High-Temperature Corrosion Behavior of Heat-Resistant Alloys

in Atmosphere Containing Water Vapor and Ammonia

SuzueYoneda1✉Emails-yoneda@eng.hokudai.ac.jp

ShigenariHayashi1

Division of Materials Science and Engineering, Faculty of EngineeringHokkaido UniversityN13 W8, Kita-ku060-8628SapporoJapan

Suzue Yoneda and Shigenari Hayashi

Division of Materials Science and Engineering, Faculty of Engineering, Hokkaido University, N13 W8, Kita-ku, Sapporo 060-8628, Japan

Email: s-yoneda@eng.hokudai.ac.jp

Keywords:

high-temperature corrosion

ammonia

nitridation

Cr₂O₃-forming alloy

Abstract

Various chromia-forming alloys were exposed to atmospheres of Ar-10%H₂O and Ar-10%H₂O-2.5%NH₃ at 800°C. Protective Cr₂O₃ scale formed on SUS310S and Inconel 718 in both atmospheres, and no nitride formation was observed. However, the high-temperature corrosion resistance of Alloy 800H substantially decreased in Ar-10%H₂O-2.5%NH₃ compared with in Ar-10%H₂O, and internal oxide and thick internal nitride layers formed, despite the initial formation of Cr₂O₃ scale. These results suggest that the nitrogen permeability through Cr₂O₃ scale may depend on the alloy substrate. In Inconel 718, CrNbO₄ layers formed beneath the Cr₂O₃ scale, indicating that this layer may also act as a nitrogen diffusion barrier. In Alloy 800H, because the initially formed Cr₂O₃ scale did not suppress the nitrogen diffusion, an internal nitride layer formed beneath the Cr₂O₃ scale, resulting in the breakaway of the Cr₂O₃ scale and severe internal oxidation and nitridation.

1. Introduction

To achieve carbon neutrality, hydrogen is expected to become an alternative to fossil fuels. Ammonia is attracting attention as a hydrogen energy carrier because it is easier to transport and store than hydrogen and can be used directly [1–6]. Technological development for direct ammonia use is underway in many fields, including thermal power generation, gas turbines, and marine engines [2–3, 5–7]. These studies often focus on NOx reduction and improving combustion, which are major issues in ammonia combustion, through improvements in burners, combustors, and combustion methods. However, the high-temperature corrosion resistance of heat-resistant alloys used in ammonia combustion environments has rarely been evaluated. During ammonia combustion, heat-resistant alloys are expected to be exposed to an environment with high nitrogen potential because of nitrogen formation and the presence of unburned ammonia. Additionally, water vapor concentration increases in mixed combustion of coal and ammonia compared with that coal combustion alone [8]. Thus, the alloys are expected to be exposed to more severe environments. Cr and Al have strong affinities with nitrogen and form internal nitrides, which may result in poor high-temperature corrosion resistance because of the decreased Cr and Al content in the matrix in atmospheres with high nitrogen potential. The formation of internal nitrides may also decrease toughness. Therefore, understanding high-temperature corrosion behavior under ammonia combustion is important for high-temperature corrosion resistance and mechanical properties.

There have been several studies on high-temperature nitridation of alloys [9–14]. In a study on high-temperature corrosion in N₂-5%H₂ and 100%NH₃ atmospheres, Tjokro and Young [10] reported that, despite the nitrogen potential being lower in 100%NH₃, the parabolic rate constant was higher. For alloys with a high Fe content, the parabolic rate constant k_p in 100%NH₃ was 30 to 70 times greater than in N₂-5%H₂, which was attributed to higher nitrogen activity due to the catalytic dissociation of NH₃ at metal surface. Sand et al. [9] conducted high-temperature corrosion tests on various commercial alloys in N₂-5%H₂-(14, 45)ppmH₂O atmospheres and reported that internal nitridation was significantly suppressed in N₂-5%H₂-45ppmH₂O, in which a chromia scale formed, compared with N₂-5%H₂-14ppmH₂O, in which no chromia scale formed. The water vapor concentration increases during ammonia combustion; therefore, heat-resistant alloys are expected to be exposed to atmospheres in which a chromia scale forms and ammonia is present. However, no studies have been published on high-temperature corrosion in these atmospheres.

In this study, the high-temperature corrosion behavior of various chromia-forming alloys in atmospheres containing ammonia and water vapor was investigated to understand the high-temperature corrosion behavior during ammonia combustion.

2. Experimental

Commercial SUS310S, Alloy 800H, and Inconel 718 were used. The chemical compositions of alloys are given in Table 1. Samples approximately 1-mm-thick were cut from a φ12 mm bar and polished with 4000 grit SiC paper. Samples were finished with a 3 µm diamond paste, followed by ultrasonic cleaning in acetone.

Table 1
Chemical composition of alloys
													wt%
	Fe	Ni	Cr	C	Si	Mn	P	Al	Ti	Mo	Co	Nb	Cu
SUS310S	Bal.	19.42	25.65	0.04	0.62	1.25	0.25	-	-	-	-	-	-
Alloy 800H	Bal.	31.16	19.57	0.068	0.51	0.94	0.023	0.54	0.56	0.23	0.09	-	0.05
Inconel 718	18.32	Bal.	18.45	0.04	0.08	0.06	0.009	0.59	0.99	2.91	0.34	5.29	0.05

Figure 1 shows a schematic of the corrosion test equipment. The samples were placed in a horizontal furnace, which was evacuated and purged with Ar gas (99.9999%). The samples were heated at a rate of 10°C/min to 800°C in Ar gas, and Ar gas was immediately switched to Ar-10%H₂O or Ar-10%H₂O-2.5%NH₃. Ar gas was bubbled into distilled water at about 55°C or 62°C, then cooled to 46.4°C or 52°C to obtain the Ar-10%H₂O or Ar-13.5%H₂O gas mixture, respectively. Ar-10%NH₃ gas was added to Ar-13.5%H₂O to obtain the Ar-10%H₂O-2.5%NH₃ gas mixture. The corrosion test was conducted for 100 h. After the corrosion test, samples were analyzed by optical microscope, field-emission scanning electron microscopy (FE-SEM), scanning transmission electron microscopy (STEM), electron probe micro analyzer (EPMA), and energy-dispersive X-ray spectroscopy (EDS). FactSage 8.1 was used for thermodynamic calculations.

Fig. 1

Schematic of the horizontal furnace

3. Results and Discussion

3.1 Oxidation and Corrosion Kinetics

Figure 2 shows the mass gain curves in (a) Ar-10%H₂O and (b) Ar-10%H₂O -2.5%NH₃ at 800°C over 100 h. In Ar-10%H₂O without NH₃, the oxidation mass gain of SUS310S and Inconel 718 gradually increased, and the mass gain was smaller for Inconel 718 than for SUS310S. The mass gain of Alloy 800H was similar to that of SUS310S up to after 25 h of oxidation but started to increase rapidly and became approximately twice that of SUS310S at 100 h. In Ar-10%H₂O-2.5%NH₃, there was no significant difference in the mass gain of SUS310S and Inconel 718 compared with Ar-10%H₂O. The mass gain of Alloy 800H increased rapidly after corrosion for 4 h and became one order of magnitude greater than that in Ar-10%H₂O, indicating that NH₃ decreased the high-temperature corrosion resistance of Alloy 800.

Fig. 2

Oxidation and corrosion kinetics of SUS310S, Alloy 800H, and Inconel 718 at 800°C in

(a) Ar-10%H₂O, and (b) Ar-10%H₂O-2.5%NH₃

3.2 Cross-sectional Microstructures

Figure 3 shows cross-sectional FE-SEM images of each alloy after oxidation in Ar-10%H₂O for 100 h. The oxide scale formed on SUS310S was approximately 2 µm thick, and EDS revealed that the oxide scale mainly consisted of Cr₂O₃. (Mn,Fe,Cr)₃O₄ spinel was also partially formed above the Cr₂O₃ scale. A semi-continuous SiO₂ layer was observed at the scale/alloy interface and in the alloy. Inconel 718 also formed Cr₂O₃ scale that was thinner than that on SUS301S, consistent with the smaller mass gain for Inconel 718. Internal oxides of Al and Ti were observed below the Cr₂O₃ scale, and these oxides tended to be formed along the alloy/δ-Ni₃Nb interface. In Alloy 800H, Cr₂O₃ scale, semi-continuous SiO₂, and nodule formation were observed. The nodules consisted of outer (Fe,Ni)₃O₄ and inner (Cr,Fe,Ni)₃O₄ spinel. The inner (Cr,Fe,Ni)₃O₄ spinel also contained a small amount of Al and Mn. EDS analysis showed that Cr₂O₃ scale contained several percent of Fe and Ni, and (Fe,Ni)₃O₄ was formed above the Cr₂O₃ oxide scale. These results indicated that Fe, Mn, and Ni cations diffused through the Cr₂O₃ oxide formed on Alloy 800H. Nodule formation began after oxidation for 25 h, and the area fraction of nodules occupying the alloy surface and the thickness of the nodules increased with exposure time, resulting in a rapid increase in the oxidation mass gain of Alloy 800.

Fig. 3

Cross-sectional FE-SEM images of each alloy after oxidation at 800°C in Ar-10%H₂O for 100 h.

(a) SUS310S, (b) Alloy 800H, and (c) Inconel 718

Figure 4 shows cross-sectional FE-SEM images of each alloy after corrosion in Ar-10%H₂O-2.5%NH₃ for 100 h. The oxide scales formed on SUS310S and Inconel 718 were similar to those formed in Ar-10%H₂O. No nitride formation was observed beneath the oxide scale in these alloys. By contrast, a Cr-rich oxide scale formed on Alloy 800H in some areas, but most of the area was covered by a porous oxide scale with an internal oxide layer. A thick internal nitride layer was also formed below the internal oxide layer. The internal oxide and nitride layers formed after corrosion for 4 h, and the thickness of these layers increased with exposure time. A metallic phase composed of Ni and Fe was observed on the surface, which may have been formed by extrusion of the matrix because of the formation of internal oxides [15–17]. Figure 5 shows STEM images of the area near the internal oxidation zone (IOZ)/internal nitridation zone (INZ) interface and near the tip of the INZ formed in Alloy 800H after corrosion for 100 h. STEM-EDS analysis revealed that (Cr,Fe)₃O₄ spinel was formed near the internal oxidation/nitridation interface. The internal nitrides consisted of Si and Mn granular nitrides less than 1 µm in size and fine needle-like Cr₂N (Fig. 5(a)). A study on nitriding of steel containing Si and Mn reported MnSiN₂ formation in the steel [18]. Based on the previous study and the present STEM-EDS results (Table 2), the Si and Mn nitrides may have been MnSiN₂. The formation of needle-like Cr₂N was also observed near the tip of INZ, and granular CrN was formed at the grain boundary (Fig. 5(b)). The amount of MnSiN₂ decreased and the size of needle-like Cr₂N increased toward the alloy substrate.

Fig. 4

Cross-sectional FE-SEM images of each alloy after corrosion at 800°C in Ar-10%H₂O-2.5%NH₃ for 100 h. (a) SUS310S, (b) Alloy 800H, and (c) Inconel 718

Fig. 5

Cross-sectional STEM images of Alloy 800H after corrosion at 800°C in Ar-10%H₂O-2.5%NH₃ for 100 h. (a) Near the IOZ/ INZ interface, and (b) near the tip of the INZ

Table 2
STEM-EDS point analysis of the internal oxide and nitride formed in Alloy 800H after corrosion at 800°C in Ar-10%H₂O-2.5%NH₃ for 100 h. Figure 5 shows the analysis points.
									at.%
	O	N	Si	Ti	Cr	Mn	Fe	Ni	Al
1	-	27.1	1.3	0.0	71.3	0.0	0.2	0.1	-
2	10.9	20.8	31.0	0.1	8.7	24.4	0.7	0.1	3.2
3	-	44.7	1.3	0.0	55.6	0.0	0.3	0.0	0.2
4	-	28.1	1.5	0.6	51.7	0.2	10.3	6.5	2.0

3.3 High-temperature Oxidation Resistance of Alloys in Ar-10%H₂O

A protective Cr₂O₃ scale was formed on SUS310S and Inconel 718 after oxidation for 100 h in Ar-10%H₂O, but breakaway of Cr₂O₃ scale and nodule formation were observed in Alloy 800 (Fig. 3). On SUS310S and Alloy 800H, a semi-continuous SiO₂ layer was formed beneath the Cr₂O₃ scale as mentioned in Section 3.2. Si improves the high-temperature oxidation resistance of alloys [19–22]. This SiO₂ layer acts as a barrier to the outward diffusion of Cr or Fe, decreasing the growth rate of the Cr₂O₃ scale [22]. Figure 6 shows the STEM cross section and STEM-EDS mapping of (a) SUS310S and (b) Alloy 800H after oxidation for 25 h in Ar-10%H₂O. The SiO₂ layer formed on Alloy 800H was more continuous, and the Cr₂O₃ scale was thinner than those on SUS310S. However, some areas also showed the start of nodule formation. The Cr₂O₃ scale that formed on Alloy 800H could not prevent the outward diffusion of Fe and Ni, indicating that this Cr₂O₃ scale was less protective than that formed on SUS310S. Furthermore, the Cr content in Alloy 800H is lower; therefore, although Cr₂O₃ formed at the early stage of oxidation, it could not be maintained for longer, resulting in breakaway. The Cr content of Inconel 718 was lower than that of Alloy 800H, but a thin Cr₂O₃ scale remained during the oxidation for 100 h. Kuo et al. [23] reported that Nb addition to Ni-Fe-Cr alloys strongly promotes Cr₂O₃ scale formation because CrNbO₄ acts as an inward diffusion barrier for oxygen. Inconel 718 contains approximately 5% Nb, and STEM-EDS mapping after oxidation for 25 h revealed the enrichment of Cr, Nb and O at the same place in some areas (Fig. 6(c)). Thus, Nb in Inconel 718 may have promoted the Cr₂O₃ scale formation. These results suggest that nitridation may have occurred in Alloy 800H because of the poor long-term performance of the Cr₂O₃ scale.

Fig. 6

Cross-sectional STEM image and STEM-EDS mapping of (a) SUS310S, (b) Alloy 800H and (c) Inconel 718 after oxidation at 800°C in Ar-10%H₂O for 25 h

3.4 Relationship between Cr₂O₃ Scale Formation and High-temperature Nitridation

The oxygen and nitrogen potentials used in the present study were

$\:{P}_{{\text{O}}_{2}}=2.0\:\times\:\:{10}^{-5}$

atm in Ar-10%H₂O, and

$\:{P}_{{\text{O}}_{2}}=2.12\:\times\:\:{10}^{-18}$

$\:{P}_{{\text{N}}_{2}}=1.22\:\times\:\:{10}^{-2}$

, and

$\:{P}_{{\text{N}\text{H}}_{3}}=2.91\:\times\:\:{10}^{-7}$

atm in Ar-10%H₂O-2.5%NH₃. Figure 7 shows the phase-stability diagrams of Alloy 800H and Inconel 718 at 800°C calculated by FactSage 8.1. For the calculations, the Cr activities of each alloy determined using FactSage were used (0.290 for Alloy 800H and 0.406 for Inconel 718). The Cr activity for Inconel 718 was derived from the composition of the matrix, excluding δ-Ni₃Nb, as determined by compositional analysis. The present potential is shown in Fig. 7. Cr₂O₃ can be formed on both alloys. Because the oxide is more stable than the nitride, the oxides form on the alloy surface and the nitrides form below the oxide. If the nitrogen were not to penetrate through Cr₂O₃ scale and Cr₂O₃ acted as a diffusion barrier for nitrogen, nitridation would be suppressed. However, nitrides form below Cr₂O₃ scale [9], which indicates that nitrogen diffuses through Cr₂O₃ scale. Assuming that nitrogen diffuses easily through Cr₂O₃ scale and the nitrogen potential at the Cr₂O₃/alloy interface is similar to that at the gas/Cr₂O₃ interface, the Gibbs free energy change for Cr₂N formation shown in Eq. 1 was negative in all alloys. Therefore, Cr₂N can be formed thermodynamically. For the ΔG calculations, the Cr activity of each alloy calculated using FactSage was used, and the activity of Cr₂N was assumed to be 1. Figure 8 shows the cross-section and EPMA mapping of Alloy 800H after 4 h of corrosion. Cr-rich oxide still formed, and nitride formation was observed below the Cr-rich oxide scale, which indicated nitrogen diffusion through the Cr-rich oxide scale and sufficiently higher nitrogen potential at the scale/alloy interface. Thus, the diffusion path can be expressed by the red curve shown in Fig. 7(a). In contrast, nitridation did not occur in SUS310S and Inconel 718, which suggests that nitrogen potential decreased toward at the scale/alloy interface and became lower than the equilibrium nitrogen potential between Cr and Cr₂N. Thus, nitrogen permeability of the Cr₂O₃ scale may depend on the Cr₂O₃ scale formed on different alloys, and the Cr₂O₃ scale formed on SUS310S and Inconel 718 may act as a diffusion barrier for nitrogen. The differences of nitrogen permeability could also explain why nitridation was suppressed in SUS310S and Inconel 718. The grain sizes of Cr-rich oxide scale formed on Alloy 800H and Inconel 718 seemed to be larger than that formed on SUS310 (Fig. 6), suggesting that it is difficult to evaluate the difference in nitrogen permeability from the microstructure and that other contributing factors should be considered.

Fig. 7

Phase stability for the Cr-O-N system at 800°C

(a) Alloy 800H (a_Cr = 0.290) and (b) Inconel 718 (a_Cr = 0.406)

Fig. 8

Cross-section and EPMA mapping of Alloy 800H after 4 h of corrosion

3.5 High-temperature Nitridation Behavior of Alloy 800H in Ar-10%H₂O-2.5%NH₃

Internal MnSiN₂ and Cr₂N were formed near the IOZ/INZ interface (Fig. 5). Only Cr nitrides were formed near the tip of the INZ. These results indicate that the Cr nitride was more stable than MnSiN₂. The standard Gibbs free energy changes, ΔG⁰, of Cr₂N and MnSiN₂ are expressed by Eqs. 1–4.

$\:4\text{C}\text{r}+\:{\text{N}}_{2}={2\text{C}\text{r}}_{2}\text{N}$	(1)
$\:\varDelta\:{G}^{0}=\text{R}T\text{l}\text{n}\frac{{{a}_{\text{C}\text{r}}}^{4}{P}_{{\text{N}}_{2}}}{{{a}_{{\text{C}\text{r}}_{2}\text{N}}}^{2}}$	(2)
$\:\text{S}\text{i}+\text{M}\text{n}+\:{\text{N}}_{2}=\text{M}\text{n}\text{S}\text{i}{\text{N}}_{2}$	(3)
$\:\varDelta\:{G}^{0}=\text{R}T\text{l}\text{n}\frac{{a}_{\text{S}\text{i}}{a}_{\text{M}\text{n}}{P}_{{\text{N}}_{2}}}{{a}_{\text{M}\text{n}\text{S}\text{i}{\text{N}}_{2}}}$	(4)

Figure 9 shows the Ellingham diagram of Cr₂N and SiMnN₂ calculated by FactSage 8.1. When the activities of Cr, Si, and Mn were set to 1, MnSiN₂ was more stable than Cr₂N (thin line). However, the activities should be less than 1 because an alloy was used. Considering the activity, the stability was reversed and Cr₂N was more stable in Alloy 800H (thick line), supporting the experimental results. The change in the thickness of the IOZ and INZ is shown in Fig. 10. Both the IOZ and INZ grew thicker parabolically, and the parabolic rate constant was 2.4 × 10^− 11 and 2.5 × 10^− 10 cm²/s, respectively. The growth of the INZ was one order of magnitude faster than that of the IOZ, which indicates that nitrogen diffuses faster than oxygen in Alloy 800H. Generally, when the alloys are exposed to two oxidants, such as oxygen and nitrogen, new oxides are formed at the internal oxide/nitride interface by the oxidation of nitrides because oxides are more stable than nitrides [24]. Thus, the growth of the IOZ occurs via the oxidation of nitrides. In this reaction, nitrogen is released, which diffuses inward and forms new nitrides at the IOZ/alloy interface. Figure 11 shows EPMA mapping at the IOZ/INZ interface of Alloy 800H after corrosion for 100 h. Si enrichment was observed in the granular oxides in the IOZ, and their shape and size were similar to the Mn-Si nitrides in the INZ. Therefore, the IOZ likely grew through the oxidation of nitrides under the present conditions. Figure 12 shows a schematic of the proposed high-temperature corrosion behavior of Alloy 800H in Ar-10%H₂O-2.5%NH₃ at 800°C based on the present results. The Cr₂O₃ scale forms initially but cannot suppress the nitrogen diffusion, resulting in internal nitride formation beneath the Cr₂O₃ scale. Cr consumption increases because of the Cr nitride formation, which leads to breakaway of the Cr₂O₃ scale. Once breakaway occurs, less protective Fe-rich spinel oxide scale forms, and the inward diffusion flux of oxygen and nitrogen increases, resulting in the formation of the IOZ and MnSiN₂.

Fig. 9

Ellingham diagram of Cr₂N and SiMnN₂

Fig. 10

Change in IOZ and INZ thickness

Fig. 11

EPMA mapping at the IOZ/INZ interface of Alloy 800H after corrosion for 100 h

Fig. 12

Schematic of high-temperature corrosion behavior of Alloy 800H in Ar-10%H₂O-2.5%NH₃ at 800°C

4. Conclusion

High-temperature corrosion behavior of various chromia-forming alloys in ammonia- and water vapor-containing atmospheres was investigated. The results can be summarized as follows:

Ammonia did not affect the high-temperature corrosion resistance of SUS310S and Inconel 718, and no nitride formation was observed. However, the high-temperature corrosion resistance of Alloy 800H decreased substantially in an ammonia-containing atmosphere, and severe internal oxidation and nitridation occurred.

The severe internal oxidation and nitridation in Alloy 800H may be attributed to the poor long-term stability of the Cr₂O₃ scale and to differences in the permeability of nitrogen through Cr₂O₃.

In Alloy 800H, the initially formed Cr₂O₃ scale did not suppress the nitrogen diffusion, resulting in internal nitride formation beneath the Cr₂O₃ scale and breakaway of the Cr₂O₃ scale.

Acknowledgement

This work was financially supported by the Kubota Young Researcher Support Program and the “Steel Carbon Neutral Research Grant” from the Iron and Steel Institute of Japan.

Author Contribution

S.Y. performed all experiments and analysis and wrote the first draft of manuscript. S.H. commented and reviewed the manuscript.

Date Availability

No datasets were generated or analyzed during the current study.

Declarations

Conflict of interest

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

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