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ORIGINAL PAPER
Na2SO4-Deposit-Induced Corrosion of Mo-ContainingAlloys
B. S. Lutz1,2• J. M. Alvarado-Orozco1,3
•
L. Garcia-Fresnillo1,4• G. H. Meier1
Received: 2 August 2016 / Revised: 23 January 2017 / Published online: 6 February 2017
� Springer Science+Business Media New York 2017
Abstract Disk alloys used in advanced gas turbine engines often contain significant
amounts of Mo (2 wt% or greater), which is known to cause corrosion under Type I
hot corrosion conditions (at temperatures around 900 �C) due to alloy-induced
acidic fluxing. The corrosion resistance of several model and commercial Ni-based
disk alloys with different amounts of Mo with and without Na2SO4 deposit was
examined at 700 �C in air and in SO2-containing atmospheres. When coated with
Na2SO4 those alloys with 2 wt% or more Mo showed degradation products similar
to those observed previously in Mo-containing alloys, which undergo alloy-induced
acidic fluxing Type I hot corrosion even though the temperatures used in the present
study were in the Type II hot corrosion range. Extensive degradation was observed
even after exposure in air. The reason for the observed degradation is the formation
of sodium molybdate. Transient molybdenum oxide reacts with the sodium sulfate
deposit to form sodium molybdate which is molten at the temperature of study, i.e.,
700 �C, and results in a highly acidic melt at the salt alloy interface. This provides a
negative solubility gradient for the oxides of the alloying elements, which results in
continuous fluxing of otherwise protective oxides.
Keywords Hot corrosion � Disk alloys � Mo-effects � Alloy-induced acidic fluxing
& G. H. Meier
1 Mechanical and Materials Engineering Department, University of Pittsburgh, Benedum
Engineering Hall, Pittsburgh, PA 15261, USA
2 Present Address: University of California, Santa Barbara, Santa Barbara, CA, USA
3 Present Address: Centro de Ingenierıa y Desarrollo Industrial, Queretaro Campus,
76130 Queretaro, Mexico
4 Present Address: AUTOVISION Servicios, Autovıa A-2, Km 585, 08760 Martorell, Spain
123
Oxid Met (2017) 88:599–620
DOI 10.1007/s11085-017-9746-0
Introduction
Advanced gas turbine engines are currently being developed for reduced gas
emissions and reduced fuel consumption (higher efficiency). This can be done by
increasing the combustion temperature and pressure or employing improvements in
cooling technologies. Not only will the temperatures and pressures of the turbine
blades be increased, but also the temperatures and pressures surrounding the disk
rotors will increase. Currently, high-strength nickel alloys are used for disk rotor
hardware because of their high-temperature oxidation resistance and strength.
Exposing disk alloys to engine environments for long periods of time will cause
severe oxidation and hot corrosion due to salt deposits. Increasing the temperature
and pressure will have even greater effects on disk integrity [1]. Many of the disk
alloys contain significant amounts of Mo, which can cause significant corrosion at
high temperatures (900 �C), and as will be seen, even at low temperatures (700 �C).
Sodium Sulfate-Induced Hot Corrosion
Alloys used in the combustion process of gas turbine engines, especially those used
in marine applications, can undergo an aggressive form of corrosion associated with
the formation of a salt deposit, which is usually based on sodium sulfate, on the
surface of the metal or thermally grown oxide. This type of corrosion is called hot
corrosion. The amount of hot corrosion caused is determined by the amount of
Na2SO4 deposit, gas atmosphere, temperature, cycling, erosion, and alloy compo-
sition [2]. The composition of the deposit is also important and can be influenced by
impurities coming either from the alloy or the environment.
The hot corrosion process occurs in two stages: initiation stage and propagation
stage. During the initiation stage, the alloy undergoes a process similar to simple
oxidation with the difference that in the hot corrosion process the oxidizing species
comes from the salt deposit. The interaction between the oxides formed on the alloy
and the salt can lead to extremely corrosive conditions as the salt becomes basic or
acidic. Na2SO4 exhibits an acid–base character with SO3(g) being the acidic
component and Na2O(s) the basic component. The composition of the Na2SO4 melt
at a fixed temperature can therefore be described by the oxygen partial pressure and
the activity of Na2O in the melt, aNa2O, or the SO3 partial pressure, because it is
related to the activity of Na2O by the equilibrium constant for Eq. 1 [3].
Na2SO4 ¼ SO3 þ Na2O ð1Þ
The salt may become more basic (higher Na2O activity) or more acidic (lower
Na2O activity) when interacting with the base alloy and its oxides, i.e., species
dissolving into the salt can alter its basicity. A phase stability diagram for the Na–S–
O system (Fig. 1) shows the compositional changes which occur in the salt deposit
during the initiation stage. Once the salt deposit becomes sufficiently basic or acidic
rapid degradation can occur, i.e., the propagation stage starts. In many cases, the end
of the initiation stage occurs when the deposit becomes liquid, penetrates the oxide
scale and spreads along the alloy/scale interface.
600 Oxid Met (2017) 88:599–620
123
The propagation stage is dependent on what happens to the alloy and the oxides
during the initiation stage. Rapp [4] established solubility curves for a number of
oxides in Na2SO4 as a function of the activity of Na2O in the melt (Fig. 2). Rapp
and Goto [5] have proposed a criterion by which the continued self-sustaining hot
corrosion attack can occur. This is known as the Rapp–Goto criterion, and it is given
by Eq. 2 shown below, where Coxide is the solubility of the protective oxide and x is
the distance into the molten salt deposit from the oxide/salt interface.
Fig. 1 Thermodynamic stability diagram for the Na–O–S system at constant temperature [2]
4 6 8 10 12 14 16 181.0
1.5
2.0
2.5
3.0
3.5
4.0
SiO2 (900°C)
Al2O3
Cr2O3
Fe2O3
NiO
Log
Con
c. p
pm (m
ole
met
al Io
ns/m
ole
Na 2S
O 4)
- Log aNa2O
Co3O4
Fig. 2 Measured oxide solubilities in Na2SO4 at 927 �C and 1 atm O2. Adapted from [4]
Oxid Met (2017) 88:599–620 601
123
dCoxide
dx
� �x¼0
\0 ð2Þ
When the solubility gradient is positive, the salt can become saturated with oxide
and a protective scale is able to form over the metal surface. A negative gradient in
the oxide solubility at the oxide/scale interface results in dissolution of the
protective oxide scale and reprecipitation of the oxide as discontinuous non-
protective particles in regions of the molten salt deposit where the solubility is
lower.
Once a deposit has been formed on the surface of combustion hardware, the
amount of corrosion depends significantly on whether or not the deposit melts. If the
temperature of the combustion environment is above the melting point of Na2SO4
(Tm = 884 �C) [6], the corrosion is called Type I hot corrosion or high-temperature
hot corrosion. Below the melting point of Na2SO4, the salt deposit can become
molten because of a reaction between the combustion gas and the oxide scale grown
on the alloy. This type of corrosion is called Type II hot corrosion or low-
temperature hot corrosion.
Type I hot corrosion is typically assumed to occur at high temperatures of
approximately 900–1000 �C. There are two common propagation modes for this
type of hot corrosion. These occur when there is dissolution of the protective oxide
into a highly basic molten salt deposit where aNa2O is high, and when there is
dissolution of the protective oxide into a highly acidic molten salt deposit where
aNa2O is low. These are called basic fluxing and alloy-induced acidic fluxing,
respectively. These two forms of Type I hot corrosion occur due to interactions
between the salt deposit and the underlying alloy substrate and its oxides. A
description of the corrosion mechanism for alloy-induced acidic fluxing Type I hot
corrosion will now be discussed in more detail since some of the alloys tested
showed characteristics of alloy-induced acidic fluxing even at low (700 �C)temperatures.
The oxide solubility plots presented in Fig. 2 show that a protective oxide scale
can be dissolved due to liquid sulfate deposits with low aNa2O (acidic melts). The
Na2SO4 deposit can become acidic due to SO3 in the gas atmosphere, or by the
dissolution of transient oxides of Mo, W, or V, which are added to superalloys for
solid solution strengthening. Alloy-induced acidic fluxing is due to the latter.
Bornstein et al. [7] studied the effects of various elemental additions to nickel-based
alloys on the Na2SO4-induced hot corrosion at temperatures between 800 and
1000 �C. The oxidation rates of alumina forming alloys were significantly
accelerated when there were additions of Mo or V in the alloys or when Na2SO4
was deposited with MoO3 or V2O5. They concluded that the attack was due to the
fact that the acidic transient MoO3 and V2O5 oxides were molten at the temperatures
tested. The molten oxides flux the Al2O3 scale, which causes rapid degradation of
the substrate. Goebel and Pettit [8] studied the oxidation of some commercial and
model nickel-based alloys with different elemental additions in order to determine
their effect on Na2SO4-induced hot corrosion. They oxidized Ni–Al and Ni–Cr–Al
alloys with additions of Mo, W, and V at 1000 �C in air with Na2SO4 deposits. All
602 Oxid Met (2017) 88:599–620
123
of the alloys tested were catastrophically degraded. The refractory element additions
were found to be concentrated at the alloy/scale interface, and the attack was
initiated by these phases. The authors also performed tests in which oxides of Mo,
W, V, and Cr were mixed with Na2SO4 in an alumina crucible and heated to
1000 �C. The crucibles with WO3, MoO3, and V2O5 all lost weight and traces of Al
were found in the salt after the test. The refractory oxides lower the Na2O activity
significantly so that the reaction in Eq. 3 can occur.
Al2O3 þ 3MoO3 ¼ Al2 MoO4ð Þ3 ð3Þ
Based on these tests the authors proposed a corrosion mechanism for alumina
forming nickel-based alloys, which contain refractory elements [8]. During the first
stages of oxidation, transient MoO3 and NiO oxides as well as Al2O3 form at the
alloy surface. MoO3 reacts with the oxide ions in the Na2SO4 salt deposit to
decrease the Na2O activity of the melt, which prevents basic fluxing mechanism
from occurring. This reaction is shown below in Eq. 4.
MoO3 þ Na2O ¼ Na2MoO4 ð4Þ
MoO3 will decrease the Na2O activity of Na2SO4 to levels where Eq. 3 and
acidic fluxing of the protective Al2O3 scale can occur. The attack initiates near Mo-
rich particles because the activity of MoO3 in the Na2SO4 is highest in these areas.
The fluxing of the Al2O3 scale causes Al3? and MoO4
2- ions to diffuse through the
salt where Al2O3 reprecipitates as a porous non-protective scale, and the MoO3
evaporates. This form of attack is self-sustaining, because MoO3 is able to continue
to form at the alloy/salt interface and evaporate at the salt/gas interface [8]. This
maintains a negative solubility gradient for the oxides which dissolve and
reprecipitate. Fryburg et al. [9] performed chemical analysis of the corrosion
products from Mo-containing alloys and found that rapid corrosion occurred after
all of the Na2SO4 was converted to Na2MoO4.
The previous section described alloy-induced acidic fluxing at high temperatures
of around (900–1000 �C). It is also possible to have hot corrosion at lower
temperatures between 650 and 750 �C, and the attack can be significant at these
lower temperatures. Pure Na2SO4 has a melting point of 884 �C, so, unless it
contains other sulfates as impurities which lower its melting point, it does not form a
molten salt and remains solid at these temperatures. In order for severe degradation
to occur, the salt must become liquid, and so severe corrosion would not be expected
at these temperatures, but often occurs in turbine components. This type of low-
temperature corrosion is called Type II hot corrosion. Type II hot corrosion is not
completely understood, but the accelerated corrosion is generally believed to be
caused by the formation of a Na2SO4–MSO4 eutectic melt (e.g., M = Ni, Co) that
has a melting point well below that of Na2SO4 [10]. It is also known that a partial
pressure of SO3 of about 10-5 atm is required for this melt to form [11]. This
amount of SO3 in the gas atmosphere is not uncommon in industrial gas turbines.
The research presented in this paper highlights the importance of Mo content for
the hot corrosion resistance of the alloy. It studies the low-temperature oxidation
(700 �C) as well as Na2SO4-induced hot corrosion morphologies and mechanisms of
Oxid Met (2017) 88:599–620 603
123
several model and commercial Ni-based disk alloys which contain different
amounts of Mo. The degradation mechanism is compared to the Type I alloy-
induced acidic fluxing hot corrosion mechanism described previously.
Experimental Procedures
Three commercial nickel-based superalloys (IN-738, IN-617 and RR1000) and two
model NiCrAl alloys were initially tested. The alloys were of differing types. IN-
738 is a chromia-forming c–c0 alloy. IN-617 is a chromia-forming c-solid solution
alloy. RR1000 is a modern disk alloy whose equilibrium microstructure varies
significantly with temperature. The major phases over the temperature range of the
tests conducted are c and c0, and the minor phases are the carbides M23C6 and MC
[12]. RR1000 is reported to form chromia scales [1] although the composition
suggests it could form alumina under some conditions. The NiCrAl and NiCrAl Mo
model alloys are c–c0 alloys and are borderline alumina formers. The Mo content in
the alloys should be noted, as it is the focus of this study. IN-738 is a commercial
Ni-based superalloy with relatively low Mo content (1.75 wt%), IN-617 has a
significant amount of Mo (9 wt%), while the disk alloy RR1000 has 5 wt% Mo. The
model alloys were a Ni–6Cr–8Al alloy and an alloy with the same chemical
composition, but with 6 wt% Mo. After the initial experimental results, three
additional model alloys with composition similar to that of RR1000 but varying
amounts of molybdenum were tested. The equilibrium microstructures of these
alloys are similar to RR1000, but there are no carbides. The aim of these tests was to
do a more systematic study and get further insight into the role that molybdenum
plays in the hot corrosion resistance of Ni-base alloys. The commercial alloys
together with the NiCrAl(Mo) model alloys will in the following be designated as
disk alloys and the model alloys, which contain 20% Co, with varying amount of
Mo will be designated as Mo-containing model alloys (second set). The
compositions of all alloys tested are shown in Table 1.
Table 1 Nominal alloy compositions (wt%)
Ni Co Cr Al Mo Other
Disk alloys
IN-738 Balance 8.5 16 3.5 1.75 3.5Ti, 2.5W, 1.75Ta, 0.8 Nb
IN-617 Balance 12 22 1.2 9 0.3Ti
RR1000 Balance 18.5 15 3 5 3.6Ti, 2Ta, 0.5Hf, 0.03C
NiCrAl Balance – 8 6 – –
NiCrAlMo Balance – 8 6 6 –
Mo-containing model alloys (second group)
Ni–20Co–15Cr–3.5Al–2Mo Balance 20 15 3.5 2 –
Ni–20Co–15Cr–3.5Al–4Mo Balance 20 15 3.5 4 –
Ni–20Co–15Cr–3.5Al–6Mo Balance 20 15 3.5 6 –
604 Oxid Met (2017) 88:599–620
123
All commercial alloys were cut into rectangular coupon specimens of approx-
imately 15 9 10 9 2 mm3. The model NiCrAl and NiCrAlMo alloys were received
as cast bars that were heat treated at 1000 �C for 24 h in sealed quartz tubes under a
small partial pressure of argon. The cast bars were cut into circular coupons of
approximately 1.5 mm thick and 18 mm in diameter in the case of the NiCrAl
model alloy and 1.5 mm thick and 19.4 mm in diameter for the NiCrAlMo model
alloy. The Mo-containing model alloys (second set) were received as cast bars of
approximately 10 mm in diameter, which had been annealed for 24 h at 1000 �C in
vacuum. These were cut into circular coupons approximately 0.6 mm thick. All of
the specimens were ground with SiC sandpaper to a 1200 grit finish and
subsequently cleaned and degreased in ethanol using an ultrasonic cleaner. The
specimens were then dried and weighed before any deposits were applied or testing
was conducted.
Na2SO4 was the salt deposit used. The deposit was prepared by mixing 80 g of
Na2SO4 powder per 16 oz. (473 ml) of distilled water. A 2.75 ± 0.25 mg/cm2
deposit was applied to the surface of the specimens by heating the specimens with a
heat gun and spraying them with the prepared Na2SO4 solution. During the process,
the specimens were weighed to ensure that the desired amount was deposited on the
surface.
All experiments were conducted in horizontal resistance-heated furnaces whose
reaction tube was either made of sintered alumina or silica glass with an internal
sintered alumina sleeve. The hot zone was maintained within 3 degrees of the test
temperature. The gas atmospheres employed during testing consisted either of
laboratory air or oxygen with 1000 ppm SO2. For the exposures in laboratory air,
the specimens were placed in alumina boats which were laterally pushed into the hot
zone. In the case of the O2 ? 1000 ppm SO2 exposures, a slightly more complex
setup, shown schematically in Fig. 3, was used. The specimens were placed in a
quartz rod that could be cycled in and out of the hot zone of the furnace manually by
using a magnet. The O2 ? 1000 ppm SO2 gas flowed into the tube at a constant
flow rate of 15 ml/min (0.0125 cm/s) and passed over a platinum honeycomb
catalyst placed in the hot zone of the furnace to establish the equilibrium pSO3.
Assuming that equilibrium was attained, the equilibrium pSO3values for the
temperatures and gas atmospheres tested are those given in Table 2.
When exiting the furnace, the gas was bubbled through a Na2CO3 plus water
mixture before entering the fume hood. This removes the SO3 from the gas, as
Fig. 3 Schematic of horizontal tube furnace apparatus for fireside corrosion tests
Oxid Met (2017) 88:599–620 605
123
shown by the reaction in Eq. 5. The sodium sulfate product precipitates out in the
bubbler, and the exiting gas was predominantly carbon dioxide.
SO3 þ Na2CO3 ¼ Na2SO4 þ CO2 ð5Þ
After exposure, the surface of selected specimens was investigated by photo-
stimulated luminescence spectroscopy (PSLS) using a micro-Raman mapping
spectrometer (Renishaw InVia) connected to a Leica microscope equipped with a
633 nm line-focus laser. Details of the PSLS technique can be found in [13]. All
specimens were prepared for metallographic analysis. Some of the specimens were
PVD coated with a 20-nm-thick layer of palladium and subsequently electrochem-
ically coated with a 100-lm-thick layer of nickel using a Watt’s bath [14]. These
coatings were applied in order to support the oxide scale during further specimen
preparation and to improve the contrast between the scale and the mounting material
during analysis. All specimens were embedded in an epoxy resin and cross sections
were prepared using conventional techniques, i.e., grinding, coarse polishing, and
fine polishing. In the case of the specimens with Na2SO4 deposits, oil-based
polishing solutions were used in order to preserve the corrosion products. The
metallographic cross sections were characterized using scanning electron micro-
scopy (SEM) with energy dispersive X-ray spectroscopy (EDS). Limited charac-
terization was also performed by wavelength dispersive X-ray spectroscopy (WDS).
Results
Disk Alloys
Exposure in Air Without Na2SO4 Deposit
Before any tests were conducted on the selected alloys with Na2SO4 deposits, each
of the disk alloys was isothermally oxidized in air at 700 �C for 100 h without any
deposit. This was used as a baseline for comparison, so that the effect of the salt
deposit could be determined.
All commercial alloys developed an extremely thin protective oxide as can be
seen in cross-sectional views in Fig. 4. IN738 formed a very thin oxide scale
consisting of NiO, Al2O3, and Cr2O3. The white areas in the substrate of the
specimen are niobium carbide (NbC). Both IN-617 and RR1000 formed very thin
oxide scales consisting of Cr2O3 and transient NiO on top. The results for RR1000
Table 2 Equilibrium SO3 partial pressures at experimental temperatures and gas atmospheres
Temperature (�C) O2 ? 1000 ppm SO2 (atm)
650 7.2 9 10-4
700 4.5 9 10-4
606 Oxid Met (2017) 88:599–620
123
are comparable to those obtained by Encinas-Oropesa et al. [1]. The scale thickness,
which was estimated from TGA results in Ref. [1], after 100 h at 700 �C, was0.5 lm. That observed in Fig. 4c is of the same magnitude. The NiCrAl model alloy
formed a slightly thicker oxide scale than the commercial alloys which consisted of
transient NiO on top of an Al2O3 scale. The scale formed on the NiCrAlMo model
alloy was considerably thicker. Oxide nodules of NiO formed with internal Al2O3
underneath. The Mo in the alloy is apparently retarding the development of
continuous alumina; see Fig. 4d, e. The oxidation microstructures shown in Fig. 4
will be used for comparison with the tests in Na2SO4.
Exposure in Air with Na2SO4 Deposit
After exposure of the alloys in air without a Na2SO4 deposit, other disk alloys were
exposed under similar conditions, i.e., isothermally exposed at 700 �C in air for
100 h, with a Na2SO4 deposit.
The results presented in Fig. 5 show NiCrAl and IN-738 (Fig. 5a, d) did not
undergo any substantial attack (the Na2SO4 remained solid). They grew very thin
oxide scales similar to the those observed for the exposures without the deposit.
There were some small localized areas of internal oxidation and sulfide penetration,
but not to the level of the other alloys tested. These are the two alloys that contained
either no or minimal Mo. In air, without the presence of SO3, a liquid NiSO4–
Na2SO4 melt is not able to form and cause corrosion. The Mo-containing disk alloys
(NiCrAlMo, IN-617 and RR1000) experienced more degradation. NiCrAlMo
suffered the most severe corrosion as can be seen in Fig. 5e. A NiO scale grew on
the outside over the salt deposit. A porous internal oxide scale consisting of NiO,
Cr2O3, and Al2O3 grew with some chromium sulfides penetrating into the substrate.
EDS measurements also revealed the presence of some Mo-oxides in the corrosion
pit at the metal/oxide interface. IN-617 and RR1000 (Fig. 5b, c) showed similar
microstructures, but the degradation was less severe than that observed for the
NiCrAlMo alloy. A NiO outer layer formed by fluxing of Ni through the deposit and
Fig. 4 SEM images of disk alloys without Na2SO4 deposit after exposure in air at 700 �C for 100 h
Oxid Met (2017) 88:599–620 607
123
covered an internal scale of Cr2O3 and Al2O3. Traces of Mo-oxides could also be
seen in the internal oxide scale. The NiO was fluxed and reprecipitated in the melt
near the gas/salt interface and internal pits which were rich in Cr, Al, Mo, S, and O
formed in the alloy. This microstructure is similar to the one, which might be seen
with alloy-induced acidic fluxing Type I hot corrosion, which was described earlier.
Given enough time, it is believed that each of these three Mo-containing alloys
would suffer severe degradation, and the microstructures would look exactly like
that of a specimen that has undergone alloy-induced acidic fluxing, i.e., NiCrAlMo
alloy. The temperatures used here are in the Type II hot corrosion range. However,
the presence of Mo is causing these alloys to corrode at 700 �C in air even without
the presence of SO3. A liquid is generally needed for severe degradation to occur.
The transient MoO3 reacts with the salt to form sodium molybdate, releasing SO3.
This is shown in the reaction given in Eq. 6.
MoO3 þ Na2SO4 ¼ Na2MoO4 þ SO3 ð6Þ
The melting point of sodium molybdate is 687 �C, so at the temperature tested
(700 �C) it becomes molten. The molten Na2MoO4 is able to dissolve the protective
oxide scale and cause corrosion of the alloy substrate as described by Fryburg et al.
[9] for higher temperatures.
To confirm that the formation of liquid sodium molybdate caused the corrosion
of the alloys at 700 �C, a test was conducted on the three alloys with higher Mo
contents at 650 �C. This temperature is below the melting point of sodium
molybdate (687 �C), and thus a liquid should not form and cause severe corrosion.
This was indeed the case, as can be seen from the results of this test shown in Fig. 6.
Each of the alloys grew a thin oxide scale consisting predominantly of Cr2O3 and
transient NiO. There were some small localized areas of internal Cr2O3 and Al2O3
growth, but fluxing and extensive corrosion did not occur as it did at 700 �C. Thepresence of Ni and Al was also detected in the Cr-rich scale on RR1000.
Fig. 5 SEM images of disk alloys with Na2SO4 deposit after exposure in air at 700 �C for 100 h. Somescale spalled on cooling from IN-738 and NiCrAl
608 Oxid Met (2017) 88:599–620
123
The model NiCrAl and NiCrAlMo alloys were also exposed at 800 and 900 �Cfor 100 h in air with Na2SO4 deposited on the surface to examine the effect of
temperature on the amount of degradation. At these temperatures, the NiCrAl alloy
should remain fairly protective, although the amount of oxidation should increase
with temperature. The NiCrAlMo alloy should be more extensively degraded,
because not only should Na2MoO4 become molten, but MoO3 has a melting point of
795 �C, so it should become molten, even in the absence of Na, and add to the
degradation. The results for each of the NiCrAl and NiCrAlMo alloys at 800 and
900 �C are presented in Fig. 7.
The NiCrAl alloy maintained a fairly protective scale rich in Ni, Al, and Cr
oxides. As the temperature increased, the amount of oxidation increased and some
internal oxidation occurred. At 900 �C, there were some localized areas across the
specimen with severe degradation, which is the result of basic fluxing. The
Fig. 6 SEM images of Mo-containing disk alloys with Na2SO4 deposit after exposure in air at 650 �C for100 h
Fig. 7 SEM images of Ni–8Cr–6Al and Ni–8Cr–6Al–6Mo alloys deposited with Na2SO4 exposed in airat 800 �C (a, b) and 900 �C (c, d) for 100 h
Oxid Met (2017) 88:599–620 609
123
NiCrAlMo alloy was catastrophically degraded at both 800 and 900 �C, and
significant spallation was observed upon cooling the specimens. The scales which
did remain were thick Ni oxide rich scales with remnant Na2SO4, as well as MoO3
mixed in. There were internal pits rich in Ni, Mo, Al, Cr, and S, and the sulfur
content increased with depth into the pits forming a sulfide rich layer at the base
along with internal sulfidation. The extent of degradation increased monotonically
as the temperature was increased.
Exposure in O2 ? 1000 ppm SO2 with Na2SO4 Deposit
Preliminary experiments were conducted in which the disk alloys were exposed
under hot corrosion conditions with SO3 added to the oxidizing gas (700 �C for
100 h in O2 ? 1000 ppm SO2). The results can be seen in Fig. 8. The attack on
each of the alloys tested in the SO2-containing atmosphere (Fig. 8) was greater than
that observed when the alloys were tested in air (Fig. 5). The NiCrAl alloy
developed a thick NiO scale which grew over top of internal oxide pits which are
rich in Ni, Al, Cr, O, and S. The sulfur content increased deeper into the pits and
chromium sulfides could be seen penetrating into the substrate. This is typical Type
II hot corrosion attack in which SO3 causes the formation of a liquid Na2SO4–MSO4
(M = Ni and/or Co) melt. IN-738, NiCrAlMo, IN-617, and RR1000 suffered severe
spallation upon cooling of the specimens. EDS analysis showed that at those points
where there was severe spallation, a layer rich in Al, S, and O remained. The
spallation on the IN-738 alloy was so severe that hardly any intact oxides or full
corrosion product could be found. It should be noted that IN-738 contains W, in
addition to Mo, which also forms an acidic oxide, WO3. Where there was no
spallation on the NiCrAlMo alloy a thick external NiO oxide scale grew with
remnant Na2SO4 deposit intermixed into the scale. This turned into internal pits rich
in Al, Cr, S, and O. The base of the pits was identical to the sulfide layer found
Fig. 8 SEM images of disk alloys with Na2SO4 deposit under Type II hot corrosion conditions, i.e.,exposure in O2 ? 1000 ppm SO2 at 700 �C for 100 h
610 Oxid Met (2017) 88:599–620
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where spallation occurred. IN-617 suffered attack similar to the NiCrAlMo alloy,
but the degradation was greater. Where the oxide had not spalled away, a thick NiO
oxide scale grew with circular areas of remnant Na2SO4 deposit intermixed into the
scale. This turned into deeper internal pits rich in Cr, O, and S than in the case of the
other alloys. Some circular Na2SO4 deposit could also be seen in the internal pits as
well. A lighter region that contained large amounts of Cr and S was observed at the
base of the pits. RR1000 had an attack, which was similar to IN-617 and NiCrAlMo.
External NiO grew over internal pits rich in Cr2O3 and S. There was some fluxing of
CoO and NiO through remnant parts of the sodium sulfate deposit forming an oxide
layer on the outside of the deposit. This is similar to what was seen in the absence of
SO2 as was described earlier. The presence of Mo at the base of the scale could not
be confirmed by EDS because of overlap between the peaks for Mo and S.
Mo-Containing Model Alloys (Second Set)
The commercial and model disk alloys tested have different base chemical
compositions as well as different Mo contents. In order to get further insight into the
role of molybdenum on the degradation of Mo-containing disk alloys, a series of
model alloys with a base composition of Ni–20Co–15Cr–3.5Al (similar to RR1000)
with additions of 2, 4, and 6 wt% Mo was tested.
Exposure in Air Without Na2SO4 Deposit
As with the disk alloys, before any tests were conducted on the Mo-containing
model alloys with Na2SO4 deposits, the alloys were isothermally oxidized in air at
700 �C for 100 h with no deposit. The results from the alloys without deposit were
used to compare with those when a Na2SO4 deposit was applied, with the aim of
determining the effect of the salt on the degradation of the alloys.
Fig. 9 Weight change values of Mo-containing model alloys with and without Na2SO4 deposit after 100h exposure in air at 700 �C
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Figure 9 shows weight change values of the Mo-containing model alloys without
Na2SO4 deposit measured after exposure in air at 700 �C for 100 h. Weight change
values increased with increasing Mo content of the alloy. The weight gain for
RR1000 (not shown) was approximately 0.02 mg/cm2. Thus, the weight gains for
the model alloys, at a comparable Mo content, appear to be somewhat larger
although still quite small.
After exposure, a thin oxide layer formed on the top of the specimens and the as-
cast structure could clearly be seen in the macrophotographs (Fig. 10). Photo-
stimulated luminescence spectroscopy (PSLS) analyses were carried out on the
surface of the specimens (Fig. 10). Several measurements were made in order to
cover different alloy surface regions. The PSLS spectra showed only peaks
corresponding to a-Al2O3, but the intensity of the peaks depended on the surface
region considered. Thus, the scales consisted of Al2O3 in addition to Cr2O3 and NiO
which were detected by EDS.
SEM images of the Mo-containing alloys after exposure in air at 700 �C for
100 h (Fig. 11) revealed that the Ni–20Co–15Cr–3.5Al–2Mo alloy formed a thin
mixed oxide scale containing Ni, Cr, and Al. The PSLS measurements indicated the
presence of alumina. Nonetheless, this alumina layer must be very thin since it
could not be distinguished by EDS measurements. A similar oxide scale formed on
the Ni–20Co–15Cr–3.5Al–4Mo alloy, but some oxide nodules were present along
the oxide layer. These oxide nodules consisted of two different zones, an upper
oxide rich in Cr, Co, and Ni and a Mo-containing oxide below (6 wt% Mo was
Fig. 10 Macrophotographs and PSLS spectra of Mo-containing model alloys after exposure in air at700 �C for 100 h
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found in the oxide scale vs the 4 wt% contained in the bulk alloy). The Ni–20Co–
15Cr–3.5Al–6Mo alloy formed a thicker oxide scale than the other model alloys.
The oxide scale was a Ni-, Cr- Al-rich oxide scale containing up to 12 wt% Mo.
NiO was found on top of the mixed oxide scale.
Exposure in Air with Na2SO4 Deposit
Even though weight change values are not the most accurate measure of alloy
degradation when a deposit is applied (since phenomena such as evaporation or
deposit spallation must be considered) the weight change values of the Mo-
containing model alloys with Na2SO4 deposit measured after 100 h exposure in air
at 700 �C indicated that the weight change, and thus the degradation of the alloy,
increased with the Mo content of the alloy (Fig. 9). The corresponding weight
change for RR1000 (not shown) was approximately 0.2 mg/cm2, which is
comparable to the model alloy with 4 wt% Mo.
Figure 12 shows SEM images of the Mo-containing model alloys with Na2SO4
deposit after exposure in air at 700 �C for 100 h. The Ni–20Co–15Cr–3.5Al–2Mo
alloy did not undergo substantial attack. Nonetheless, a thin layer of fine Ni-, Cr-
rich oxide particles formed on top of the deposit layer as well as what appears to be
the edges of the Na2SO4 particles. With increasing Mo content of the alloy, more
degradation was observed. Both Ni–20Co–15Cr–3.5Al–4Mo and Ni–20Co–15Cr–
3.5Al–6Mo alloys formed a NiO layer on top of the deposit as well as surrounding
some of the Na2SO4 deposit particles. An internal oxide scale containing Cr, Al, and
Ni formed below the Na2SO4 deposit layer. Below this oxide scale, a Mo-containing
oxide with the presence of S was detected. The presence of Mo and S in the same
region was confirmed by limited WDS measurements. This latter oxide contained up
to 6 and 17.7 wt% Mo on the Ni–20Co–15Cr–3.5Al–4Mo and the Ni–20Co–15Cr–
Fig. 11 SEM images of Mo-containing alloys without Na2SO4 deposit after exposure in air at 700 �C for100 h
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3.5Al–6Mo alloy, respectively. No Mo-enrichment was found at the deposit/alloy
interface on the Ni–20Co–15Cr–3.5Al–2Mo alloy.
Comparing the results of the Mo-containing model alloys with those of the disk
alloys previously presented it is clear that the presence of a certain amount of Mo in
the alloy, more than 2 wt%, is responsible for the observed alloy degradation. The
degradation is similar to alloy-induced acidic fluxing Type I hot corrosion, but it
occurred at lower temperature, i.e., 700 �C, and without the presence of SO2/SO3 in
the atmosphere. The microstructure exhibits ‘‘veins’’ of NiO threading through the
deposit and connected to the NiO outer layer. The deposit adjacent to these veins
was essentially pure Na2SO4, and was presumably solid at the exposure
temperature. The NiO outer layer is similar to that reported by Misra [15] for
U-700 (Ni–14.3Cr–15Co–4.3Al–3.6Ti–4.1Mo wt%) at 950 �C, at which tempera-
ture the deposit would have been entirely molten. The degradation increased as the
Mo content of the alloy increased, but other alloying elements also affected the
behavior of the alloy. Both the NiCrAlMo model alloy and the Ni–20Co–15Cr–
3.5Al–6Mo model alloy contain 6 wt% Mo. Nonetheless, the degradation of the
NiCrAlMo alloy (Fig. 8) was greater than that for the Ni–20Co–15Cr–3.5Al–6Mo
alloy (Fig. 11). The higher content of Cr and Co in the Ni–20Co–15Cr–3.5Al–6Mo
model alloy reduced the degradation of the alloy. This point is supported in
considering that the degradation on IN-617, which contains higher amounts of Mo,
Co, and Cr than the NiCrAlMo alloy, was also less than that for the NiCrAlMo
alloy. The Ni–20Co–15Cr–3.5Al–4Mo model alloy has a chemical composition
similar to RR1000. The degradation of both alloys after exposure at 700 �C for
100 h with Na2SO4 deposit was comparable. The main difference was that a Cr-, Al-
, Ni-rich oxide formed at the alloy/salt interface on the Mo-containing model alloy
(Fig. 11), whereas some internal oxides were present in the commercial alloy
(Fig. 5).
Fig. 12 SEM images of Mo-containing alloys with Na2SO4 deposit after exposure in air at 700 �C for100 h
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Discussion
Exposure in Air Without Na2SO4
The results from the air exposures at 700 �C of the Mo-containing alloys without
Na2SO4 deposit indicated that Mo played a role in the oxidation behavior of the
alloys even when no sulfate salt deposit was present. Less protective oxide scales
grew on those alloys with significant Mo-amounts (Figs. 5, 11). This is due to Mo
becoming incorporated in the oxide scale and resulting in the formation of a mixed
oxide instead of a more protective Al2O3 or Cr2O3 layer. Contrary to Al2O3 or
Cr2O3, refractory metals, e.g., Mo, have poor oxidation resistance and follow a
linear oxide growth rate and form oxides which have high volatility.
Exposure in Air with Na2SO4
The results from the Mo-containing commercial and model alloys with Na2SO4
deposit exposed in air at 700 �C (Figs. 6, 12) indicated that the alloys which
contained an appreciable amount of Mo (4–9 wt%) suffered severe degradation. The
corrosion microstructures are somewhat similar to what has been reported for alloy-
induced acidic fluxing Type I hot corrosion. A typical alloy-induced acidic fluxing
Type I hot corrosion microstructure from a 900 �C exposure is shown in Fig. 13 for
the NiCrAlMo alloy compared to the results from the same alloy tested at 700 �C in
air. The degradation is much greater for the Type I hot corrosion image. The
appearance of Mo-rich oxides at the alloy/oxide interface is a key characteristic of
Fig. 13 Comparison of corrosion of NiCrAlMo model alloy with a Na2SO4 deposit at 700 �C (left) in airto Type I hot corrosion at 900 �C (right)
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alloy-induced acidic fluxing. The Mo-rich oxide is more pronounced on the alloy-
induced acidic fluxing image, but Mo-rich oxide was detected not only on the
NiCrAlMo model alloy but on all alloys with a Mo content higher than 2 wt% tested
in air at 700 �C. If the tests on the Mo-containing disk alloys were conducted for
longer times, the microstructure would likely become similar to the Type I hot
corrosion image.
The tests were conducted below the melting point of Na2SO4, and there is no SO3
in the gas atmosphere to form a low melting MSO4–Na2SO4 solution. A liquid is
forming to cause corrosion. However, transient MoO3 reacts with the salt to form
sodium molybdate releasing SO3. Sodium molybdate has a melting point (687 �C)which is below the temperature tested, so it becomes liquid. This reaction forming
sodium molybdate reduces the Na2O activity and causes the melt to become highly
acidic, producing a negative solubility gradient at the oxide/salt interface which
results in dissolution and fluxing of the NiO or Al2O3 oxide scales. The fluxing of
both the transient NiO and protective Al2O3 scale causes Ni2?, Al3?, and MoO42-
ions to diffuse through the salt where Al3? and/or Ni2? are reprecipitated where the
pO2 is higher as a non-protective scale, and the MoO3 evaporates. The presence of
Ni-rich oxide was not only observed as an outside layer but also at what appeared to
be the deposit particles edges (Fig. 12). EDS and WDS measurements indicate that
there are significant areas of pure Na2SO4 present in the deposits. These presumably
remained solid during the exposure. Apparently, Mo has only interacted with the
Na2SO4 at the salt/alloy interface and boundaries between the Na2SO4 particles, and
the Na2MoO4 is able to wet the interparticle boundaries. The interparticle veins are
the Ni transport paths through the deposit to form the NiO observed at the salt/gas
interface. Degradation is less than that observed at 900 �C where the entire deposit
is liquid and provides a much larger transport cross section for Ni. The lack of a
protective oxide scale results in rapid oxidation/corrosion and the observed
corrosion products. This form of attack is self-sustaining, because MoO3 is able to
continue to form at the alloy/salt interface and evaporate at the salt/gas interface.
The alloy degradation was observed to increase with increasing Mo content of the
alloy. Higher Mo content led to the formation of more transient MoO3 at the alloy/
salt interface, which reacts with the Na2SO4 deposit resulting in formation of more
sodium molybdate and thus more degradation. The released SO3 from the reaction
in Eq. 6 may be the cause of the internal sulfides observed in the corrosion products,
i.e., sulfur dissolves in the alloy and then reacts with Cr to form sulfides. The
mechanism which operates for the Mo-containing disk alloys at 700 �C with a
Na2SO4 deposit in air is explained very similarly by the mechanism for alloy-
induced acidic fluxing Type I hot corrosion proposed by Goebel et al. [8]. This is
important, as alloy-induced acidic fluxing is believed to only occur at high
temperatures (above 900 �C). A low-temperature (700 �C) alloy-induced acidic
fluxing mechanism is able to occur for Mo-containing Ni-based alloys in air
atmospheres.
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Effect of Temperature
The alloys exposed at 700 �C underwent significant corrosion if their Mo content
was greater than 2%. The tests performed at 650 �C, below the melting point of
sodium molybdate, produced no severe corrosion and tend to confirm that Na2MoO4
is the cause of corrosion. The results of the NiCrAl and NiCrAlMo alloys tested at
800 �C and 900 �C showed that the amount of degradation increases with
temperature, and the Mo-containing alloy was catastrophically degraded. The
Mo-containing alloy has more degradation at 800 and 900 �C because not only does
molten sodium molybdate form, but also MoO3 becomes molten at 795 �C, somolten MoO3 contributes to the degradation at 800 �C. The entire salt deposit is
molten at 900 �C, which is above the melting point of pure Na2SO4. Figure 14
shows a schematic diagram of the corrosion rate versus temperature for NiCrAl and
NiCrAlMo alloys. The ‘‘double peak’’ generally observed for NiCrAl alloys [2] is
absent for NiCrAlMo alloys. (Of course, the 700 �C peak only occurs for NiCrAl in
the presence of SO3.) Since the salt can become molten at any temperature above
the melting point of Na2MoO4 (687 �C) and the dissolution of Mo into the salt
creates acidic conditions at the salt/alloy interface alloy-induced acidic fluxing can
over the entire temperature range with the corrosion rate increasing monotonically
with temperature.
Effect of Alloy Composition
The experimental results indicate that not only Mo played a role on the degradation
of the Mo-containing disk alloys, but that the extent of the degradation also
depended on the additional components of the alloy. It was observed that for a given
Mo content, i.e., 6 wt%, the degradation was greater in the Ni–8Cr–6Al–6Mo alloy
Fig. 14 Schematic diagram comparing the temperature dependence of the hot corrosion NiCrAl alloys(two peaks) and NiCrAlMo alloys (single peak)
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than in the Ni–20Co–15Cr–3.5Al–6Mo alloy which contains higher amounts of Cr
and Co but less Al. The higher amount of Cr is likely responsible for the better
resistance of the Ni–20Co–15Cr–3.5Al–6Mo. The degradation for alloy RR1000
and Ni–20Co–15Cr–3.5Al–6Mo alloy, which have similar chemical composition,
was as can be seen in Figs. 5 and 12 although slightly more oxide formed for the
model alloy. Some Cr2O3 and Al2O3 internal oxides formed in RR1000, whereas the
Ni–20Co–15Cr–3.5Al–6Mo alloy formed a Cr-, Al- Ni-rich oxide as well as some
Mo-containing oxide at the alloy/salt interface. The formation of such oxide scale
morphology in the latter is likely due to the slightly higher Mo content of the Ni–
20Co–15Cr–3.5Al–6Mo alloy.
Exposure in O2 1 1000 ppm SO2 with Na2SO4 Deposit—Type II HotCorrosion
When the alloys were exposed to Type II hot corrosion conditions (700 �C in
O2 ? 1000 ppm SO2 with Na2SO4 deposits) the degradation was more severe in the
SO2-containing atmosphere than in air. For the alloys with little or no Mo, the SO3
in the gas atmosphere allows for the low melting MSO4–Na2SO4 eutectic to form
and cause dissolution and fluxing of the oxide scales by mechanisms described for
Type II hot corrosion proposed by Luthra [16] and Chiang et al. [17]. The Mo-
containing alloys were also more degraded after exposure in the SO2-containing
atmosphere. The mechanism in this case is the same as that described above for air
exposures except that more liquid is available because of the formation of the
MSO4–Na2SO4 eutectic. Misra [18] observed that the corrosion of U-700 at 950 �Cin O2 ? SO2 was dependent on the gas atmosphere. Small amounts of SO2
(*1000 ppm) accelerated the corrosion while large amounts (*10,000 ppm)
retarded it. At this temperature, all of the species in the deposit are liquid. The
acceleration was explained by SO3 providing a second species (S2O7-2) to transport
oxygen through the melt. The retarding effect at high SO2 concentrations was
explained by the reverse of Eq. 6 consuming sodium molybdate, i.e.,
Na2MoO4 þ SO3 ¼ Na2SO4 þ MoO3 ð7Þ
Kuepper and Rapp [19] have also reported that the ionization state of Mo in
Na2SO4 at 927 �C depends on the SO3 partial pressure in the gas phase. However,
high concentrations of SO2 were not included in the present study.
Conclusions
The corrosion resistance of Mo-containing disk alloys with and without Na2SO4
deposit was examined at 700 �C in air and in SO2-containing atmospheres. The
results of the tests can be summarized as follows:
• Exposure in air at 700 �C without Na2SO4 deposit:
• Less protective oxide scales formed with increasing Mo content of the alloy.
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• Exposure in air at 700 �C with Na2SO4 deposit:
• Those alloys containing sufficient Mo (4–9%) underwent degradation in air
atmospheres similar to alloy-induced acidic fluxing Type I hot corrosion.
• The tests were conducted below the melting point of Na2SO4, and in the
absence of an SO3-containing atmosphere, a low melting eutectic would not
be expected to form.
• Sodium molybdate can form from a reaction with transient molybdenum
oxide and the sodium sulfate deposit. Sodium molybdate has a melting
temperature of 687 �C, and thus is molten at a typical Type II testing
temperature (700 �C).• The formation of liquid sodium molybdate causes the melt to be highly
acidic, resulting in acidic dissolution and fluxing of the protective oxide
scales.
• Mo-containing alloys undergo a low-temperature alloy-induced acidic
fluxing mechanism in the absence of a SO3-containing atmosphere.
• The degradation of the alloy not only depended on the Mo content of the
alloy but also on other alloying elements, e.g., Cr and Co.
• Exposure in O2 ? 1000 ppm SO2 at 700 �C with Na2SO4 deposit:
• Low Mo alloys exposed to O2 ? 1000 ppm SO2 underwent Type II hot
corrosion, caused by the formation of the low melting MSO4–Na2SO4
eutectic and fluxing of the oxide scales.
• The high-Mo alloys were also more degraded after exposure in the SO2-
containing atmosphere because more liquid is available from the formation
of the MSO4–Na2SO4 eutectic.
Acknowledgements The authors gratefully acknowledge the Office of Naval Research for support of
their participation in this collaboration under ONR Contract N00014-10-1-0661, David A. Shifler,
Scientific Monitor.
Compliance with Ethical Standards
Conflict of interest The authors declare that they have no conflict of interest.
Human and Animals Rights There were no human participants or animals in this study.
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