High-temperature corrosion of pure Ni 3 Al and its alloyed (2.99 wt.%Ti) in Ar-0.2%SO 2 gas environment

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Original Article

High-temperature corrosion of pure Ni 3 Al and its alloyed (2.99 wt.%Ti) in Ar-0.2%SO 2 gas environment

Poonam Yadav

^a

, Muhammad Ali Abro

^b

, Dong Bok Lee

^c

, Jonghun Yoon

^a,d,^*

aDepartment of Mechanical Engineering, BK21FOUR ERICA-ACE Center, Hanyang University, Ansan-si, Gyeonggi- do, 15588, Republic of Korea

bDepartment of Mechanical Engineering, MUET, SZAB Campus, Khairpur Mir's, 66020, Pakistan

cSchool of Advanced Materials Science&Engineering, Sungkyunkwan University, Suwon, 16419, Republic of Korea

dAIDICOME Inc., 55, Hanyangdaehak-ro, Sangnok-gu, Ansan, Gyeonggi 15588, Korea

a r t i c l e i n f o

Article history:

Received 6 May 2021 Accepted 11 February 2022 Available online 17 February 2022 Keywords:

Corrosion Diffusion High-temperature Kinetics

Spallation Oxide scale

a b s t r a c t

In this work, the high-temperature corrosion behaviour of pure Ni3Al (PeNi3Al) and alloyed (2.99 wt.%Ti) Ni3Al (TieNi3Al) was investigated in Ar-0.2%SO2gas at 900e1100C for up to 100 h. The corrosion kinetics of PeNi3Al and TieNi3Al reveal that Ti addition increased the total weight gain at all temperatures by approximately 10 times than without Ti (PeNi3Al).

The alloy initially gained more weight with the increase in temperature; but later on, the corrosion kinetics changed. Because of the extensive scale spallation during cooling, which causes the creation of large and deep geometric voids, the corrosion kinetics of PeNi3Al deviated from the parabolic rate law. At all temperatures, Ti strengthened the scale adherence as it occupied the Al substitutional sites with a broad atomic radius, and facilitated the creation of ordered phases known as the gamma prime phase (g⁰). Owing to the ordered structure, it was assumed that the diffusion of occupying atoms would be slower, thereby increasing the scale adherence. Darker inclusions were found in TieNi3Al at the scaleematrix interface, which were rich in TiS owing to inward sulphur diffusion.

1. Introduction

Nickel aluminide (Ni3Al) is an intermetallic compound that has attracted significant technological interest owing to characteristics such as high strength, improved high-temperature oxidation and corrosion resistance, and low density and production cost [1,2]. Ni3Al alloys are among the most promising structural materials with extensive applications at high temperatures [3‒5]

because of their ordered L12 structure based on the face-centred cubic lattice (A1), with a lattice parameter ofa¼ 0.3572 nm.

Several studies have been conducted on the oxidation and oxide composition of Ni3Al [6‒13]. Ni3A1 has high oxidation resistance and parabolically oxidises with very a low weight gain [11,12]. At 900C, Ni3Al is oxidised to NiO towards the surface,a-A12O3to- wards the matrix, and NiA12O4 between NiO and a-A12O3. However, at 1200C, Ni3Al exhibits poor scale adherence and low oxidation resistance, leading to the evolution of voids at the

*Corresponding author.

E-mail address:[email protected](J. Yoon).

Available online atwww.sciencedirect.com

journa l hom epage: www.elsevier.com /locate/jmrt

https://doi.org/10.1016/j.jmrt.2022.02.057

creativecommons.org/licenses/by-nc-nd/4.0/).

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matrixeoxide interface due to the Kirkendall effect [11,12,14,15].

With the addition of Ti, the oxidation rate of Ni₃Al increased at 1027C and 1127C, but decreased at 1227C. The oxide scales that formed became protective despite the development of interfacial voids, which is the reason for poor scale adherence [12]. In another study, Cr, Co, and Ti were added to pure Ni3Al to explore their characteristics in high-temperature oxidation at 1 atm and 1000, 1100, and 1200C in air. The results were then compared with those of pure Ni3Al [13]. At each temperature, isothermal oxidation exhibited parabolic kinetics. At 1000C, the initial oxide scales on the surface of the Ni3Al-based alloys consisted mainly of NiO,q-A12O3, and NiA12O4. However, as the reaction progressed, the outer layer changed into NiA12O4

containing small NiO particles, and the inner layer changed into q-A12O3[13]. At 1100C,a-A12O3was the main phase, while NiO, NiA12O4, and TiO2were the minor phases [13].

The SO2corrosion rate of pure Ni3Al was significantly higher than that of oxidation, resulting in the formation of scales containing oxide and sulphide mixtures [16]. These results show that pre-oxidation significantly enhanced the high- temperature corrosion resistance of Ni3Al in a wet hydrogen atmosphere. Ni3Al oxidation and sulphidation at 605, 800, and 1000C were also investigated in air and air-1%SO₂. This study found that sulphur substitution at the oxygen sites in NiO ap- pears to reduce the concentration of electron holes in the valence band. With the partial oxygen constant, a decrease in electron hole concentration increased the Ni cation vacancies, resulting in a higher Ni flux at the gas-oxide interface. The self- diffusion rate of Ni^2þions, which determines the oxidation rate, also increased [17]. Owing to the high mobility and/or concentration of nonstoichiometric sulphide point defects, sulphidation is usually a more serious issue than oxidation, with sulphidation rates 10e100 times higher than oxidation rates [2,18]. The objective of this study is to examine the high- temperature corrosion kinetics and developed scales for pure Ni3Al and Ni3Alþ2.99 wt.%Ti alloy in Ar-0.2%SO2gas.

2. Materials and methods

Pure Ni3Al (PeNi3Al) and alloyed (2.99 wt.%Ti) Ni3Al (TieNi3Al) were prepared from high-purity Ni (99.97%), Al (99.993%) and Ti (99.79%) metals in a vacuum induction melting furnace with a nominal composition (wt.%) of 84.56Ni-13.29Al and

84.44Ni-12.58Al-2.99Ti respectively. Specimens measuring approximately 10102 mm³were cut from alloy ingots, placed in a Vycor tube, vacuum-sealed, and homogenised for 24 h at 1200C. These specimens were then cooled in air after homogenisation; ground using 1500 grit SiC paper to eliminate any oxides; and thoroughly rinsed in acetone. Each specimen was corroded in flowing commercial-grade Ar (99.9999% purity)-0.2%SO2 (99.9% purity) gas mixtures at 900, 1000, and 1100C for up to 100 h. A thermogravimetric analyser (TGA, Cahn Thermax 2141, USA) was used to evaluate the weight changes as a function of time during corrosion. A platinum wire was used to suspend each sample and expose it to the uniform zone of the TGA-connected furnace. Field-emission scanning electron microscopy (FE-SEM) with an energy dispersive spectrometer (EDS), a high-power X-ray diffrac- tometer (HP-XRD) with Cu-Ka radiation, a field-emission electron probe micro-analyser (FE-EPMA) with EDS, and a high-resolution transmission electron microscope with EDS were used to analyse the morphology of the samples.

3. Results and discussion

The optical microscopic images of as-cast PeNi₃Al and TieNi3Al after etching in aqua regia solution for 5 min are shown in Fig. 1(a) and (b), respectively. Both PeNi3Al and TieNi3Al display theg⁰-phase during the solidification process.

PeNi3Al precipitates theg⁰-phase in the dendritic core, while TieNi3Al forms equiaxed grains with theg'-phase. Thus, this phase is favourable for high-temperature applications.

Figure 2 shows the corrosion kinetics of PeNi3Al and TieNi3Al as a function of time in terms of the mass gain per unit surface area at 900, 1000, and 1100C for 100 h in Ar-0.2%

SO2gas. For the overall weight gain over time, PeNi3Al adop- ted the parabolic time law and initially gained weight. How- ever, the weight gain curves later fluctuated because of the scale spallation arising from sulphur inclusion in the scale, as shown inFig. 2(a). The presence of Ti in TieNi3Al increased the overall weight gain at all temperatures by approximately 10 times. Owing to increased diffusion and reaction rates, the weight gain increased as the temperature increased. As shown inFig. 2(b), the corrosion kinetics did not follow the well-known parabolic rate law because of early stage varia- tions in corrosion rate; Ti escape; aggressive corrosion of Ni, Fig. 1eOptical microscope images (etched by aqua regia solution) of as-cast (a) pure Ni3Al and (b) Ni3Alþ2.99 wt.% Ti.

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Al, and Ti; and local microscopic scale spallation during the later stage of corrosion.

In general, the corrosion kinetics is represented by a power law in the parabolic oxidation process [11]:

Dm¼ ffiffiffiffiffiffiffiffi kp:t q

(1) whereDm,t, andkpare the mass gain per unit surface area of the sample, corrosion time, and parabolic rate constant, respectively. Kpfollows the equation's Arrhenius relation and is defined as

kp¼koexp Q_eff.

RT

(2) whereQeffis the effective corrosion activation energy,kois the material constant,Tis the absolute temperature, andRis the gas constant.

Figure 3shows a logarithmic scale graph of k_pversus 1/T for PeNi3Al and TieNi3Al, including other reference data [19‒ 23]. The activation energies are provided by the slope of the plot as 109.81 kJ/mol and 171.56 kJ/mol, respectively. The higher the activation energy, the lower the corrosion rate;

hence, the addition of Ti improves the corrosion resistance of TieNi3Al. Moreover, Ti strengthens the scale adherence.

As shown in Fig. 3, PeNi3Al oxidises to NiO and a-Al2O, TieNi3Al oxidises to NiO, TiO2, anda-Al2O3, while Ni and Ti sulphidise to NiS and TiS. The sulphidation rate is much higher than the oxidation rate [24,25]. The overall oxidation rate suggests that although TieNi3Al strengthens the scale adherence and improves the corrosion resistance, it also provides a sulphur ingress path where sulphur diffuses inward and forms NiS and TiS. This leads to the formation of inclusions, which are poor barriers to corrosion. Corrosion is primarily regulated by oxidation rather than sulphidation;

compared with PeNi3Al, TieNi3Al shows better corrosion resistance.

The SEM/TEM results of PeNi3Al following Ar-0.2%SO2gas corrosion at 900C for 100 h are shown inFig. 4(a)e(f).Figure 4 shows that the scale has two layers. InFig. 4(a), a closer study P-Ni3Al

(a) (b) Ti-Ni3Al

Fig. 2eCorrosion kinetic curves of (a) pure Ni3Al and (b) Ni3Alþ2.99 wt.% Ti exposed to Ar-0.2%SO2gas at 900, 1000, and 1100C for 100 h.

Fig. 3eTemperature dependence of logarithm parabolic constants of pure Ni3Al, Ni3Alþ2.99 wt.% Ti, and other reference data [28‒32].

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of the SEM image of the top-view reveals that the corrosion products began to expand through the creation of a very thin layer of corrosion. This corrosion layer is discontinuous because of scale spallation. Scale discontinuity and spallation is confirmed by the TEM cross-sectional image (Fig. 4 (b)), which shows voids between the outer and inner scales. In Fig. 4(b) and (c), EDS spots 1e5 show that the outer layer primarily consists of NiO. Al was enriched with an inner layer (spots 7e13), and a protective Al2O3layer was formed. The EDS concentration profile inFig. 4(c) and EDS mapping inFig. 4(d) show the same results.Figure 4(e) and 4(f) depict the selected area (electron) diffraction (SAED) patterns of spots 3 and 10, and indicate the formation of the outer NiO scale and enriched inner Al2O3scale due to the outward diffusion of Ni^2þand Al^2þ, or inward diffusion of O².

Figure 5shows the SEM/EDS results of PeNi3Al following Ar-0.2%SO2 gas corrosion at 1000 C for 100 h. As seen in Fig. 5(a), the oxide scales spalled in large flakes on PeNi3Al.

The magnified SEM image inFig. 5(b) is the square region in Fig. 5(a), where it shows outer equiaxed fine NiO grains following scale spallation with numerous geometric voids in the exposed region. The primary cause of scale spallation is geometric void formation. The primary defect in NiO oxide is Ni vacancy [15]. The cross-sectional SEM image in Fig. 5(c) shows that the scales (~3.14mm thick) can be split into spot 1 as the outer scale, spot 2 as the middle scale, and spot 3 as the inner scale.Figure 5(d) shows that spots 1, 2, and 3 consist of 55.5Ni-44.5O, 1.2Ni-51.4Al-47.4O, and 18.7Ni-44.0Al-37.33O, respectively.

All compositions are subsequently represented as atomic percentages (at.%). The NiO outer scale as the outer layer (spot 1), NiAl2O4mid-scale as the intermediate layer (spot 2), and Al2O3 inner scale as the innermost layer (spot 3) are well balanced by such a scale composition. The formation of the outer NiO scale was regulated by the external diffusion of Ni^2þ ions, which depleted Ni at the middle and innermost layers;

Fig. 4ePure Ni3Al after corrosion at 900C for 100 h in Ar-0.2%SO2gas: (a) SEM top-view, (b) cross-sectional TEM image, (c) concentration profile across the scale of (b); (d) TEM mapping of (b); (eef) SAED pattern of spot 3 and spot 10 of (b).

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hence, NiAl2O4serves as a diffusion barrier against outward- migrating Ni ions, indicating a slow oxidation rate [12].

Therefore, the scale consisted mainly of NiO, NiAl2O4, and Al2O3without sulphides in the presence of Ar-0.2%SO2gas.

This suggests that the tendency of oxides is greater than that of sulphides.

PeNi3Al showed strong corrosion resistance at 1000 C owing to such oxide stability and the rapid establishment of protective oxide scales with sulphide suppression; however, the scale was vulnerable to microcracking and spallation. At low concentrations, sulphur joins NiO as S⁰ions instead of S⁰⁰ ions. S⁰ions are the most common form by which sulphur in most oxides enters the matrix [16]. Sulphur can serve as an electron donor under the equilibrium:

1/2S2þOoþh^$¼SO^$þ1/2O2 (3)

On an anion site,Oois a natural oxygen ion,h^$is the concentration of electron holes, and SO^$is the sulphur ion at the site of an oxygen anion.

Eq.(3)and the following reaction:

1/2O2(q)¼V^⁰⁰Niþ2h^$þOo (4)

Could be formed with the dissolution of S into NiO.

V⁰⁰Niare the cation vacancies, i.e. vacancies in the Ni site with a charge of^-2 compared with the usual occupation of the site.

The equilibrium constant forEq. (4) can be written as [V^⁰⁰Ni] [h^$]²¼K(constant)PO21/2 (5)

Sulphur substitution at the oxygen sites in NiO tends to decrease the concentration of electron holes in the valence

band. A decrease in [h^$] increases V⁰⁰Ni(Eq.(5)) with a partial oxygen constant. Thus, there will be greater Ni flux to the gas- oxide boundary. Consequently, the oxidation rate increased with a low sulphur concentration [16,17]. As Kirkendall voids are driven by vacancies, they grow on the substrate that causes scale spallation. The wide and deep voids found in PeNi3Al substrates serve as stress concentrators. When localised separation occurs in a void, the stress increases beyond the adhesion strength upon cooling, resulting in the separation of large oxide flakes [11].

Figure 6 shows the XRD/SEM/EPMA results for PeNi3Al following Ar-0.2%SO2gas corrosion at 1100C for 100 h. The XRD results indicate NiO,a-Al2O3, and NiAl2O4oxides as the main phases, and Ni3Al as the minor phase. Because the underlying matrix was exposed after extensive scale spallation, no sulphides were detected. The scales remained discontinuous owing to the stress caused by the huge amount of scaling, the thermal stress produced by the mismatch of thermal expansion coefficients between the scale and alloy, and the incorporation of sulphur into the scale (Fig. 6(b)).

Figure 6(c) presents the SEM image of fractured oxide scales detached from the matrix, which shows the division of the scale into three layers: the outer scale with fine and equiaxed NiO grains; the inner scale with coarse and columnara-Al2O3

grains toward the matrix side; and the middle scale with NiAl₂O₄and lateral oxide growth between the inner and outer scales. The decrease in nucleation rates at the scaleematrix interface was correlated with the thickening of columnar grains of the inner scale. In general, such growth is considered to be controlled by the inward diffusion of oxygen ions [17,26‒ 29]. The entire scale was ~5.81mm thick (Fig. 6(d)). Owing to the outward diffusion of Ni^2þions and the anisotropic volume expansion, voids formed close to the scaleematrix interface, and between the scales. Because the Kirkendall effect is Fig. 5ePure Ni3Al after corrosion at 1000C for 100 h in Ar-0.2%SO2gas: (a) SEM top-view, (b) magnified image of area marked in (a); (c) SEM cross-sectional image; and (d) EDS spectra of①,②, and③shown in (c).

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driven by vacancies, the growth of A12O3was correlated with the preferential consumption of aluminium in the alloy near the matrixeoxide interface [30], and substantial voids were observed on the substrate surface. Such Kirkendall voids can provide a path for oxygen and sulphur with fast inward diffusion. The growth of geometric voids, as seen inFig. 6(c) and (d), lead to extensive scale spallation after corrosion during cooling. The EPMA line profiles in Fig. 6(d) and (e) show that the outer scale primarily consisted of NiO (on the surface) and Al₂O₃-rich NiAl₂O₄ (towards the inner side), while the inner scale primarily consisted ofa-Al₂O3. The Ni depletion in the Al2O3-rich NiAl2O4middle layer was caused by the preferential oxidation of Ni on the scale through the outward diffusion of Ni^2þions. The EPMA line profiles of O, Ni, Al, and S along with the EDS quantitative composition analysis shown inTable 1at spots 1, 2, and 3 further verified that the outermost scale was Al-rich with the dissolution of Ni; the middle scale was Al-enriched and Ni-deficient; and the voids consisted of mixed NieAl oxides.

The focused ion beam (FIB)-SEM/EPMA/TEM/EDS results of TieNi₃Al following Ar-0.2%SO₂ gas corrosion at 900 C for 100 h are presented inFig. 7 (a)e(d). The corrosion layers were

analysed using the FIB cross-sectional SEM image (Fig. 7(a)), EPMA cross-sectional back-scattered electron images and line profiles (Fig. 7(b) and (c)), and TEM cross-sectional images along with EDS concentration profiles and SAED patterns (Fig. 7 (d)e(g)). Figure 7(a) shows that, unlike PeNi3Al, the adherent oxide scales covered the alloy surface with equiaxed micrometre-sized NiO grains; the oxide scales did not split into pieces. The cross-sectional portion showed that the oxide layer was compact and remained attached to the surface and matrix, thus providing the alloy with adequate protection.

However, there were a few sub-micrometre-sized voids on the scale. The addition of Ti increased the adherence and conti- nuity of the oxide scales, which can be considered as the Fig. 6ePure Ni3Al after corrosion at 1100C for 100 h in Ar-0.2%SO2gas: (a) XRD pattern, (b) SEM top-view, (c) fracture image of oxide scale, (d) back-scattered electron cross-sectional image, and (e) EPMA line profile along A-B of (d).

Table 1eEPMA quantitative analysis of PeNi3Al following Ar-0.2% SO2gas corrosion at 1100C for 100 h.

Spot No. Al (at. %) Ni (at. %) Total (at. %)

1 13.75 3.38 17.16

2 15.82 0.28 16.1

3 12.93 4.60 17.53

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Fig. 7eNi3Alþ2.99 wt. % Ti after corrosion at 900C for 100 h in Ar-0.2%SO2gas: (a) cross-sectional (FIB)-SEM image;

(b) EPMA cross-sectional image; (c) EPMA line profile along A-B of (b); (d) TEM cross-sectional image of corrosion layer;

(e) TEM-EDS concentration profiles along spots 1e15 marked in (d); and (feg) SAED pattern of spots 4 and 8 of (d).

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primary obstacle to reducing the permeability of corrosive products. During corrosion, most thermodynamically stable Al2O3oxides are initially formed with the dissolution of Ni and Ti. Over time, Ni ions diffuse outward and form NiO as a top layer. The oxide scale has an outer and inner layer with a thickness of approximately 1.0mm and 1.04mm, respectively (Fig. 7(b)).Figure 7(c) indicates that the outer layer is rich in Ni with high oxygen affinity, while the inner layer is rich ina- Al2O3from the dissolution of Ni and Ti, with the possibility NiA12O4 and TiO2oxide formation. Figure 1(b) showed that during the solidification process, equiaxed grains were formed because of the addition of Ti. The presence of defects (voids, porosity, and microcracks) in oxide scales provided a pathway for the flow of oxygen and sulphur into the alloy, which pro- motes intergranular corrosion. At the interface of the oxides and metal surfaces, an oxygen-affected zone was formed.

These interfaces are Al-depleted areas because of the consumption of Al to Al2O3. Thus, because of the presence of grain boundaries, corrosion occurred along the grain boundaries, with the possibility of NiO and TiO2formation. As TiO2is a reducible oxide [31], it has a tendency to form oxygen vacancies occupied by sulphur, and the TiS that forms along the grain boundaries and oxygen released by TiO2participates in the formation of the internal oxide zone. As a result of the intergranular corrosion along the grain boundaries, darker

inclusions were observed at a depth of approximately 4.8mm;

these were enriched with TiS due to the diffusion of sulphur into the matrix. According to the TEM-EDS concentration profile (Fig. 7(d) and (e)) across the scale that developed on TieNi3Al, spots 1e6 strongly correspond to NiO owing to the outward diffusion of Ni^2þions. Spots 7e9 may have formed NiAl2O4 with a small amount of Ti and S, which acts as a diffusion barrier against outward-migrating Ni ions. The composition of spots 10e12 strongly correspond to Al2O3. This suggests that the entire oxide scale grew with the NiO outer layer. Below this, an intermediate layer consisting of NiAl2O4

with a negligible amount of TiO2 grains, and an innermost layer of a-Al2O3oxide, were formed. The SAED patterns of spots 4 and 8 inFig. 7(d) confirm the development of the outer scale of NiO and inner scale ofa-Al2O3.

3.1. Mechanism of scale adherence in Ti containing Ni₃Al

Because Ti has a large atomic radii that occupies the Al substitutional sites (i.e. at the cube corners in Ni₃Al), the inclusion of Ti in Ni3Al increases the scale adherence. This facilitates the formation of ordered phases known as the g⁰-phase, which plays a vital role in strengtheningg'[32,33]. However, as the key defect in TiO2 is oxygen vacancies, the defect Fig. 8eNi3Alþ2.99 wt. % Ti after corrosion at 1100C for 100 h in Ar-0.2%SO2gas: (a) XRD pattern, (b) SEM top-view, (c) EPMA cross-section, and (d) EPMA line profile along A-B of (c).

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concentration (n_e) varies with the oxygen partial pressure pO₂¹⁼⁶, expressed as

n²_e¼ ð2KÞ¹⁼³pO21=6 (6)

From this relationship, it is inferred that the concentration of vacancies will decrease as more oxygen is supplied to the alloy [34]. Owing to the inclusion of Ti at the Al site with larger radii, the relaxation of geometric results of the local expansion at the defect site resulted in fewer Kirkendall voids relative to PeNi3Al, and improved adherence property.

The XRD/SEM/EPMA results of TieNi3Al following Ar-0.2%

SO2gas corrosion at 1100C for 100 h are shown inFig. 8. The XRD results inFig. 8(a) show that NiO anda-Al2O3are major phases along with NiAl2O4and TiO2via Ni3Al peak retarda- tion. The reduction of Ni3Al peaks suggests that the presence of Ti improved the adherence property, owing to the reduction in geometric voids. The SEM top-view (Fig. 8(b)) showed that during corrosion, round and elliptical grains were formed, which consisted of NiO mixed with NiAl2O4, a-Al2O3, and rutile-TiO2. No sulphides were detected in the outer scale, as shown in the XRD and EPMA line profile.Figure 8(c) and 8(d) show that with the dissolution of Al, Ni, and O, the outer scale (approximately 0.44mm thick) was Ti-rich. With the dissolution of Ti and Ni, the inner scale (approximately 54.55 mm thick) was Al2O3-rich, with the possibility of NiAl2O4and TiO2

formation, as confirmed by the XRD pattern inFig. 8(a). An oxygen-affected zone was formed beneath the inner scale with a large number of darker inclusions. When the corrosion temperature increased, these inclusions increased and were

identified as TiS due to the inward diffusion of sulphur pre- sent in the Ar-0.2%SO2gas atmosphere.

3.2. Schematic

The overall corrosion sequence of PeNi3Al and TieNi3Al at 900e1100C for 100 h in Ar-0.2%SO2gas is illustrated inFig. 9.

During the initial oxidation (stage I), a two-layered scale formed on PeNi3Al with an outer NiO layer and an inner Al2O3

scale. TieNi3Al also developed a two-layered scale,viz.Stage I with an outer NiO and an inner (TiO2/NiAl2O4/Al2O3)-mixed scale, and an oxygen-affected zone beneath the inner scale. Ti diffused outward from the matrix along with Al in TieNi3Al.

As shown by the results inFig. 4‒6 and 8(i.e. stage I), scale adherence deteriorated owing to sulphur incorporation [35].

Scale spallation occurred because of the following reasons.

Initially, stress developed due to the dissolution of foreign elements and phase transformation of Al under NiO (i.e. stage I) or NiAl2O4 (i.e. stages IIeIII) from the substrate to Al2O3. Later on, thermal stress was generated between the oxide scales, i.e. NiO/NiAl2O4/Al2O3(Figs. 4 and 5) due to thermal expansion mismatch and alloy contraction, which led to microcracking at the oxide layer/matrix interface, i.e. (NiO/

NiAl2O4/Al2O3)/(matrix) during the cooling process (Figs. 6 and 8). Scale spallation occurred primarily in PeNi3Al and increased with the increase in corrosion temperature. How- ever, on TieNi₃Al, the scale was rather thick owing to foreign cation dissolution in NiO, and thermal expansion mismatch of NiO with the underlying rutile-TiO2layer (Fig. 7(a) and (b), and

Fig. 9eSchematic diagram of corrosion sequence of PeNi3Al and TieNi3Al after corrosion at 900e1100C for 100 h in Ar- 0.2%SO2gas.

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gen vacancies and forms darker inclusions (TiS), as shown in Figs. 7(c) and Fig. 8(d). The oxygen-affected zone widened internally towards the substrate with the increase in corrosion temperature. TiS also increased because of inward sulphur diffusion. With increasing corrosion temperature, the diffusion of Ni cations increased, resulting in the formation of NiO grains on the outer surface (Figs. 5(b), 6(c), and 8(b)). Oygen enters the scale easily and strengthens its affinity to form the thermodynamically stable inner oxide Al2O3layer. In stage I, a sufficient amount of oxygen remains to form less stable oxides (NiAl₂O₄) in PeNi₃Al. TiO₂/NiAl₂O₄embedded an inner oxide Al2O3 layer in TieNi3Al when most of the Al was consumed at a particular volume. Therefore, the scale increased continuously with increasing corrosion time. The spallation rate also increased the number of geometric voids along the scale, and loosened the surface scale, which peeled off. At this point, the oxygen in the Ar-0.2%SO2gas continuously diffused through the Kirkendall voids to the matrix (i.e.

stage III, where the alloy corroded at 1100C), which increased the Al2O3layer thickness. Outward Ni diffusion was hindered by the formation of semi-protective TiO2on the less stable NiAl2O4with fewer voids than in TieNi3Al. This resulted in a greater thickening of the Al2O3layer in TieNi3Al relative to PeNi3Al. Compared with PeNi3Al, this is the same reason for the increasing thickness of the outer oxide layer in TieNi3Al.

The oxide scale formation and its spallation continued to advance with increased corrosion temperature; thus, the corrosion rate tends to become more severe. The defects caused by scale spallation, void formation, and microcracking became simple routes for oxygen diffusion, and caused a much smaller amount of sulphur to diffuse inward towards the matrix, resulting in increased internal attacks. At the same time, Ni and Al continued to diffuse from the matrix to the scale. Oxide regeneration, scale breakdown, and advancement of the oxygen-affected zone acted as competi- tors, eventually causing serious damage to the oxide scale (i.e.

stage III).

4. Conclusions

Corrosion tests on PeNi3Al and TieNi3Al alloys in Ar-0.2%SO2

gas at 900, 1000, and 1100 C for 100 h showed that scale development on the surface followed the parabolic time law for the total weight gain over time. The addition of Ti in Ni3Al increased the overall weight gain by approximately 10 times at all temperatures. In PeNi3Al, the oxide scale formed at 900C for 100 h was divided into two layers. The outer layer (layer I)

alloy surface at 900C slowly shifted towards the NiA12O4outer layer (layer I), which contained a small amount of NiO particles, and the A12O3inner layer (layer II). The outer sections of the oxide layer enriched with NiO, and deeper in the oxide layera- A12O3, was the most dominant phase with a small amount of NiAl2O4, NiO, and TiO2in TieNi3Al at high temperatures. EPMA line profile analysis showed a few darker inclusions (TiS) at the matrixeoxide interface because of the inward diffusion of sulphur into the base material.

Author contributions

Poonam Yadav: Conceptualisation, Experimentation, Methodol- ogy, Investigation, Writing - original draft,Muhammad Ali Abro:

Methodology, Data curation, Visualisation, Writing - original draft, preparation,Dong Bok Lee: Investigation, Writingereview and editing, Jonghun Yoon: Project administration, Funding acquisition, Data Review, Revision, and Draft finalization.

All authors have read and agreed to the submitted version of this manuscript.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could appear to influence the work reported in this paper.

Acknowledgments

This work has been supported by the research fund of Hanyang University (HY-2020).

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