Application of a dispersed-reservoir concept for the development of oxidation-resistant alloys

Corrosion Science, Vol. 35, Nos 5-8, pp. 841-854. 1993 Printed in Great Britain.

0010-938X/93 $6.00 + 0.00 © 1993 Pergamon Press Ltd

APPLICATION OF A DISPERSED-RESERVOIR CONCEPT FOR THE DEVELOPMENT OF OXIDATION-RESISTANT ALLOYS I. G. WRIGHT,* V. NAGARAJAN* and J. STRINGER'~ * Battelle Columbus Laboratories, 505 King Avenue, Columbus, OH 43201, U.S.A. + Electric Power Research Institute, Palo Alto, CA, U.S.A.

Abstract--Resistance to high-temperature oxidation is typically provided by either alloying with the desired scale-forming element or elements, or by applying a coating rich in those elements. For some alloy systems, the alloying approach requires very large additions of the oxide-forming elements and can lead to drastic reduction in the mechanical properties and melting temperature of the alloys. The dispersedreservoir concept is based on the incorporation of a reservoir of the oxide-forming element as a second phase, which is inert to the alloy substrate, but which reacts with the environment to form a protective outer scale. The selection of the reservoir phase is primarily governed by thermodynamic considerations; its effectiveness also depends on other factors such as its size and distribution in the matrix, which are controlled by processing parameters. The research represents an examination of the applicability of this concept to the development of inherent oxidation resistance in three different classes of alloys. In each case considered, the mode of growth of the scale of the parent alloy is different, and occurs by: (i) the predominant outward diffusion of cations; (ii) the predominant inward diffusion of anions; or (iii) the formation of a volatile oxide. Observations of the high-temperature oxidation behavior of these different systems has provided some insights into the ways in which the dispersed reservoir phase exerts its effect.

INTRODUCTION

TRE focus of the research reported here was on the design of high-temperature alloys which would form protective SiO 2 and A1203 scales on exposure to hightemperature, oxidizing environments. One way to promote the growth of such scales is to incorporate sufficient amounts of Si or A1 in the alloy substrate; typically, additions of approximately 35-45 wt% (all alloy compositions are given in wt% unless stated otherwise) of Si or A1 would be required to form the respective scales, in the absence of additional protective-scale-forming elements. At such levels of alloying additions, the formation of intermetallic phases is likely, which may have undesirable features such as low melting temperatures or poor mechanical properties. Attempts to alleviate these problems have been largely unsuccessful. ~ The research described was an attempt to address these problems from a fundamentally different point of view, through the use of a dispersed-reservoir phase. The dispersed-reservoir approach 2-4 is an attempt to effectively decrease the critical volume fraction of protective scale required to ensure the formation of a continuous external scale by physically reducing the area to be covered by the protective oxide. This approach is based on previous research, 5 in which it was demonstrated that both a binary Co-25Cr alloy and a ternary C o - 2 5 C r - C alloy, which contained sufficient C to precipitate virtually all of the Cr in the alloy as chromium carbides, oxidized at the same rate and formed protective Cr203 scales. In the ternary alloy, the chromium carbides acted as local reservoirs of Cr, and effectively concentrated the desired element (Cr) in the location where it was 841

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I.G. WRXGHT,V. NAGARAJANand J. STRINGER

r e q u i r e d for scale f o r m a t i o n , a n d r e d u c e d t h e d i s t a n c e t h a t the C r m u s t diffuse to t h e r e q u i r e d site. A g e n e r a l i z e d a p p l i c a t i o n o f this a p p r o a c h to t h e d e s i g n of o x i d a t i o n - r e s i s t a n t alloys r e q u i r e s t h e s e l e c t i o n o f r e s e r v o i r p h a s e s w h i c h a r e t h e r m o d y n a m i c a l l y stable in t h e a l l o y , b u t a r e u n s t a b l e in t h e p r e s e n c e o f o x y g e n , so t h a t t h e d e s i r e d e l e m e n t can t a k e p a r t in t h e s c a l e - f o r m i n g p r o c e s s . T h e o r i g i n a l c o n c e p t was t h a t the r e s e r v o i r p h a s e w o u l d d e c o m p o s e in t h e p r e s e n c e o f a s u i t a b l e o x y g e n g r a d i e n t , t h e r e b y r e l e a s i n g t h e d e s i r e d e l e m e n t to diffuse to o r n e a r the free alloy surface, a n d to f o r m t h e d e s i r e d scale. A s will b e s h o w n , this is o n e o f t h e a p p a r e n t m e c h a n i s m s that is o b s e r v e d in p r a c t i c e . EXPERIMENTAL METHOD To demonstrate the viability of the alloying approach described, a series of Ni-, Mo- and Nb-based alloys was prepared by powder metallurgical techniques. These alloys contained a dispersed-reservoir phase chosen from thermodynamic considerations, and intended to lead to the formation of an SiO2 scale on Ni and Mo, and an A120 3 scale on Nb. Ni-9SiC alloy specimens were oxidized in dry air in a thermobalance. Oxidation exposures were made for durations ranging from 0.5 to 250 h, at temperatures ranging from 800to 1400°C. Specimens of the Ni9SiC alloy were also Oxidized in a horizontal tube furnace maintained at 1050°C, using H2/H2 O mixtures which gave oxygen partial pressures lower than the dissociation oxygen partial pressure of NiO, which is 10 -9 atm at this temperature. Since the oxygen dissociation partial pressure of SiO2 is 3 × 10-27 atm at 1050°C, Si and SiC were expected to oxidize to form SiO2. The durations of these oxidation runs were varied between 2 and 44 h to allow the progress of the oxidation process to be monitored. Some of these specimens were pre-oxidized at an oxygen partial pressure of 5 x 10-18 atm for 24 h at 1050°C, and then subjected to further oxidation in air, in the thermobalance, at temperatures of 1000, 1050, 1100 and 1150°C for up to 100 h. The oxidation exposures of the refractory metal specimens were carried out in a horizontal tube furnace equipped with an alumina reaction tube, and the specimens were contained in alumina or zirconia boats. Weight change data were generated by periodicallywithdrawing the specimens from the furnace. EXPERIMENTAL RESULTS AND DISCUSSION N i - S i C alloy T h e w e i g h t c h a n g e d a t a o f s p e c i m e n s o f an N i - 9 S i C alloy s u b j e c t e d to i s o t h e r m a l o x i d a t i o n in air f r o m 800 to 1150°C a r e s h o w n in Fig. l ( a ) . T h e d a t a a r e p l o t t e d as w e i g h t g a i n v e r s u s the s q u a r e r o o t o f t i m e , a n d t h e s l o p e o f t h e s t r a i g h t line t h r o u g h t h e p o i n t s is t h e p a r a b o l i c r a t e c o n s t a n t , kp. C l e a r l y , t h e s p e c i m e n s initially e x h i b i t e d p a r a b o l i c b e h a v i o r . T h e r a t e i n c r e a s e d with i n c r e a s i n g t e m p e r a t u r e , with t h e e x c e p t i o n o f t h e d a t a for 1100°C. A n A r r h e n i u s p l o t o f log k p a s a f u n c t i o n o f t h e i n v e r s e of a b s o l u t e t e m p e r a t u r e s u g g e s t e d an a c t i v a t i o n e n e r g y o f a p p r o x i m a t e l y 73 kcal mo1-1. F o r h i g h - p u r i t y Ni, t h e a c t i v a t i o n e n e r g y for o x i d a t i o n to f o r m N i O r a n g e s f r o m 41 to 68 k c a l mo1-1, w h e r e a s the w i d e l y - u s e d v a l u e is 45 k c a l mo1-1. O t h e r d a t a r e p o r t e d in t h e l i t e r a t u r e i n d i c a t e an a c t i v a t i o n e n e r g y for o x y g e n diffusion in SiO 2 fibers o f a p p r o x i m a t e l y 71.2 kcal m o l - 1 , 6 w h e r e a s t h e v a l u e for o x y g e n diffusion in q u a r t z is a p p r o x i m a t e l y 55 kcal m o l - 1 . 7 G i v e n t h a t t h e p a r a b o l i c r a t e c o n s t a n t s w e r e c a l c u l a t e d f r o m t h e b e s t s t r a i g h t line fit to all t h e o x i d a t i o n d a t a at e a c h t e m p e r a t u r e , it is p o s s i b l e t h a t , in s o m e cases, t h e r a t e - c o n t r o l l i n g p r o c e s s was n e i t h e r s i m p l y diffusion o f Ni t h r o u g h N i O n o r o x y g e n diffusion t h r o u g h SiO2, b u t i n s t e a d was diffusion o f Ni t h r o u g h a b l o c k i n g l a y e r o f SiO2 p a r t i c l e s , o r e v e n a combination of these processes. T h e w e i g h t c h a n g e d a t a o f t h e s p e c i m e n s o x i d i z e d in the low o x y g e n p a r t i a l p r e s s u r e e n v i r o n m e n t s s u g g e s t e d t h a t e i t h e r v e r y small w e i g h t gains o r small w e i g h t

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losses occurred. The oxidation kinetics of specimens preoxidized in a low-oxygen partial pressure and then subjected to oxidation in air are shown in Fig. l(b). Comparison of these data with those for air oxidation without a preoxidation treatment [Fig. l(a)] clearly indicates that the protective oxidation behavior initiated during oxidation in the low-oxygen partial pressure (Hz/H20) persists, and that the oxidation rates of the preoxidized and non-preoxidized specimens were very similar. Also, whereas the oxidation rates were significantly lower than those for unalloyed Ni, they were faster than those associated with oxidation controlled by the growth of an SiO2 scale. This observation suggests that the rate-controlling layer on this alloy possibly contains some regions of Ni-rich oxide, which provide paths for continued diffusion of Ni. Figure 2 shows a secondary electron image (SEI) and X-ray maps for Ni, Si and O of a cross-section of a non-preoxidized specimen oxidized in air at 1000°C. It is apparent that the specimen formed a thick, outer scale of NiO, underneath which is an essentially continuous layer of Si-rich oxide, presumably SiO2. In fact, the morphology of this SiO2 subscale in some areas resembled a collection of oxidized SiC particles. The carbon map showed regions of carbon enrichment, which may have been due to graphite precipitation as a result of the decomposition of SiC particles upon release of Si to form the oxide layer. Very little spallation of the outer NiO layer was observed. Figure 3 shows a cross-section of the scale formed on a nonpreoxidized specimen oxidized at 1100°C. This particular specimen appeared to have formed a more uniformly thick and apparently more compact SiO 2 subscale than that formed at 1000°C, with few of the stringers (apparently of internally-oxidized SiC particles) that were found in the 1000°C specimen. The cross-sections of the specimens oxidized in the low-oxygen partial pressure conditions revealed the presence of a thin, adherent oxide film, which was shown by electron-probe microanalysis (EPMA) to be rich in Si and O, as is shown in Fig. 4, and were most likely SiO2. As expected, Ni was not detected in these layers. Beneath the SiO2 layer, there was a depleted zone at the inner edge of which were large, elongated, dark-appearing particles that were shown, by EPMA, to be essentially carbon (graphite), apparently precipitated on prior SiC particle boundaries. Overall, both the weight gain data and the microstructural data suggest that the oxidation behavior of the Ni-SiC alloy in air and in the Hx/H20 mixtures was controlled by diffusion through a basal layer of Si-rich oxide, which appears to consist of Si-rich (SiO2) areas with some interdispersed Ni (NiO). Recently, Stott et al. s have reported the formation of SiO2 scales on Ni-Si alloyes containing 1, 4 and 7Si oxidized at 1000°C in oxygen. The SiO 2 layers formed on those alloys contained imperfections, were not fully protective and tended to spall on cooling. The morphological observations suggest that the mechanism of formation of the SiO2 scale in air on the Ni-SiC alloy is different from that in Ni-Si alloys. These differences may account for the different behavior of the SiO2 scale on these alloys. The tentative mechanism suggested for the formation of SiO2 on the Ni-SiC alloy oxidized in a low-oxygen partial pressure environment is based on the decomposition of the SiC particles to release Si, as follows: on initial exposure, the SiC particles in the alloy matrix decompose to Si (dissolved in Ni matrix) and C (precipitated as graphite). The Si dissolved in the matrix diffuses toward the surface of the alloy. When the Si and O activities exceed the solubility product of SiO2, Si oxidizes to form SiO2 particles. This reaction may occur at or beneath the alloy surface,

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depending on diffusivities and solubilities. Particles of SiO2 link up and establish a continuous oxide layer at or slightly beneath the alloy surface. After a continuous and protective SiO2 layer is established, the oxidation behavior is probably controlled by the diffusion of oxygen through SiO2, so that the specimen undergoes very little further weight change. The microstructure presented in Fig. 4 is consistent with this proposed mechanism. The depleted zone beneath the SiO2 scale, which is devoid of SiC particles, confirms that the appearance of the SiO2 scale corresponds with the decomposition of the SiC particles. However, carbon was deposited as graphite beneath the depleted zone, for reasons that are not known at present. The mechanism of protective scale formation on the non-preoxidized alloy in air is thought to rely on morphological factors, as suggested schematically in Fig. 5. When a clean surface of the alloy is exposed to an oxygen-containing gas (at time t = 0), oxygen dissolves in the alloy and diffuses inward. Simultaneously, Ni atoms and SiC particles on the surface oxidize to form NiO and SiO2, respectively. The NiO nuclei grow more rapidly than the SiO2 nuclei and, since NiO grows by cation diffusion, its growth causes the injection of vacancies into the alloy surface. Although SiC is considered to be a stable phase, decomposition is thermodynamically feasible. Since a finite time is necessary for SiC to decompose and for the released Si to diffuse to the oxidation front, in the initial stages of oxidation it is considered likely that the SiC particles near the free alloy surface are oxidized in situ, resulting in SiC particles surrounded by an SiO2 shell. At t = 2, the outward diffusion of Ni 2+ ions has resulted in the formation of a continuous outer layer of NiO, which reduces the partial pressure of oxygen at the oxide-metal interface to that corresponding to the dissociation partial pressure of NiO. At l l00°C, the resulting oxygen partial pressure would be approximately 10 s atm, a level sufficiently high that oxidation of SiC particles or, depending on the rate of ingress of the oxidation front, decomposition of these particles followed by oxidation of the released Si, can continue beneath the outer NiO layer. At t = 3, consumption of Ni from the surface, probably assisted by vacancy condensation/void formation around the SiC particles, leads to the accumulation of oxidized SiC particles at the metal-NiO interface. The SiO2-coated SiC particles coalesce, probably sinter together, and form a semi-continuous layer. At this point, oxygen ingress into the alloy is greatly reduced, which may allow time for the SiC particles beneath the semi-continuous layer to decompose, and the released Si to diffuse to gaps in the SiO2 layer where it is oxidized and contributes to the completion of the surface layer. As a result, at t = 4, the alloy has formed a continuous and protective basal scale of essentially SiO2, which controls the subsequent oxidation behavior. Mo-Si3N4-Si alloy The weight change data of specimens of an Mo-15Si3N4~Si alloy oxidized in air at 1000, 1200 and 1400°C are shown in Fig. 6. All the specimens initially lost weight due to the formation of volatile MOO3, the rate of loss increasing rapidly with increasing temperature. However, after a few hours, the rate of weight loss abruptly stopped, indicating that formation of the volatile MoO 3 had been suppressed due to the formation of a protective, solid oxide scale. Figure 7 shows a cross-section of a specimen oxidized for 111 h at 1000°C. The SEI and X-ray maps for Si and O indicate that the solid oxide layer consists of Si and O, and is presumably SiO2.

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A suggested mechanism for the formation of a protective SiO 2 scale on this Mo alloy is shown schematically in Fig. 8. The sequence of events that occurs at time t = 1 is as follows. Mo oxidizes and forms the volatile MOO3; simultaneously, oxygen dissolved in the Mo and diffuses inward, and Si and Si3N4 particles at the surface oxidize to form SiO2. Also, Si from the alloy diffuses toward the surface and, where the Si and oxygen activities are sufficiently high, oxidizes to form SiO2. Decomposition of Si3N4 is suppressed by the presence of Si in the alloy. At time t = 2, volatilization of MoO3 has caused recession of the alloy surface which, in turn, causes the oxidized Si3N4 particles to accumulate and agglomerate and possibly sinter together. This increasingly dense, blocking layer reduces oxygen ingress to the alloy

The dispersed-reservoir concept

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so that, at t = 3, Si in the alloy can diffuse to this layer, where it oxidizes and seals the gaps that exist between the agglomerated Si3N 4 particles. At t = 4, an essentially continuous SiO: based layer has been formed, and the specimen is protected from further oxidation and suffers little weight loss. NbAI3-AIN alloy Protective oxidation behavior is not exhibited by Nb-A1 alloys when oxidized in air at Ai levels below about 46 wt%, which corresponds to the intermetallic NbA13 (46.6A1). This alloy initially forms an alumina scale, but such scales are rapidly disrupted and oxidation at a linear rate ensues as a result of the formation in the A1depleted region beneath the scale of the next-lowest intermetallic, NbsA13, which has poor oxidation resistance. 9 Al-rich NbA13 was demonstrated to form consistently protective scales. An NbA13-20AIN alloy was found to consistently form a protective scale with a continuous basal layer of alumina, when oxidized for times in excess of 100 h at 10001400°C in air. The values of the parabolic rate constants are shown in the Arrhenius plot in Fig. 9. At 1000°C the rate constant was very similar to that found by Singheiser et al.l° for cast NbAl 3, and at 1000 and 1200°C values align with those for the formation of A120 3 on Ni-AI. 11 Figure 10 shows an optical micrograph of a specimen of NbAI3-20A1N oxidized at 1000°C for 31 h. It is quite clear that the specimen formed a protective Al-rich scale, with patchy areas of an Nb-A1 oxide at the scale gas interface; X-ray diffraction analysis confirmed that the scale was a-AlzO 3. Any explanation of the effects of the dispersed-reservoir phase on the development and maintenance of a protective alumina scale on NbAI 3 must account for the facts that, since Nb oxides grown by anion diffusion, there is no injection of vacancies, nor accumulation of oxidized reservoir particles beneath the outer scale, as in the case of the Ni-SiC alloy, and that there is no volatilization to assist the accumulation of dispersoid particles at the alloy surface, as in the case of the Mobased alloy. For NbA13, the formation of a protective scale is thought to be strictly controlled by the initial composition of the alloy, and the initial distribution of A1N in 50 1000°C

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0 0 0 0 0 Fro. 8. Suggested mechanism for the formation of a protective SiO2 scale on Mo-15Si3N4--6Si. the alloy. One possible function of the A1N reservoir phase in the N b A I 3 - A I N alloy may have been to compensate for the A1 consumed to form the protective scale, and to prevent the degradation of N b A I 3 to NbsAI3. CONCLUSIONS

The experimental results reported clearly demonstrate that the ability to form a protective oxide scale can be incorporated into a range of alloy types by using the dispersed-reservoir approach. The use of appropriate reservoir phases, for instance, should allow the formation of any of the three practically useful oxides: CreO3, AlzO 3 or SlOe or, in principle, any other desired oxide. However, this ability comes at the price of incorporating into the alloy a dispersion of second-phase particles,

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which has implications for the mechanical properties of the alloy. A t the level and size used in the N i - S i C alloy, the dispersed particles resulted in a reduction of the r o o m t e m p e r a t u r e ductility of the alloy to 3%. While the alloys investigated in this p r o g r a m were not optimized in terms of the m i n i m u m v o l u m e fraction and particle size of the dispersed second phase, its effect on mechanical properties is obviously an i m p o r t a n t consideration. A l t h o u g h the dispersed-reservoir alloys described were m a d e by high-energy milling and p o w e r consolidation techniques for expediency, there appears to be no reason why processes that are perhaps better suited to higher p r o d u c t i o n volumes, such as p o w d e r p r o d u c t i o n by rapid solidification or co-precipitation and spray drying, could not be used. Obviously, further research is required to translate these concepts to the design of practical, h i g h - t e m p e r a t u r e c o m p o n e n t s , but the underlying principles a p p e a r to offer promise. A s an example, the introduction of a dispersed-reservoir phase to an alloy with otherwise p o o r resistance to oxidation, and the application of that dispersed-reservoir-containing alloy as a coating to protect the p a r e n t alloy, offers the advantages of eliminating c o a t i n g - s u b s t r a t e interdiffusion p r o b l e m s and minimizing thermal expansion mismatch. Acknowledgements--The results described for the Ni-SiC alloy were obtained from programs supported by the Electric Power Research Institute under RP 2426-03 and by Battelle. The NbAI3 alloy was produced in a program supported by the Naval Air Development Center, Warminster, PA, on Contract No. N62269-89-C-0252.The data for the Mo-based alloy were from studies supported by Battelle. We are very grateful for this support, and would like to thank Dr J. Stringer of EPRI and Mr M. K. Thomas of NADC for the encouragement and support. The experimental work was accomplished by Mr R. K. Beale, Mr P. N. Mincer and Mr R. D. Smith.

REFERENCES 1. G. H. MEIER,in Oxidation ofHighoTemperature Intermetallics (eds T. GROBSTEINand J. DOYCHACK). TMS (1988).

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