Corrosion Behaviour of Fe-based and Ni-based Alloys in Wet CO2 Gas with and without Chloride Deposits at 750°C

Corrosion behaviours of Fe-25Cr, Fe-25Cr-2Mn-1Si, Fe-25Cr-20Ni, 310SS, Ni-25Cr, and Ni-25Cr-2Mn-1Si (all in wt. %) with and without NaCl-KCl deposits in Ar-60%CO 2 -20%H 2 O gas at 750 o C were studied. Without salt deposits, Fe-25Cr performed protectively, while Fe-25Cr-20Ni and Ni-25Cr underwent breakaway oxidation with multi-layered scales formed. Adding alloy elements Si + Mn increased the corrosion resistance of all alloys by forming additional Mn-rich oxides and silica. Surface deposits of NaCl-KCl accelerated corrosion, forming porous Fe-rich oxide nodules for Fe-25Cr and thick, porous scales and internal oxidation zones for all other alloys. The protective effect of Si + Mn alloying disappeared in the presence of chlorides. Limited intergranular carbides were observed for all alloys in the gas-only condition. The extent of carburisation increased with the presence of chloride deposits for all Fe-based alloys, but remained unchanged for Ni-based alloys. Corrosion of these alloys at 750 o C is compared with that at 650 o C. The effect of chlorides in volatilising metals at 750 o C is discussed.


Introduction
Reducing carbon emissions is a current global trend in attempting to limit global warming.
Although renewable and clean energy sources are increasingly being utilized, the traditional fossil fuel energy could remain in use for decades due to the enormous energy demand.Oxy-fuel combustion is a potential technology for limiting CO2 emissions in thermal power generation.In this process, coal is burnt in oxygen diluted with recycled flue gas, instead of air, which therefore produces flue gas mainly consisting of CO2 and H2O, allowing easier CO2 collection and sequestration [1][2][3].However, this flue gas is very corrosive at high temperatures.Recent research revealed that both Fe-based and Ni-based Cr-containing alloys with even 25 wt.% Cr are not sufficient to form a protective scale in such an environment [4][5][6][7].However, corrosion resistance could be significantly improved by alloying with Si and/or Mn [8,9].
In addition to the gas environment, solid particles like chloride salts and coal ash are also found to be deposited on combustion equipment surfaces, where they can affect alloy corrosion [10][11][12].Our recent work [8,9] on chloride corrosion of Fe-based and Ni-based alloys in wet CO2 at 650C showed clearly that chloride salts cause severe material degradation due to active oxidation, the volatilisation of metal as chlorides [13,14].This paper examines corrosion behaviour of ferritic alloys, austenitic, and Ni-based alloys coated with chloride salt deposits, comparing them with salt-free surfaces in wet CO2 at 750C.This temperature was chosen to match the likely requirement for advanced ultra-supercritical boilers.The results are compared with those at 650C [8,9] to analyse the temperature effect.
The alloys were cut into approximately 9 x 7 x 2 mm rectangles, and ground to a 1200-grit finish, followed by polishing to a 3-µm finish and finally electropolishing in 15% hydrochloric acid.The purpose of electropolishing was to remove the work-hardened surface zone introduced during sample preparation.Alloy grain boundaries were revealed after electropolishing [8,9].Alloy grains in Fe-25Cr, Fe-25Cr-2Mn-1Si, Fe-25Cr-20Ni, and Ni-25Cr were elongated, with sizes ranging from hundreds of microns to several millimetres, while the 310SS grains were much finer and more uniform, of order tens of microns.
For the gas-only experiment, sample coupons were laid on a ceramic crucible with the upper face exposed to the gas atmosphere.This was done for comparison with the subsequent saltdeposited experiments.The ceramic crucible was placed in a horizontal alumina tube furnace for isothermal oxidation at 750C, and exposed to flowing gases for 300 h.The reaction gas was Ar-60%CO2-20%H2O (vol.%) with a linear gas flow rate of 2 cm/s at 750C.Argon and carbon dioxide inputs were controlled by mass flow controllers.Wet gases were generated by passing the Ar + CO2 mixture heated water to first produce an excess amount of water vapour, which was then condensed inside a distillation column at a temperature set to achieve the required   2  value.
For salt-deposit experiments, NaCl and KCl were dissolved in distilled water in a weight ratio of 1:1.25, in order to obtain the lowest liquidus temperature [9].The solution was sprayed onto preheated specimens which were then dried using a heat gun.The coated samples were weighed after spraying, ensuring salt deposits in the range 2.5 ± 0.5 mg/cm 2 .Sample coupons were laid on a ceramic crucible with coated face exposed to the flowing reaction gas.Coated alloy samples were exposed in the same way as described above for non-coated specimens, for a reaction time of 300 h.In the presence of chlorides (NaCl + KCl), the calculated equilibrium gas partial pressures [15] including   and   and aC are listed in Table 1.

Table 1. Equilibrium minority gas partial pressures (atm) and aC at 750°C, 1 atm
Reacted sample surfaces were analysed by X-ray diffraction (XRD; PANalytical Empyrean with Co Kα radiation).All samples were then cold-mounted in epoxy resin for metallographic preparation.
The salt deposited specimens were prepared using an oil-based lubricant to protect any remaining chlorides.Cross-sections were observed and analysed by a series of techniquesoptical microscopy (OM); Raman microscopy (Renishaw inVia Raman Microscope; wavelength 532 nm with argon laser); field emission scanning electron microscopy (FEI Nova NanoSEM 450) and energy-dispersive X-ray spectroscopy (EDS; Bruker); and transmission electron microscopy (TEM: Philips CM200).Samples for TEM were prepared by focused ion beam (FIB: FEI Nova Nanolab 200) milling with an accelerating voltage of 30 kV.

Gas-only Condition
Figure 1 shows metallographic cross-sections of all alloys after 300 h exposure in wet CO2 gas.
Protective behaviour was also found for Fe-25Cr-2Mn-1Si (Fig. 1 this thin layer produced the results shown in Fig. 2. From the TEM bright field image shows the oxide scale to consist of coarse-grained outer and fine-grained inner layers.As seen from the EDS analysis and phase identifications published previously [16], the scale consists of an outer Mn-rich oxide layer, an intermediate MnCr2O4, and an inner Cr2O3 layer.A thin, continuous SiO2 layer was also found at the scale-alloy interface [16].The Fe-25Cr-20Ni austenitic alloy formed a thick multilayered scale (Fig. 1(c)), consisting of external Fe2O3, intermediate spinel, and inner chromium rich oxides [16].While for 310SS, a protective oxide scale interrupted by small, local Fe-rich oxide nodules was observed (Fig. 1(d)).The TEM-EDS mapping results for the protective scale on 310SS (Fig. 3) revealed an outermost MnCr2O4 layer with Cr2O3 below it, and discrete islands of SiO2 at the oxidation front [16].A multi-layered scale formed on Ni-25Cr (Fig. 1(e)), incorporating a metallic Ni layer within the oxide [17].For Ni-25Cr-2Mn-1Si, a much thinner oxide was observed, and no Ni layer appeared (Fig.

1(f))
. High magnification image revealed that this thin multi-layered scale contains an outer NiO layer, a middle thin Mn-rich oxide layer, and an inner (Cr, Mn)-rich oxide layer (Fig. 4).There were silica precipitates at the reaction front (Fig. 4).In addition to this structure, a thin protective layer was also found to be mainly a chromia layer with a small amount of Mn3O4/MnCr2O4 at its surface and a silica sublayer beneath, occupying about 10% of the whole alloy surface [17].
Mixed oxides with different contrast are seen in Fig. 7(b), which mainly consist of Fe2O3 (Pa in Fig. 7(b)) and Cr2O3 (Pb in Fig. 7(b)), with an inner oxide, most likely spinel and chromia (Pd in Fig. 7(b)) mixed with a small amount of Fe2O3 (Pc in Fig. 7(b)).The area mapping in Fig. 8 reveals that the light contrast oxide is Fe-rich; some Cr is located in the oxide outside the original alloy surface; the inner oxide layer is enriched with Mn; and Si is located at the reaction front.Of note, Si is also enriched at the scale surface (Fig. 8).No signal for K, Na, or Cl was detected.Clearly both external and inner oxide layers are very porous with significant local scale damage.Moreover, the depletion of Fe and the enrichment of Cr confirmed the grey shaded elongated precipitates below the scale-alloy interface to be Cr-rich carbides, as no oxygen enrichment was detected.

Fe-25Cr-20Ni
Images of Fe-25Cr-20Ni after salt affected reaction (Fig. 9(a-b)) show the whole surface was covered by a multi-layered scale with localized deep inner oxide protrusions, as shown in Fig. 9(b).
According to the Raman results, Fe2O3 is located at the external surface (Pa in Fig. 9 An enlarged SEM view of Fig. 9(b) confirms that the external scale consists of Fe-rich oxides (P1 in Fig. 10), (Fe, Ni)-rich oxide (P2 in Fig. 10), and dark contrast Cr2O3 (P3 in Fig. 10).The pitting area contains an IOZ (internal oxidation zone) consisting of large (Fe, Cr, Ni)3O4 spinel (P4 in Fig. 10) and Ni metal (P5 in Fig. 10) particles.A thick, dense chromia band (P6 in Fig. 10) is located at the interface  SEM-EDS mapping of the outer layer (Fig. 12) revealed that the light shaded oxide is Fe-rich, while the dark contrast oxide is Cr2O3.An IOZ was found where the metal and spinel particles are finer than the ones observed on Fe-25Cr-20Ni.In addition, silicon enrichment was seen in the black precipitate area ahead of the main reaction front.Oxide at the main reaction front was rich in Cr (Fig. 12).

Ni-25Cr
Metallographic cross-sections of Ni-25Cr after 300 h exposure (Fig. 13(a, c)) revealed two different reaction morphologies.The first structure is shown in Fig. 13    In addition to the above typical structure, another localised structure (about 10% of the alloy surface) (Fig. 14(c)) was found with external Cr-rich oxide (P4 in Fig. 14(d)), containing dispersed nickelrich particles in its inner region.Some voids containing silicon (P5, P6 and P7 in Fig. 14(d)) were located internally, and chlorine (4.6 wt.%) was detected in one of them (P7 in Fig. 14(d

Discussion
The corrosion behaviour of all these alloys at 650 o C has been reported [8,9], and is compared with that at 750 o C in Table 2.The performance of austenitic and Ni-based alloys after reaction in the absence of salt is similar at the two different temperatures, but ferritic Fe-25Cr behaved differently, protective at 750 o C, but non-protective at 650C.Adding Mn and Si increased corrosion resistance of all three types of alloys in the gas only condition at both temperatures.However, in the presence of chloride deposits, this beneficial effect did not exist and breakaway corrosion occurred (Table 2).The thicknesses of oxide scales for all six alloys after reactions with and without chloride deposits at 750C are compared in Table 3.Without chlorides, the additional Mn and Si significantly reduced the scale thicknesses of Fe-25Cr-20Ni and Ni-25Cr alloys.This alloying effect was not apparent for Fe-25Cr ferritic alloy which in any case performed protectively at this reaction temperature.In comparing the behaviour of 310SS with Fe-25Cr-20Ni, it is recognised that the fine-grained microstructure of the steel is another factor behind its better performance.

Table 3. Corrosion product thicknesses (µm) after 300h in Ar-60CO2-20H2O at 750 o C
In the presence of chloride deposits, the pattern of alloy effects on corrosion resistance at 750 o C is rather different.Adding Mn+Si to Fe-25Cr is worse than ineffectual, leading to a substantial increase in scale thickness.Comparing the two austenitics, 310SS and Fe-25Cr-20Ni, one sees somewhat worse performance for the steel, with a slightly thicker scale and a more uniformly attacked internal zone (Figs. 9 and 11).In contrast, the presence of Mn+Si in the Ni-base alloy leads to some improvement, particularly in largely suppressing internal oxidation.Only 10% of the alloy surface underwent the internal attack, developing voids and precipitates.Oxide scale morphologies formed on the three alloys without Mn or Si are different under the gas-only condition.Unlike its failure at 650C [8,9], Fe-25Cr (Fig. 1(a)) formed a protective chromia layer at 750C.Both Fe-25Cr-20Ni (Fig. 1(c)) and Ni-25Cr formed an IOZ with a multilayered outer scale for the former, and an expelled Ni layer overgrown by a continuous NiO layer for the latter.After adding Mn and Si to the alloys, all three alloys behaved protectively with a thin dense protective scale formed.

Gas only
When reacted under chloride deposits, all alloys experienced more extensive corrosion than those in the gas only condition.Similar to what was reported for 650C, more porous oxides were observed and significant relocation of Cr from the metal substrate to the external oxide scale was identified for both ferritic and austenitic iron-based alloys.However, for the two Ni-based alloys (Figs. 13(a, c), 14(a, c)), porosity within the oxide scales is very small.Localised alloy porosity was found at the reaction front for both austenitic iron-based alloys (Figs.10) and Ni-based alloys (Figs.13(a, c), 14(a, c)).
The oxidation performance of alloys in the gas-only condition is first reviewed and compared with that at 650 o C. Effects of chloride salt deposits and the temperature effect in the presence of chlorides are then examined.

Corrosion of alloys under gas-only condition
To understand the different corrosion behaviour of alloys in the gas-only condition at 650 o C [9] and 750 o C, the critical Cr concentrations for protective chromia formation were calculated from Wagner's equations [18,19].Critical values are listed in Table 4 [8,9,16,17].Here,   (1) is the critical value for external chromia formation rather than its internal precipitation, and   (2) for the maintenance of continued growth of that scale.
The results indicate that the Cr concentration of the alloy is high enough for chromia formation and maintenance for Fe-25Cr, but not for Fe-25Cr-20Ni and Ni-25Cr alloys at either temperature.Comparing the critical Cr values at 650C and 750C in Table 4, there is a significant decrease in both   (1) and   (2) for Fe-25Cr with increasing temperature, due to the temperature sensitivity of alloy Cr diffusion.The prediction is in line with the observed protective oxide formation at the higher temperature.Although the calculation suggests also protective oxide formation for Fe-25Cr at 650C, a thick multilayered scale was grown instead.This was attributed to the formation of internal chromium carbides, leading to a substantially lowered subsurface Cr concentration [9].Clearly, the decrease in the critical Cr concentrations at 750C benefits the formation of a protective oxide scale.The beneficial effects of Si and Mn alloying are significant at both temperatures, promoting protective chromia scale formation for all alloys.These effects have been well explained by forming an additional silica sublayer or silica precipitates at the reaction front which provide a diffusion barrier effect.The fast diffusion and high solubility of Mn in chromia make it possible to form Mn-bearing oxides such as (Cr,Mn)2O3 or Mn and Cr spinel, which further increase the protective value of the scale [20].These alloying effects are demonstrated by the microstructure analyses shown in [8,9] at 650C and in Figs.2-4 at 750C.
The observed Mn-enriched oxide located above the inner (Cr, Mn)-rich oxide layer (Fig. 4) could also be attributed to the diffusion coefficient of Mn (~10 -17 cm 2 s -1 ) [21] higher than that of Cr in Cr2O3 (~10 -30 cm 2 s -1 ) [22] in Cr2O3 at 750 o C. Either MnO or MnCr2O4 is a stable phase which allows the Mn enrichment at the top of the inner layer [23].

Effects of Chloride Deposit
As mentioned at the beginning of the discussion, porous oxides and/or alloy subsurface pores were found at both 650C [8,9] and 750C (Figs. 5, 6(a), 7b, 9b, 10, 11b, 13(b), and 13(d)).This phenomenon can be explained by the so-called active oxidation mechanism [9,[24][25][26][27].In the presence of chlorides (NaCl and KCl), Cl2 can be produced from reaction between salt and water vapour, hence initiating the formation of metal chlorides by the reaction within and beneath the scale of metal or metal oxides with Cl2.These metal chlorides are volatile and can transport outward to be reconverted to oxides at high   2 values near the gas-oxide interface.The Cl2 formed from this conversion will then react with the remaining metal oxide and this reaction cycle will be repeated, continuing active corrosion: where the forward Eq. 2 represents oxide volatilisation, favoured where   2 is high, and the reverse reaction describes chloride re-oxidation and deposition externally where   2 is high.
As shown in Tables 2 and 3, during corrosion under a salt deposit, all alloys underwent breakaway oxidation at 650 o C, while Fe-25Cr and Ni-25Cr-2Mn-1Si exhibited partially protective behaviour at 750 o C, and all other alloys were non-protective.
A protective Cr2O3 scale (Fig. 1(a)) was formed on Fe-25Cr under the gas-only condition at 750C.However, in the presence of chloride deposits, a thin double-layered scale was formed, together with local nodules of thick, porous Fe-rich oxide (Fig. 5(b)).In contrast, at 650 o C, thick multilayered scales formed on Fe-25Cr in both gas-only and under-deposit conditions.As discussed in Section 4.1, the difference is attributed to the faster diffusion of alloy Cr and the consequently enhanced ability to form chromia at 750 o C. Evidently the presence of chloride salt mitigates the improvement in corrosion resistance at the higher temperature.
At the oxide-gas interface,   2 = 1.4 x 10 -7 atm, and equilibrium values of   2 = 7.7 x 10 -21 atm and   2  2 = 2.9 x 10 -17 atm were calculated based on Eqs.(3)(4)(5).It is concluded on this basis that if an external protective chromia layer is able to form before Cl2 is produced, the volatilization would be suppressed.
At the oxide-alloy interface,   2 is controlled by the local Cr/Cr2O3 equilibrium.Approximating aCr ≈ 0.25, a value of   2 = 3.30 x 10 -29 atm is calculated.If the   2 value at the oxide-gas interface (1.0 x 10 -13 atm) is used as the maximum value for the estimation, the resulting   2  2 value is 1.1 x 10 -22 atm, hence the vaporization of this species can be neglected.However,   2 is found to be 1.4 x 10 -4 atm, indicating the occurrence of appreciable vaporization [28][29][30].It is therefore concluded that if Cl2 can enter the chromia scale, volatilisation of the oxide as CrCl2 is very likely to happen, which explains the relocation of Cr to the external oxide scale observed (Figs.6(a)) and the loss of protection of chromia scale at 750C.This conclusion should also apply for Fe-25Cr at 650 o C if chromia forms.
However, no chromia scale formed at 650C, and the porous iron oxide scale formation was attributed to vaporisation of iron-rich oxides [8].If such a process occurred at 750 o C, it would also affect the oxide scale morphology.The possibility of iron oxide volatilisation at this temperature is now considered.
Figure 17 shows the Fe-O-Cl predominance diagram at 750 o C, calculated from FactSage 8.1 [15].A dashed line added on the Fe-O-Cl stability diagram represents the vapour pressure for FeCl2 (g) of 10 −4 atm, above which substantial volatilisation is generally assumed [28][29][30][31].The point marked as a star labelled A on the diagram represents gas phase conditions at the scale surface and is in the  * scale, leading to the predicted outward transport of Cr, as shown by the formation of a thick, porous, outer chromia scale layer (Fig. 13(d)).The observation of an additional thin spalled oxide layer containing varying concentrations of Ni on the top of the thick chromia layer (P4 and P5 in Fig. 13(d)) is attributed to the initial stage of oxidation.
In addition to the porous oxides formed in the presence of chlorides, a considerable amount of external Cr-bearing oxide formed on all six alloys (Figs. 6, 8, 10, 12, 13(d), 14(d)) which is attributed to the vaporisation of chromia by chlorine.This phenomenon was found at both temperatures, but with more Cr transported outward at 750C.For Fe-25Cr ferritic alloy, the amount of Cr carried out was very limited at 650 o C [9], but more Cr was relocated to the external oxide scale at 750 o C (Fig. 6).
For Fe-25Cr-20Ni, the outer Cr-rich oxide was localised at 650C [8], but became continuous although non-uniform at 750 o C (Fig. 10).For Ni-25Cr, the area of the structure with thick chromia oxide was increased from 40% at 650 o C [8] to 60% at 750 o C (Fig. 13(c)).Those differences are attributed to the increased partial pressure (two to ten times the values at 650 o C) of CrCl2 produced at 750 o C, which increases the rate of volatilization.
With the presence of chloride salts (Figs.8,12,14), the beneficial effects of both Mn and Si observed in the gas-only condition (Figs.2-4) disappeared on all alloys at 750 o C, as they did also at 650 o C. It is noteworthy that the oxide layers (Table 3) formed on Fe-25Cr-2Mn-1Si (Fig. 7(b)) and 310SS (Fig. 11(b)) are even thicker than those formed on the Si-and Mn-free alloys (Figs.5(b), 9(b)).An explanation can be found in the volatilisation of the additional SiO2, Mn3O4, and (Mn, Cr)3O4 spinel products: The predicted partial pressures of SiCl4 and MnCl2 can be calculated, by using   2 at Cr/Cr2O3 equilibrium, which yields   4 = 2.3 x 10 -8 atm, and   2 ≫ 1 atm.It is concluded that only Mn could be volatilised at the very low oxygen partial pressure at the Cr/Cr2O3 equilibrium.Moreover, the diffusion rate of Mn is fast in Cr2O3 [22].These two processes explain why Mn could be detected externally.The existence of Cl at the porous region (P7 in Fig. 14(d)) indicates the possible formation of SiCl4 locally, which explains why the protective effect of Si disappeared.

Carburisation
Comparison of the carburisation results (Figs. [15][16] for the gas-only and chloride coated cases, reveals that the extent of Fe-based alloy carburisation is significantly affected by the chloride deposits.Whilst only sparsely distributed intergranular carbides were found after gas-only reaction, some intragranular carbides were also found in the subsurface regions of Fe-based model alloys beneath chloride deposits (Fig. 16(a-c)).
Carbide fractions measured by ImageJ [32] in reacted 310SS seemed very similar in the two cases (gas-only: fv = 0.036 ± 0.001, and chloride-deposited: fv = 0.037 ± 0.001).This reflects the fact that most carbides are intergranular, which come from the carbon originally present in the steel.
Significantly, intragranular carbides developed near the surface in the presence of chlorides (Fig.
The increased carburisation of Fe-based alloys when exposed under chloride deposits reflects the greater carbon permeability of the non-protective scales grown in these conditions.Just as oxidation is faster in the presence of chloride, so too is carbon ingress.
The difference in carburising behaviour between Fe-based and Ni-based alloys is due to significantly different carbon diffusion coefficients in Fe-based (ferritic and austenitic: 3.9 x 10 -7 [33], 1.8 x 10 -8 cm 2 s -1 [34], respectively) and Ni-based alloys (6.6 x 10 -9 cm 2 s -1 [35]).Moreover, as mentioned demonstrated that the volatilisation of Mn is significant, and any silica subscale is damaged by reaction with chlorine.
Only very limited intergranular carburisation was found for all six alloys in the gas-only condition.In the presence of chlorides, the two Ni-based alloys behaved similarly to the gas-only case.
However, some intragranular carbides were also found in all four Fe-based alloys.This difference can be attributed to faster carbon diffusion in iron-based alloys and less accelerated attack on nickel-based alloys by chloride due to less volatile chloride formation.

Statements and Declarations
(b)).Analysis by TEM-EDS mapping

Figure. 9 .
Figure. 9. (a-b) Cross-section of Fe-25Cr-20Ni after 300 h reaction beneath chloride deposits with different magnifications, and (c-h) Raman analysis results for points Pa, Pb, Pc, Pd, Pe, and Pf marked in (b).
u.)Raman shift (cm -1 ) between the alloy and the oxides.Subsurface voids were associated with internal oxides identified by EDS (not shown) within the alloy.

Figure. 11 .Figure. 12 .
Figure.11. (a-b) Cross-section of 310SS after 300 h reaction beneath chloride deposits with different magnifications, and (c-f) Raman analysis results for points Pa, Pb, Pc, Pd, and Pe marked in (b).
(a), which is a multi-layered scale.Further analysis of this structure by SEM-EDS point analysis (Fig.13(b)) identifies the outer layer as NiO, located on top of the original alloy surface.In addition, an IOZ composed of (Ni, Cr) spinel and metallic Ni was found below the surface, with a Cr2O3 band at the reaction front.

Figure. 13 .
Figure. 13.Cross-sectional optical micrographs (a, c) and BSE-SEM cross-section (b, d) of Ni-25Cr after 300 h reaction beneath chlorides deposit with different structures at 750°C, the white dashed line in Fig. 13(a) denotes the original surface of alloy.

1
(a)) to a double layer with nodules for Fe-25Cr in the presence of chloride deposits (Fig.5(b)), volatilisation of Cr2O3 in the presence of chlorine is considered:

Fe2O3
field.This oxide is therefore predicted to occur at the scale surface, as was observed.If chlorine can enter the scale and diffuse inward, the resulting diffusion path is represented schematically by the dotted line A-B.When this path intercepts the dashed line, volatilisation of FeCl2 becomes significant, and oxide is consumed, leaving a porous oxide scale.Such a process accelerates porous oxide formation for Fe-25Cr, as observed in Fig.6(a).

Figure. 17 .FeCl3
Figure.17. Traditional stability diagram and quasi-stability diagram (dashed lines calculated for   2 = 10 -4 atm) of Fe-O-Cl system at 750 o C. "Star A" represents the experimental condition, dotted line AB represents an assumed diffusion path for oxygen and chlorine from the oxide/gas to oxide/metal interface.Over the majority of the surface, however, the iron oxide layer remains intact above a Cr-rich oxide in contact with the alloy.It seems therefore, that conditions are marginal for volatilisation, and

650°C[8, 9] and 750°C[16, 17]
This does not answer the question.Saying that the calculated partial pressures are very different is one thing.Concluding that one species ids more likely to volatilise than the other is correct, but it does not prove that it actually happens.You are saying that Fig 13 proves that it does happen.How does it?What does the reader see in Fig 13d which is physical (visible) evidence that Cr is volatilised but Ni is not?There is an outermost layer of Ni-rich oxide.Commented [YC6R1]: The thick outer layer in figure 13d contains only Cr and O.And the outermost Ni metal should formed at the experiment initial stage.