the EPMA results. The concrete reason
for the discrepancy will be analyzed in the future work.
The Ni-W-Cr alloy, i.e. Haynes 230 suffered intergranular corrosion
b, c). Furthermore, the thickness of Ni-rich layers increased with temperature, which could be due to the leftovers of base element Ni caused
by Cr depletion. In addition, many Mo-rich phases can be observed on
the surface of Hastelloy N at 600 °C and 800 °C based on the SEM/EDS
5
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H. Sun et al.
Fig. 5. Cross-sectional EPMA mappings of various alloys exposed in molten NaCl-KCl-MgCl2 under Ar at 800 °C for 400 h: (a)Hastelloy N; (b) C276; (c) C22; (d)
Haynes 230.
ascribed to MgO combining with the XRD results (Fig. 9d). While the
irregular phases distributed uniformly on the alloy surface mainly
consisted of Ni element, which could be attributed to the corroded alloy
base. It appeared that most of alloy base was dissolved and therefore
and no corrosion products were found on the alloy surface at 600 °C
(Fig. 8g). In contrast, two different phases were observed on the surface
of Haynes 230 at 800 °C (Fig. 8h). The large particles on the top surface
with diameters of 10∼20 μm were composed of Mg and O, which were
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H. Sun et al.
Fig. 6. EPMA mappings of O element on various alloys exposed in molten NaCl-KCl-MgCl2 under Ar at 800 °C for 400 h: (a)Hastelloy N; (b) C276; (c) C22; (d) Haynes
230.
activity in molten salt, i.e. pO 2 − = −log α O2 − . The calculation procedure
is described in detail elsewhere [17,18]. The diagram is divided into
three regions: immunity, oxidation and dissolution. The domain of
immunity on the diagram defines the combination of potential and
pO2− where corrosion does not occur. The oxidation domain corresponds to the combination of potential and pO2− where the alloying
elments are oxidized to form their oxides. The dissolution domain
means that the combination of potential and pO2− where the alloying
elements are dissolved to form their chlorides. This suggests that the
stability domains of Cr and Ni can be established as a function of potential and pO2−.
In the present work, H2O and O2 partial pressures were evaluated to
be 6.5 × 10−6 and 1.6 × 10−6 atm at 700 °C, respectively, according to
the ideal gas law pV = nRT . H2O and O2 were converted into HCl and
Cl2 by the following reactions (1) and (2) [24–27].
significant Ni and W depletion was also detected in the near-surface
region of the alloy as observed in the EPMA results (Fig. 5d). The result
further indicated that the substrate of Haynes 230 suffered severe
corrosion at 800 °C.
4. Disscusion
4.1. Thermodynamic analysis about the alloy corroison
The SEM and EPMA results revealed that the corrosion of the alloy
was mainly ascribed to the dissolution of Cr. Similar results were also
reported in the works of Ambrosek, Alkhamis and Cho et al. [8,9,14]
about the corrosion of Ni-based alloys in the molten chloride salt under
inert atmosphere. However, there is little discussion on the chemical
composition of Cr dissolved in molten chloride salt.
Indeed, the corrosion of the alloys in molten salt is electrochemical
process in natural [16]. As already mentioned, the stable corrosion
products of the alloys in molten salt can be predicted by the thermodynamic E-pO2− relationship. In order to better understand the corrosion mechanisms, therefore, the E-pO2- diagram for the active element
Cr and base element Ni at 700 °C in molten chloride salt was calculated
based on the thermodynamics software of HSC 6.0, as shown in Fig. 10.
Where E represents the equilibrium electrode potentials of various
electrochemical reactions versus standard chlorine reduction electrode
(Cl2/Cl−), and pO2− represents the negative logarithm of O2− ion
MgCl2 (l) + H2 O(g) = MgO(s) +2HCl(g)
(1)
4HCl(g) + O2 (g) = 2Cl2 (g) +2H2 O(g)
(2)
Because H2O partial pressure was about four times larger than that
of O2 and moisture in salts can also hydrolyze the salt by reaction (1),
forming HCl. Therefore, it can be assumed that O2 was converted
completely into Cl2, resulting that the cover gas above the molten salt
was mainly composed of Ar, HCl and Cl2. In this case O2 in the molten
salt was absent, and correspondingly O2− in the molten salt was close
to zero.
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H. Sun et al.
Fig. 7. SEM/EDS analysis of the precipitated phases at the grain boundaries of various alloys exposed in molten NaCl-KCl-MgCl2 under Ar at 800 °C for 400 h, etched
with 50 vol.% HCl + 50 vol.% alcohol + small amount of H2O2: (a) Hastelloy N, (b) C276; (c) C22; (d) Haynes 230.
chlorine partial pressure. Due to the most negative oxidation potential
of Cr, the element Cr at the alloy/salt interface was preferentially
electrochemically dissolved to form CrCl2 into the salt. Studies by Zahs
et al. [26] also showed that Cr was attacked preferentially in chloridizing atmospheres, followed by Fe and Ni. Based on the above analyses, the corrosion reactions of Ni-based alloys in the molten chloride
salt under the present conditions are inferred as follows:
Anodic reaction:
The Cl2 as the depolarized species was involved in the cathodic
process. The Cl2 partial pressure was calculated to be 3.2 × 10−6 atm at
700 °C according to reaction (2) and Cl− ion activity in the molten
chloride salt was assumed to be 1. According to the Nernst equation, the
equilibrium electrode potential of Cl2 reduction at 700 °C was calculated to be -0.53 V (vs. Cl2/Cl−). In addition, HCl may be dissolved
partially into the melt and corresponding H+ ion could act as a depolarized species to participate in the cathodic reaction. However, the
equilibrium electrode potential of H2 evolution can not be calculated
due to the unknown H+ ion activity and H2 partial pressure. Generally,
the cathodic reaction of corrosion depends on the depolarized species
with the lowest negative electrode potential [28], suggesting that the
electrode potential of cathodic reaction under the present conditions
will be equal or above -0.53 V (vs. Cl2/Cl−). As a result, the corrosion
potential of Ni-based alloys was between the potentials of Cr oxidation
and Cl2 reduction as labelled in Fig. 10 or shifted to more positive
potential direction. It can be seen that Cr and Ni are dissolved to form
CrCl2/CrCl3 and NiCl2 under the present conditions, respectively. Zahs,
Kawahara and Bender et al. [26,27,29] also pointed out that the elements such as Cr, Fe and Ni tended to form their chlorides in oxidizingchloridizing atmospheres with low oxygen partial pressure and high
Cr(s) → Cr 2+ (l) +2e
(3)
Cathodic reaction:
Cl2 (g) +2e → 2Cl− (l)
(4)
and/or
2H+ (l) +2e → H2 (g)
(5)
Overall corrosion reactions:
Cr(s) +Cl2 (g) = CrCl2 (l)
(6)
and/or
Cr(s) +2HCl(l) = CrCl2 (l) + H2 (g)
8
(7)
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H. Sun et al.
Fig. 8. Surface SEM images and corresponding EDS analysis results of Hastelloy N(a,b), C276(c,d), C22(e,f) and Haynes 230(g,h) exposed in molten NaCl-KCl-MgCl2
under Ar for 400 h at 600 °C(a,c,e,g) and 800 °C(b,d,f,h).
4.2. Effects of the alloying elements on the alloy corrosion
The thermodynamic prediction agrees well with the experimental
results, that is, the corrosion of the alloys primarily depends on the
dissolution of Cr to form Cr chloride into the salt, resulting in the formation of voids at the subsurface of the alloys and weight loss. It should
be mentioned that MgO with low solubility in the molten chloride salt
can be formed based on reaction (1), which can deposit into the voids of
the corroded specimens. Accordingly, the alloys showed the abnormal
weight changes with the accelerated corrosion (Fig. 2) and the O enrichment was observed in the Cr depletion regions (Fig. 6).
The present investigation found that the Ni-Mo-Cr alloys showed
better corrosion resistance than the Ni-W-Cr alloy. This result implies
that the element Mo is favorable to enhance the corrosion resistance of
Ni-based alloys in molten chloride salts under Ar gas. As before mentioned, Alkhamis [9] also found Hastelloy C276 with high Mo showed
better corrosion resistance than Haynes 230 with high W in molten
MgCl2-KCl and NaCl-KCl-ZnCl2 at 800 °C. Additionally, Bender et al.
[29] and Galetz et al. [30] also revealed that Mo containing Ni-based
alloys were almost inert against chlorine induced corrosion in reducing
atmospheres. However, a comparison of the Gibbs free energy of
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Corrosion Science xxx (xxxx) xxx–xxx
H. Sun et al.
Fig. 9. XRD patterns of the surfaces of various alloys exposed in molten NaCl-KCl-MgCl2 under Ar for 400 h at 600 °C and 800 °C: (a)Hastelloy N; (b) C276; (c) C22;
(d) Haynes 230.
phases, restricting the outward diffusion of Cr from the alloy base because of the nobility of Mo element (Fig. 11a). In contrast, it is difficult
for W to diffuse and segregate around the voids of the Ni-W-Cr alloy due
to the large atom size of W [31,32], so that the outward diffusion of Cr
can not be prevented during the corrosion process of the Ni-W-Cr alloy
(Fig. 11d).
On the other hand, the chemical compositions of carbides at the
grain boundaries of various alloys are different. According to Fig. 7, the
Ni-Mo-Cr alloys are prone to form the Mo-rich and Ni-rich carbides at
the grain boundaries, while the Ni-W-Cr alloy facilitates the formation
of Cr-rich carbides at the grain boundaries. A comparision of Gibbs free
energy of formation of various divalent metal chlorides at 800 °C (CrCl2
-265.2 kJ/mol; FeCl2 -214.2 kJ/mol; NiCl2 -145.6 kJ/mol; MoCl2
-129.9 kJ/mol; WCl2 -135.7 kJ/mol) indicates that Cr is the most
thermodynamically favored element in chloride salts compared to the
other alloying elements, suggesting that the high Cr contents are very
detrimental for the resistance of carbides against molten chloride salt.
Galetz et al. [30] also pointed out that Mo carbides were much more
resistant than Cr carbides to chlorinating attack. Therefore, the carbides
at the grain boundaries of Haynes 230 can be dissolved during corrosion, leading to severe intergranular corrosion (Fig. 11e and f). In
contrast, the Cr contents of carbides at the grain boundaries of the NiMo-Cr alloys are obviously smaller than that of Haynes 230, and
therefore the Ni-Mo-Cr alloys show better resistance to the intergranular corrosion than Haynes 230 (Fig. 11b and c). Meanwhile, the
carbides with better compatibility with the molten salt can play a better
role in preventing the outward diffusion of Cr from the alloy base,
which can further protect the alloy base from corrosion.
The present investigation also found that the corrosion resistance of
the Ni-Mo-Cr alloys increased with increasing Mo content and decreasing Cr content. This was mainly attributed to the different influence of Mo and Cr on the corrosion, that is, Mo can help to improve the
corrosion resistance of Ni-based alloys, while Cr tends to be dissolved
preferentially in molten chlorides. Additionally, the Cr content of
Fig. 10. Potential-oxoacidity (E-pO2−) diagram calculated for Cr, Ni in molten
NaCl-KCl-MgCl2 at 700 °C.
formation for divalent Mo and W chlorides indicates that Mo and W
show similar resistance against molten chloride salt from the thermodynamic viewpoint (600 °C: MoCl2 -157.9 kJ/mol; WCl2 -155.5 kJ/mol;
700 °C: MoCl2 -143.8 kJ/mol; WCl2 -145.5 kJ/mol; 800 °C: MoCl2
-129.9 kJ/mol; WCl2 -135.7 kJ/mol). Based on the present experimental
results, the different influence of Mo and W on the corrosion resistance
of Ni-based alloys mainly depends on two reasons, as shown schematically in Fig. 11.
On the one hand, the elements Mo and W show different ability in
restricting the outward diffusion of Cr. According to the EPMA, SEM/
EDS and XRD results (Figs. 4, 5, 8 and 9), Mo can diffuse from the alloy
base to the surface during the corrosion process of the Ni-Mo-Cr alloys
and be then segregated around the voids as a form of new Mo-rich
10
Corrosion Science xxx (xxxx) xxx–xxx
H. Sun et al.
Fig. 11. Schematic of corrosion of the Ni-based alloys in molten NaCl-KCl-MgCl2 under Ar gas:(a,d) the early stage of corrosion; (b,e) the later stage; (c,f) observed
situation.
but also along line and surface defects such as grain boundaries.
According to the literature [33,34], the lattice and grain boundary
diffusion coefficients of Cr in Ni-based alloys are expressed as respectively:
carbides increases with Cr content of the alloy bases (Fig. 7), which is
also responsible for worse corrosion resistance of the Ni-Mo-Cr alloys
with higher Cr content.
4.3. Effects of temperature on the alloy corrosion
-1
340kJmol ⎞ 2 -1
Dv = 1.8 × 10−2 exp ⎛⎜−
⎟m s
RT
⎝
⎠
The experimental results found that the corrosion depth of the alloys
increased and the intergranular corrosion became more outstanding
with temperature. Futhermore, the effects of the alloying elements on
the corrosion of Ni-based alloys have a strong temperature dependence.
The corrosion rate of alloy is mainly controlled by the outward diffusion
of Cr from the alloy base to the surface. The diffusion distance of Cr in
Ni could be estimated by assuming [33]
x = (Dt )1/2
(9)
-1
335kJmol ⎞ 3 -1
δDgb = 8.2 × 10−8 exp ⎜⎛−
⎟m s
RT
⎝
⎠
(10)
where Dv is lattice (volume) diffusion coefficient, Dgb is grain boundary
diffusion coefficient, δ is grain boundary width (assume δ = 0.5 nm), R
is the gas constant and T is the temperature. Based on Eqs. (9) and (10),
Dv and Dgb of Cr in Ni-based alloys at temperatures investigated can be
calculated and are shown in Fig. 12. It is clear that Dv and Dgb increase
with temperature, resulting that the Cr diffusion depths increase. Accordingly, the corrosion of the alloys is accelerated with temperature.
(8)
where x is the diffusion distance, D the diffusion coefficient and t the
immersion time. At the same immersion time, x is mainly related to D.
The solid diffusion occurs not only along point (or lattice) defects,
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Corrosion Science xxx (xxxx) xxx–xxx
H. Sun et al.
carbides are precipitated at the grain boundaries of the alloys with
increasing temperature, which also accelerates the intergranular
corrosion and enlarges the differences in the corrosion resistance of
various alloys.
Acknowledgement
This work was supported by the Strategically Leading Program of
the Chinese Academy of Sciences (Grant No. XDA02040000).
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5. Conclusions
(1) Hastelloy N, C276, C22 and Haynes 230 suffer corrosion in molten
NaCl-KCl-MgCl2 under Ar atmosphere in the temperature range of
600–800 °C for 400 h. The corrosion of the alloys is mainly attributed to the selective dissolution of Cr to form soluble Cr chlorides
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(2) The Ni-Mo-Cr alloys exhibit better corrosion resistance than the NiW-Cr alloy and the corrosion resistance of the Ni-Mo-Cr alloys increases with an increase of Mo and a decrease of Cr. These differences are mainly caused by the different reactivity of formation of
chlorides and the different roles of the alloying elements in preventing the outward diffusion of Cr from the alloy base, as well as
the alloying elements induced differences in the chemical compositions of carbides at the grain boundaries.
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corrosion resistance of Ni-based alloys have a strong temperature
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