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

ghmeier@pitt.edu

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.

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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

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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

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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

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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 –

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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

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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

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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

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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

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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

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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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