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International Journal of Corrosion
Volume 2019, Article ID 8517648, 8 pages
https://doi.org/10.1155/2019/8517648
Research Article

Isothermal Oxidation Behaviour of 69.5Fe-14Ni-9Al-7.5Cr Alloy at High Temperatures

1Department of Metallurgical Engineering, Faculty of Mining and Petroleum Engineering, Bandung Institute of Technology, Jl. Ganesha No. 10, Bandung 40132, Indonesia
2Nuclear Technology Center for Materials and Radiometry, National Atomic Agency of Indonesia, Jl. Taman Sari No. 71, Bandung 40132, Indonesia

Correspondence should be addressed to Eddy Agus Basuki; di.ca.bti.gninim@ikusab

Received 7 September 2018; Revised 11 December 2018; Accepted 30 December 2018; Published 3 February 2019

Academic Editor: Flavio Deflorian

Copyright © 2019 Eddy Agus Basuki et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Abstract

A series of experiments have been conducted to study the resistance of isothermal oxidation in 69.5Fe-14Ni-9Al-7.5Cr (% wt) alloys with some temperature variations of 800, 900, and 1000°C and each temperature is tested for 1, 10, 50, and 150 hours. Based on the analysis conducted on the test results, it is known that the higher the oxidation temperature, the faster the oxidation rate. As for the duration of oxidation, the longer the testing time is, the smaller the weight change per unit sample area is, indicated by a sloping graph. The oxide products found in this study were Fe2O3, Fe3O4, α-Al2O3, NiO, Cr2O3, and NiCr2O4. The rate of oxidation of this alloy at each temperature follows a logarithmic equation indicating that the oxide layer formed is very thin and can protect the metal surface well. The oxidation rate constants at temperatures of 800, 900, and 1000°C are 5.15E-05, 5.57E-05, and 6.74E-05 gr.cm−2, respectively.

1. Introduction

Relatively high global energy demand has forced the policies to increase the power plant efficiency, especially for those use fossil fuels, such as coal-based steam power plants. Increasing power plant efficiency can be achieved by increasing the pressure and temperature of the steam. Consequently, materials with capability to be used at higher temperature are demanding. Such materials should have high creep resistance as well as high temperature corrosion and oxidation.

The advanced ferritic steels with alloying elements of Ni, Al, and Cr have been the alternative materials used in some parts of power plant construction due to the combination of low thermal expansion coefficient and relatively low price compared with nickel-based super alloys [1]. In addition, these materials have high temperature strength resulted from the occurrence of coherent precipitate B2-FeNi(Al) in the ferrite matrix of α [2, 3]. This microstructure is essentially similar to that found in nickel-based superalloys, but these later alloys have austenitic matrix of γ. Relatively slight difference in lattice parameter α-Fe (A2) and precipitates NiAl (B2) as well as similar crystal structures of both phases provides the ordering and coherency that provide high strength [4]. Therefore, the strength or creep resistance of the materials increases with the volume fraction of B2-FeNi(Al) precipitates [5].

Stallybrass and coworkers [6] have studied the Fe-Ni-Al-Cr alloys and referred to the pseudobinary phase diagram of Fe+2.5 at.% Al – NiAl+1.25 at.% Al. The chemical composition was selected to produce various volume fraction of the alloys. Alloys with typically containing B2-FeNi(Al) dispersed homogeneously in the ferrite matrix of α are then known as ferritic superalloys [7]. These alloys have been recently studied, mostly focused on the mechanical properties as well as microstructural changes at high temperatures [810]. Various mixtures of phases in different quaternary alloy systems of Fe-Ni-Cr-Al at 900°C have been investigated by Yin et al. [11]. Three phases of α-ferrite, β-FeNi(Al), and γ-austenite were identified; alloy 70Fe3Cr12Al15Ni gave three phase of α+β+γ, while alloy 70Fe6Cr15Al9Ni produced α+β. Alloy containing two phase of β+γ had chemical composition of 70Fe3Cr9Al18Ni. Addition of Ni increases the toughness and high temperature strength, while addition of both Ni and Al promotes the precipitates of B2-FeNi(Al) that increases the creep resistance of the alloys [12]. Therefore, the strength and creep resistance of the alloy increase with increasing the volume fraction of B2-FeNi(Al).

Oxidation resistance of the ferritic alloys for high temperature application relies on the formation of a continuous slow growing protective scale based on Al2O3 [13]. Chromium addition can improve the high temperature oxidation as Cr can act as oxygen getter that reduces the critical aluminium content for external continuous layer of Al2O3 [14]. Therefore, in the case of B2-FeNi(Al) strengthened ferritic alloys, the involvement of Cr in the alloys can decrease the potency for B2 in substrate underneath the scale to dissolve.

Recently, we have developed a ferritic alloy with chemical composition of 69.5Fe-14Ni-9Al-7.5Cr that has precipitates of B2-FeNi(Al) dispersed homogeneously in the ferrite matrix of α. The alloy contains “bone-like” precipitates of B2, dispersed along certain crystallography orientation. The alloy is intended to be used for high temperature components of power generators, especially coal-based steam power plants. Most high temperature components used in such operations require high oxidation resistant properties. This paper discussed the results of isothermal oxidation behaviour of the alloy at 800, 900, and 1000°C.

2. Experimental Works

The alloy of this study had a chemical composition of 69.5Fe-14Ni-9Al-7.5Cr (otherwise stated the composition is in weight percent) or is equivalent with 72.74 at%Fe,15.40 at%Ni, 4.55 at% Al, and 7.31 at.% Cr. The raw materials for alloy making were low carbon ferronickel-shot obtained from PT ANTAM (Persero) Tbk., pure aluminium (>99%) obtained from PT. Inalum, and pure chromium and iron obtained from Sigma Aldrich. To obtain chemical composition of the alloy, raw material proportion was set up as shown in Table 1. The raw materials were melted in an argon purged single mini electric arc furnace at 220V and 75-90A for 2 minutes. In order to produce homogeneous alloy buttons, each with 20 grams weight, repeated melting was carried out for 5 times. Further homogenization was done by heating the alloy button at 1100°C for 24 in a horizontal tube furnace purged with high purity argon.

Table 1: Raw material proportion for alloy making.

The homogenized alloy buttons were then cut to produce samples of 6 mm x 6 mm x 1 mm in size using low speed diamond cutter. The samples were polished finish at 2000 grit and cleaned in acetone using an ultrasonic cleaner. Weighing and surface area measurement of each sample were carried out prior to oxidation test. Isothermal oxidation tests were carried out by heating the samples at 800, 900 dan 1000°C for 1, 10, 50, and 150 hrs. in a tube heat resistance electric furnace. Weighing of the oxidized samples was done to obtain the weight change of each sample. Microstructural analysis of the samples was carried out using optical microscope and scanning electron microscope (SEM) in which an energy dispersive X-ray spectroscopy is attached. To reveal the microstructure of the substrate, the as-homogenized alloy and the oxidized samples were etched in chemical solution containing 0.1 mL HF, 3.3 mL acetate acid, 3.3 mL HNO3, and 3.3 mL H2O. Point-count method was used to measure the volume fraction of primary B2 precipitates in the substrate of each sample based on the SEM images. X-ray diffraction analyzer (XRD) was used to identify the oxides formed on the surface of the oxidized samples. To confirm the occurrence and distribution of different oxides in the scales, X-ray mapping was applied.

3. Results and Discussion

The as-homogenized alloy sample of 69.5Fe-14Ni-9Al-7.5Cr essentially has two phases, as shown in Figure 1. It is seen clearly that rod-like ordered B2 (Fe,Ni)Al precipitates dispersed homogeneously throughout the ferrite-α (α-Fe) matrix. Identification of these two phases was based on the EDX analysis of the samples, as shown in Figures 2 and 3. However, different from these normally rod-type precipitates in ordinary alloys, the B2 (Fe,Ni)Al precipitates in this alloy have more likely bone-type precipitates, in which epiphysis-like structures are set up on both rod tips. These type precipitates have not been found in literatures discussing B2 contained ferritic steels. Most literatures indicate that B2 (Fe,Ni)Al precipitates have cuboidal or rounded type structures [6, 8, 9]. Heating the alloy at high temperatures changes the volume fraction, morphology, and size of the B2 precipitates. These microstructural changes occurred significantly at higher temperatures, especially at 1000°C, as seen in Figures 2 and 3. The as-homogenized particle size of B2 was about 1-2 μm in wide dimension and can reach up to more than 10 μm in length.

Figure 1: Dispersion of rod-like B2 (Fe,Ni)Al precipitates in the ferrite-α matrix of as-homogenized sample.
Figure 2: EDX analysis of phases in the sample heated at 800°C for 150 hrs.
Figure 3: EDX analysis of phases in the substrate sample heated at 1000°C for 150 hrs.

The volume fraction of the homogenized samples was 36.3%. Heating the alloy at different temperatures reduced the volume fraction of B2. Heating at 800, 900, and 1000°C, the volume fraction of B2 was 33.6%, 30.5, and 26.6%, respectively. Moreover, it was revealed that volume fraction of B2 precipitates decreased with time of heating, as shown in Figure 4. This decreasing volume fraction of B2 with time has been well recognized in the study of particle coarsening [15]. The stability of the B2 rods at 800°C was relatively high, as most B2 precipitates were still in rod type when the alloy was heated for 150 hrs. However, at 1000°C, the morphology of the B2 tended to be more rounded, due to the spheroidization of the B2 rods. These volume fraction and morphology changes of the B2 precipitates caused the hardness of the alloy to decrease, as shown in Figure 5.

Figure 4: Percentage of B2 particles in sample substrates after heating at different temperatures and times.
Figure 5: Hardness of the substrate in samples heated at different temperatures and times.

Exposing the alloy samples at high temperatures at atmospheric condition caused the alloy to oxidize followed by the scale formation containing various oxides on the surface of the alloy. Figure 6 show the scales on the samples oxidized at 800, 900, and 1000°C for 150 hrs. The thickness of the scales in varies from very thin of about 2 μm in sample oxidized at 800°C up to relative thick of about 8 μm in sample oxidized at 1000°C. EDS analysis of the scales shows that the scale contains mostly aluminium oxide and chromium oxide. No B2 precipitate depleted zones underneath the scales were found indicating that aluminium content in the α-Fe matrix in all samples was adequate to provide aluminium atoms to form Al2O3 in the scales. Moreover, no internal oxidation was found in all samples. The XRD analysis results, shown in Figure 7, confirmed that oxides of Fe3O4, Fe2O3, Cr2O3, Al2O3, NiO, and spinel of NiCr2O4 were detected in all sample surfaces. It is believed that the outer parts of the scales were dominated by iron oxide of Fe3O4 and Cr2O3.

Figure 6: Scale formed on the samples exposed within 150 hrs. at (a) 800°C, (b) 900°C, and (c) 1000°C.
Figure 7: X-ray diffraction patterns of samples oxidized within 150 hrs. at (a) 800°C, (b) 900°C, and (c) 1000°C.

X-ray mapping results confirmed further that most Al and Cr contained in the scales of samples heated at 900 and 1000°C, as shown in Figures 8 and 9, respectively. A particle found in the substrate of sample oxidized at 900°C, Figure 8, did not result from internal oxidation, but an Al2O3 inclusion in the alloy formed during melting is expected, as it is surrounded by relatively thick B2 phase. In scale of sample oxidized at 1000°C, it was found that aluminium oxide formed in the upper part as well as lower part of the scale.

Figure 8: X-ray mapping of scale on sample oxidized at 900°C for 150 hrs.
Figure 9: X-ray mapping of scale on sample oxidized at 1000°C for 150 hrs.

Iron oxides of Fe2O3, Fe3O4, and FeO are believed to be formed at the first stage of oxidation at all temperatures studied [16, 17], as the alloy rich in iron, even though these oxides are thermodynamically less stable compared with Al2O3 and Cr2O3. With time, underneath iron oxides depleted in Fe and consequently enrich the Ni, Al, and Cr. Even though nickel content is higher in the alloy compared with aluminium, thermodynamically Al2O3 is much stable than NiO. Enrichment of Cr in this region would drive formation of Cr2O3. However, Cr2O3 is less stable than Al2O3. Therefore, Al2O3 would form immediately bellow iron oxides scale. Formation of this aluminium oxide reduced inward diffusion of oxygen to the alloy surface as the ions have larger size compared with other anions [13], making the outward diffusion of iron ions to the oxygen rich regions decrease. The growth of Al2O3 layer caused depletion aluminium content in the alloy below the scale. Consequently, as the Cr and Ni content increased, Cr2O3 and NiO would form, dominated by chromia formation. Formation of Cr2O3 reduces the oxygen pressure on the alloy-scale interface and reduces the possibility of iron oxides to form further [18]. NiO normally formed by outward diffusion of cation [19]. However, relatively larger size of Ni2+ and higher Gibbs free energy for NiO formation a layer of NiO would not be possible to form, except dispersion of NiO particles in a layer of Cr2O3. Reaction or mutual dissolution of these two oxides might produce spinel of NiCr2O4 in this layer as also found in the study of Dudziak etal [20] in Cr-rich FeNi alloys of Hayness 6 [17] but at lower temperature of 750°C. Bensch etal [21] found this type of spinel in GTD111 Ni-base superalloy oxidized at 900°C. This is basically the result of reaction between these two oxides. Depletion of Cr and Ni contents in the alloy underneath Cr and Ni oxides induced the formation of aluminium oxide or alumina layer.

The oxidation resistance of the alloy at 800, 900, and 1000°C can be obtained from the weight gain plot with time, as shown in Figure 10. It is seen that the thickness of the scale increases rapidly at the beginning, indicating that the oxidation rate was very slow at the early stage of oxidation. With longer time, the rate of oxidation, represented by the slope of the curves, decreases with times. The kinetic characteristic of the oxidation can be analyzed from the weight gain data and three possible equations were analyzed by (1) plotting linearly of x with time for linier rate, (2) x with log (t) for logarithmic rate, and (3) x2 to t for parabolic rate. Table 2 shows all possible equations with various R2. It is found that, at all temperatures studied, the oxidation kinetics follow logarithmic rate. This indicates that the oxidation rates of the alloy were very slow, and the oxides formed are relatively compact [20]. The values of the logarithmic rate constants for oxidation at 800, 900, and 1000°C are 5.15E-05 gr.cm−2, 5.57E-05 gr.cm−2, and 6.74E-05 gr.cm−2, respectively, showing that the oxidation rate increased with temperatures.

Table 2: Oxidation rate equations and rate constants of samples at different temperatures.
Figure 10: Weight change of samples oxidized at different temperatures and times.

Kofstad [22] provided the method to evaluate the oxidation rate equation using double-logarithmic plot by plotting log dw/A versus log t. This method can be used to see the fitting of the gradient, or m, of each resulted line where 1, 2, and 3 are for linear, parabolic, or cube, respectively. The lines from plot results, as shown in Figure 11, indicate that no m values fit with the values for linear, parabolic, or cube. Therefore, logarithmic rate behaviour is confirmed for the oxidation behaviour of the 69.5Fe-14Ni-9Al-7.5Cr containing B2 (Fe,Ni)Al precipitates at 800, 900, and 1000°C. The oxides of iron oxide formed at the early stage of oxidation grow very fast and followed by very slow growing of alumina and spinel that acted as compact thin scale of less than 5 μm and significantly reduced the rate of further oxidation.

Figure 11: Double-logarithmic plot results.

4. Conclusion

The homogenized 69.5Fe-14Ni-9Al-7.5Cr alloy had B2 (Fe,Ni)Al precipitates dispersed homogeneously throughout the ferrite-α (α-Fe) matrix. The morphology of B2 was found as bone-type but tended to change into rounded type when heated at higher temperatures and longer times. In addition, after heating, the volume fraction of the B2 precipitates decreased causing the hardness of the alloy reduced. Scale formed on the surface of the samples consisted of oxides of Fe2O3, Fe3O4, α-Al2O3, NiO, Cr2O3, and NiCr2O4, but dominated by α-Al2O3 and Cr2O3. Due to relatively thin layer of oxide scales in order of 2 up to 8 μm, logarithmic type kinetics was observed for all oxidation temperatures. Oxidation rate constant for 800, 900, and 1000°C was, respectively, 5.15E-05 gr.cm−2, 5.57E-05 gr.cm−2, and 6.74E-05 gr.cm−2.

Data Availability

The data used to support the findings of this study are available from the corresponding author upon request.

Conflicts of Interest

The authors declare that they have no conflicts of interest.

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