Slow Transition from Protective to Breakaway Oxidation of Haynes 214 Foil at High Temperature
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- Young, D.J., Chyrkin, A., He, J. et al. Oxid Met (2013) 79: 405. doi:10.1007/s11085-013-9364-4
The oxidation behaviour of thin sheet specimens of the alumina forming nickel base alloy 214 in the temperature range 1,100–1,200 °C is described. Rapid transient oxidation produces a spinel oxide layer which then stops growing, as a protective alumina layer forms beneath. The slow growth of this alumina ceases when the alloy aluminium content is exhausted. Subsequent formation of an innermost chromia layer signals an increase in oxygen activity at the scale-alloy interface. The abnormally slow growth of this layer extends the alloy lifetime. Examination of individual layer growth processes revealed a complex time dependence of spinel composition as a result of Cr evaporation, and dissolution of alumina in the innermost chromia.
KeywordsNickel base alloyBreakaway oxidationAlloy 214AluminaCr-evaporationTransient spinel
Nickel base high temperature alloys are commonly used as construction materials for components operating at high temperatures in oxidizing environments. For their protection against oxidation attack, most of these alloys rely on the formation of a slowly growing chromium oxide layer on the alloy surface during high temperature service. From the viewpoint of oxidation resistance, the maximum operating temperature of such alloys is limited to approximately 1,000 °C, the exact temperature depending on the actual component design, service environment, required lifetime, etc. . Beyond that temperature, rapid oxidation attack may occur as a result of enhanced growth rates of the chromia surface scales in combination with the formation of volatile CrO3(g) and CrO2(OH)2(g), the latter being especially important in gases containing substantial amounts of water vapour [2–5].
Far better resistance against oxidation attack can be obtained from surface scales consisting mainly of aluminium oxide. There are a number of alumina forming nickel base materials which form such very slowly growing surface scales. Alloys of the NiCrAlY or NiCoCrAlY type, which typically contain e.g. 15–25 mass% Cr and 6–12 % Al, are, however, mainly used as coatings, e.g. on nickel base superalloy components in aero engines and industrial gas turbines [6, 7]. They are not suitable as construction materials mainly due to their inherent brittleness at low temperatures and the lack of formability into semi finished products such as sheets, bars and tubes. Some of the newer types of nickel base superalloys such as CXMSX-4 or Rene N5 [8, 9] may to a first approximation also be considered as alumina forming materials. However, these materials can not be processed as wrought alloys, mainly due to their high concentrations of γ′ stabilizing alloying elements such as Al, Ti and/or Ta.
Only a few wrought nickel base alloys are commercially available with aluminium levels sufficient to form alumina base surface scales during high temperature exposure. Examples of such materials are Nicrofer 6025 HT, Haynes 224 and Haynes 214 [10, 11], the last having a higher aluminium content than the others. Compared to alumina forming FeCrAl type alloys, the nickel base materials possess substantially higher creep strength, which makes them suitable as construction materials for thin walled components such as heat exchangers, honeycomb structures and catalyst carriers [12, 13].
In spite of their superior oxidation resistance, the lifetime of components manufactured from wrought alumina forming alloys is limited by oxidation during long term service at temperatures at and above about 1,100 °C [14, 15]. The reason is that during high temperature exposure, aluminium is depleted in the alloy matrix due to its selective oxidation into the protective scale. The depletion can be made worse by repeated scale spallation and re-growth induced by thermal cycling. If the remaining aluminium content is decreased beneath a critical concentration, the alloy can no longer re-form the protective scale, and breakaway oxidation—the formation of rapidly growing base metal oxides—results [11, 16]. The time at which this breakaway oxidation commences cannot be predicted from the results of commonly available non-destructive tests. Prediction requires a comparison of the scale growth rate and spallation with the available Al-reservoir in the respective component [14, 15, 17].
In the present paper, the oxidation induced lifetime of the alumina forming nickel base wrought alloy 214 has been investigated in the temperature range 1,100–1,200 °C. A number of oxidation results are available in the literature for this material [10, 18]. In a few cases breakaway times were estimated and compared with those obtained for FeCrAl type materials . However, as far as is known to the authors, detailed microstructural studies of the surface scales, the kinetics of aluminium depletion and the mechanisms of breakaway are not available. For this reason, the emphasis of the present study is on evaluation of the oxidation kinetics, oxide scale composition and morphology, and changes in bulk alloy composition with time in the temperature range 1,100–1,200 °C. The studies concentrate on thin sections, for which breakaway is expected to occur before substantial oxide scale spallation [14, 19].
Chemical composition in weight and atomic percent of alloy Haynes 214 analyzed by ICP-OES
The oxidized specimens were electroplated with nickel and metallographically sectioned. The resulting cross-sections were characterized by optical metallography and scanning electron microscopy (SEM) with energy dispersive X-ray analysis (EDX) and wavelength dispersive X-ray analysis (WDX).
For a more detailed investigation of the phase composition of the oxide scales formed on the alloy surface, the 1,150 °C specimens were analyzed by X-ray diffraction (XRD) using Cu Kα radiation in the D4 ENDEAVOR set-up from BRUKER AXS.
Results and Discussion
Protective Scaling of Alloy 214
Transition from Alumina to Chromia Formation
At 1,150 °C, oxidation is considerably faster (Fig. 1), and leads to complete consumption of the alloy aluminum content (Fig. 6b). All aluminium was removed from the alloy foil after about 800 h of reaction, and alumina scaling ceased. At 1,200 °C aluminium is completely depleted in half the time, approximately 400 h (Fig. 6c). As aluminium is no longer present in the alloy, chromium is expected to oxidize. This effect is indeed observed and is most pronounced at the highest oxidation temperature, resulting in a decrease in remnant alloy chromium concentration after 400 h (Fig. 6c).
Chemical composition (in at%) of oxide layers in Fig. 12d measured by EDX
The pseudo steady-state situation in which chromia forms at the alumina-alloy interface whilst the prior alumina is consumed is important. This stage of reaction is still slow (Fig. 1b), and thus represents an effective extension of alloy life beyond the point where its aluminum content is exhausted.
Chromia formation beneath an external alumina scale has been previously observed during breakaway studies of FeCrAl-base alloys [15, 35–37]. At temperatures in the range 1,100–1,200 °C, exhaustion of the aluminum reservoir in thick specimens of a few mm resulted in formation of rapidly growing Fe-rich oxides. However, in the case of thin specimens, e.g. 50–100 μm thick, this event was preceded by a “pseudo protective” period in which subscale chromia formation occurred. Evans and Taylor  studied breakaway processes in plasma sprayed Ni(Co)CrAlY coatings at 1,100 °C. In isolated particles, the initially formed alumina layer was destroyed by breakaway oxidation, in which the first stage of the process was subscale chromia formation. However, as far as is known to the authors, the process of alumina dissolution in the subscale chromia has not yet been described.
The resulting chromia volatilization rates are excessive at high temperatures in flowing gases [2, 3, 5, 40, 41]. Whereas chromia alone provides a scale which is inadequate for high temperature service, in combination with a prior alumina layer it serves to extend alloy section lifetime substantially. It is desirable to be able to predict the duration of this extension.
The breakaway at the specimen corners during exposure at 1,200 °C was first observed after 300 h exposure, i.e. when the Al content in the bulk specimen was virtually completely exhausted. Apparently, the complete exhaustion results at the specimen corners in formation of rapidly growing mixed oxides rather than subscale chromia formation observed on the flat areas of the specimen.
Although the onset of breakaway at specimen corners is accelerated by mechanical damage to the alumina layer because of out of plane tensile stresses , the alumina layer is also being consumed on flat specimen surfaces, as seen in Fig. 9b. Two processes are involved: dissolution of alumina into the slow growing chromia layer, and conversion to Ni-rich spinel at the spinel–alumina interface.
At temperatures above 900 °C, the miscibility of Al2O3 and Cr2O3 is complete [34, 44]. It is clear from the analysis results in Table 2 that aluminum dissolved in the chromia, but no evidence for bulk diffusion of Cr into the alumina was found.
It is apparent from the micrograph of Fig. 12 that attack on the alumina layer is most rapid at its grain boundaries. The columnar grain structure facilitates penetration, and this may be more important to layer breakdown than a uniform dissolution process. In this event, predicting the life expectancy of the alumina layer becomes difficult.
(for which ∆G° = −242100 − 74.6 T J mol−1) to occur, nickel must diffuse from the underlying alloy through both the chromia and alumina layers. Traces of nickel were detected in the chromia by EDS (Table 2), but not in the alumina. However, the SNMS profiles in Fig. 11b confirm penetration of nickel into the alumina after breakaway. It seems likely that nickel diffusion in Al2O3 is, like that of aluminum and oxygen at these temperatures, a grain boundary process . Overall concentrations of nickel in the alumina layer would in this case be small.
It is concluded therefore that the remnant alumina layer transmits oxygen, allowing chromia to form. The mechanism is presumably one of grain boundary diffusion . The unusually slow growth rate of the resulting chromia subscale results from limited oxygen availability beneath the alumina layer.
Overall diffusion of nickel through both chromia and alumina layers is driven by the gradients of both nickel and oxygen activities. It seems likely that the increase in aO within the alumina layer accompanying the onset of chromia growth accounts for the effective increase in Al2O3 nickel permeability. Further spinel growth at the alumina-spinel interface via Eq. (5) then becomes possible.
It is very likely that the decreasing Cr/Al-ratio in the spinel in the early time period is related to formation of volatile Cr-species according to Reaction (9). This loss in chromium may in fact result in a thinning of the spinel layer, which is, however, not clearly experimentally observed (Fig. 3).
Interesting in this context is the presence of substantial porosity in the spinel (Fig. 2a). The time and temperature dependence of this porosity was difficult to quantify but it seems not unlikely that void formation is affected by Cr-volatilization. Pint et al.  found more voids in the outer spinel layer grown on Haynes 214 in wet air at 1,100 °C than when reacted with laboratory air. Additionally, porosity is known to originate from reaction between various oxides in the transient oxidation stage [24, 28, 49].
After prolonged air exposure, the outer spinel layer is thus expected to consist exclusively of NiAl2O4 as found during 1,100 °C exposure. Qualitatively similar behavior is observed within the first 800 h of oxidation at 1,150 °C and in the first 300 h at 1,200 °C (Fig. 18). However, the Cr concentration in the spinel is seen in Fig. 18 to increase again after 800 h at 1,150 °C, and after approximately 300 h exposure at 1,200 °C. A matching decrease in Al concentration occurs simultaneously, whereas the nickel concentration remains constant and the spinel layer thickens. It should be mentioned that the relative error of the EDS analysis in Fig. 18 was approximately 10 %, i.e. the minimum of Cr concentration e.g. at 1,150 °C may seem doubtful as the variation of the values may fall within the measurement error. However, the reproducibility of the EDS analysis was very high for all specimens which at least confirms the qualitative shapes of the curves plotted in Fig. 18 although the given absolute concentration values may contain a systematic error.
A decrease in Cr level in the spinel layer with increasing time is also confirmed independently by the SNMS profiles (Fig. 11) after 1,150 °C oxidation.
The increase in chromium content of the spinel is obviously correlated with the formation of a (Cr,Al)2O3 layer beneath the alumina scale once the Al reservoir in the specimen is exhausted. As the alumina layer loses its protective properties, chromium starts to diffuse outwards to be incorporated into the spinel. The driving force for outward diffusion of chromium to the outer spinel layer is, of course, provided by the activity gradients in oxygen and chromium. What changes after longer exposure times is the increase in oxygen potential within and beneath the alumina layer as a result of the complete removal of Al from the underlying alloy. Evidently, this increases the solubility and/or mobility of chromium within the alumina, allowing its passage through the layer. At the same time, the alumina layer continues to shrink by interaction with the underlying chromia. Eventually, the scale should become duplex consisting of an outer Ni(Cr,Al)2O4 spinel layer and an inner mixed oxide (Cr,Al)2O3. At this point, the alloy would behave simply as a chromia-former, and lose utility at such high temperatures.
Summary and Concluding Remarks
Thin foils of Haynes 214 alloy have been oxidized in air at 1,100–1,200 °C. In the protective stage of oxidation, a duplex scale morphology was observed. An outer Ni(Al,Cr)2O4 spinel layer grows readily in the first few hours of reaction, and remains intact while a protective alumina layer thickens beneath it.
Continued selective oxidation of Al from the thin foil results after a certain time in complete depletion of aluminum in the specimen. At this point, the alloy matrix can no longer maintain equilibrium with the Al2O3 scale, the oxygen potential beneath the alumina rises, and Cr2O3 formation at the alloy-scale interface commences. The resulting chromia layer grows beneath alumina at a remarkably slow rate. Its development serves to lengthen the alloy section lifetime by preventing immediate catastrophic oxidation after complete exhaustion of the alloy Al reservoir.
The outer Ni(Al,Cr)2O4 spinel formed during the transient oxidation stage becomes depleted in chromium due to evaporation of volatile chromium species. When the underlying alumina stops growing as a result of complete Al depletion, it becomes permeable to nickel, allowing further spinel growth and partial dissolution of Al2O3. More seriously, the newly formed chromia subscale dissolves the alumina layer from beneath. Eventually, the scale becomes one of mixed spinel Ni(Al,Cr)2O4 overlying (Cr,Al)2O3, and is no longer protective at these high temperatures.
Mr. Cosler and Ms. A. Kick are kindly acknowledged for carrying out high temperature exposures and TG-analyses, Mr. M. Ziegner for XRD measurements, Mr. J. Bartsch and Mr. V. Gutzeit for metallographic preparation. The authors are grateful to Mr. M. Borzikov for carrying out SNMS analyses.