High temperature oxidation of Co-Mn alloys

Corrosion Science, Vol. 19. pp. 675 to 691 Pergamon Press Ltd. 1979. Printed in Great Britain

HIGH TEMPERATURE OXIDATION OF Co-Mn ALLOYS* F. GESMUNDO, P. NANNI a n d D. P. WHITTLE'~ Centro Studi di Chimica e Chimica Fisica Applicata alle Caratteristiche di Impiego dei Material/, Consiglio Nazionale delle Ricerche, Genova, Piazzale Kennedy, Italy Abstract--The oxidation of cobalt-manganese alloys in the range 0-45 wt yoMn corresponding to the stability of the et-phase of cobalt has been studied in the range 750-950°C as an example of binary alloys producing solid solution scales. The alloys oxidize according to a parabolic rate law with a rate constant intermediate between those of the pure metals. The scales were formed mainly by a CoOMnO solid solution with a spinel phase present either as dispersed particles or a continuous layer in the cubic oxide, particularly for the Mn-rich alloys: some degree of internal oxidation has also been observed. The scale properties show a time dependence at constant temperature which is unusual in these systems and is probably related to the presence of the spinel phase. The internal oxidation and the structure and composition of the external scale are discussed with reference to the phase diagram for this system and to the recent theories for the growth of solid solution scales. INTRODUCTION THE OXIDATION o f b i n a r y alloys f o r m i n g solid solution scales has been studied in detail in recent years, b o t h f r o m a theoretical 1-6 a n d e x p e r i m e n t a F -17 p o i n t o f view. I n particular, the dependence o f the metal c o n c e n t r a t i o n profiles in scale a n d alloy phases, the oxygen activity profile in the scale, a n d the p a r a b o l i c g r o w t h rate o f the scale u p o n the t r a n s p o r t a n d t h e r m o d y n a m i c p r o p e r t i e s o f b o t h the alloy and scale have been established. C o - M n alloys, in which the c o m p o n e n t oxides M n O a n d C o O f o r m a c o n t i n u o u s solid solution, have n o t however been studied. This system has a further c o m p l i c a t i n g factor since in a d d i t i o n to the solid cubic oxide solution, a spinel phase, showing a c o n t i n u o u s solid solution in the range CoaO4-Mn304, can also be formed, a n d it is o f interest to establish h o w the presence o f this a d d i t i o n a l phase modifies the overall scaling behaviour, and in p a r t i c u l a r the c a t i o n distributions in the cubic oxide solution. T h e present w o r k investigates the o x i d a t i o n b e h a v i o u r o f C o - M n alloys in the range 0-45 w t ~ o M n at t e m p e r a t u r e s in the range 750-950°C a n d at 10sPa 02. T h e limiting c o m p o s i t i o n on the M n - r i c h side is dictated b y the limited solubility o f M n in the or-phase o f cobalt. ~s EXPERIMENTAL METHOD The alloys were prepared by melting 99.97 ~oMn and 99.998 ~oCo in a H.F. furnace and casting into a water-cooled mould. The analysed compositions of the alloys were (wt. ~oMn): 6.5, 23.72 and 46.19Y0. The ingots were annealed for 12 h at 800°C in vacuo ( ~ 1.3 × 10-spa), and then cut into discs approximately 0.8 mm thick and polished on SiC papers. Just before the runs, they were given a final polishing to 6/0 emery paper. The specimens were then washed in distilled water, acetone and all-ethyl ether and finally dried in warm air. *Manuscript received 30 October 1978. "l'Department of Metallurgy and Materials Science, University of Liverpool, Liverpool L69 3BX, England. 675

676

F. GESMUNDO,P. NANNIand D. P. WmTTLE

Oxidation was carried out at I x 10sPa oxygen under static conditions in a Mettler thermoartalyser. The balance chamber was flushed through with high purity Argon for 2 h, and evacuated to 8 x 10-~Pa. The samples were heated up to 700°C at 10°C rain-x and then to the selected temperature at 100°C rain-t, to avoid Mn loss at high temperatures. At the selected temperature, oxidation was commenced by admitting high purity oxygen. Runs were performed at 750, 850 and 950°C for periods up to 15 h. At the end of the experiments, the specimens were cooled down in vacuo, in situ; no further changes in weight were detected. Transverse sections of the oxidized samples were examined by optical and electron microscopy; concentration profiles were measured by an EDAX attachment to a JEOL SEM. X-ray diffraction was used for phases identification. EXPERIMENTAL RESULTS

Oxidation kinetics O x i d a t i o n o f the three alloys follows a p a r a b o l i c law. The scales f o r m e d generally showed g o o d adherence at temperature, a l t h o u g h scale spallation on cooling was often observed. This was p a r t i c u l a r l y the case for samples oxidized at 750°C when the scales a l m o s t completely spall off. The p a r a b o l i c o x i d a t i o n constants (g~cm-4s -1) as functions o f t e m p e r a t u r e a n d alloy c o m p o s i t i o n are r e p o r t e d in T a b l e 1: these are r e p r o d u c i b l e within 4- 10--15%. Values for p u r e C o a n d p u r e M n have also been measured a n d are included for c o m p a r i s o n . T h e p a r a b o l i c rate constants increase only slightly with the M n content o f the alloy. This is expected since the o x i d a t i o n rate o f pure M n is only slightly greater t h a n t h a t o f pure C o u n d e r c o r r e s p o n d i n g conditions. T~LE 1.

PARABOLIC

OXIDATION CONSTANTS AT 10 5 P a OXYGEN

(g~ cm -~ s-X) 750°C Co Co--6.5~,Mn Co-23.72~0Mn Co-46.19~,Mn Mn

2.4 2.3 4.4 3.2 4.3

x x x x x

10-lo 10-x° 10-1° 10-1° 10-1°

850°C 1.4 3.7 4.5 1.6 4.3

x x x x x

10-° 10.9 I0 -a 10-° 10-°

950°C 1.1 3.7 1.4 3.7 2.0

x x x x x

10.8 10-a 10-s 10-a 10-a

The p a r a b o l i c o x i d a t i o n constants have also been p l o t t e d a c c o r d i n g to an A r r h e n i u s e q u a t i o n to evaluate the activation energies. F o r the 6.5 a n d 23.72 % M n alloys a g o o d fit was observed over the whole range o f temperature, whereas for the 46.19 % M n alloy, the activation energy could be estimated only for s h o r t t e r m (2 h) exposures; rate constants f r o m long t e r m exposures showed a b r e a k in the A r r h e n i u s plot. The 6 . 5 % M n alloy gives an activation energy o f 262 k J m o l - a ; with the 23.72 % M n alloy the value was s o m e w h a t d e p e n d e n t on the exposure time increasing f r o m 184 to 206 a n d 215 kJ mo1-1 after 2, 5 a n d 15 h exposure, respectively. A value o f 240 k J m o l - t was f o u n d for the 46.19% alloy. Identical values for p u r e C o a n d M n were f o u n d in the present investigation: 201 kJ mo1-1.

X-rajp diffractometry, metallography and micro-analyses Diffraction analyses were usually carried o u t on the scales as f o r m e d on the oxidized alloys. I n a few cases p o w d e r e d scales were examined. I n b o t h m e t h o d s , only the diffraction p a t t e r n s for C o O - M n O solid solution were detected. It was n o t possible to analyse samples which showed extensive scale spallation. T h e results for the individual alloys are presented separately.

t

10 IJm o

2o m FIG. 1.

Cross-section of the scale formed on C o - 6 . 5 ° o M n oxidized for 15 h at 850 C.

FIG. 2.

Cross-section of the scale formed on Co-6.5 ~'.;,Mn oxidized for 15 h at 950 C.

10 pm (a)

Fla. 4.

Cross-section of the scale formed on Co-23.72 7oMn oxidized for 5 11at 950 'C, (a) optical image; (b) Co Kct image; (c) Mn Ket image.

7

10pm 5

FIG. 5.

Cross-sect on of the scale formed oft Co-23.72 ~ M n oxidized for 15 h at 750°C.

FiG. 6.

Cross-section of the scale formed on Co~,6.19 ~ M n oxidized for 5 h at 850'-C.

~'.lk

8(0), +

:

+

+ ' I!' .

•, .

.~. +~

.

..

FIG. 8. Cross-section of the scale formed on Co-46.19 ~oMn oxidized for 5 h at 950°C, (a) optical image. FIG. 9.

Cross-section of the scale formed on Co-46.19 ~oMn oxidized for 15 h at 950°C, (a) optical image.

High temperature oxidation of Co-Mn alloys

681

Co-6.5 ~oMn Scales formed at 750°C spalled off completely on cooling, so that no microscopic investigations could be carried out. Optical observations of cross-sections of samples oxidized at 850°C showed a rather highly porous zone at the alloy-scale interface, whilst the external layer was very compact. Micro-analysis of the scales formed on samples oxidized for either 2 or 5 h indicated little variation in composition between the two layers; the Co/Mn ratio was practically constant across the section, corresponding to approximately 70 ~oCo and 5 ~oMn. X-ray diffraction revealed a cubic NaC1 structure and identified the scale as a Co-rich, solid solution of CoO-MnO. Figure 1 shows a cross-section of the scale formed after 15 h exposure, and an additional, light grey, fine-grained phase was found embedded in the outer compact zone of the scale. Quantitative micro-analysis showed that the composition of the cubic oxide was again uniform across the section and similar to that of the scales formed in shorter exposure times. The precipitated phase, however, contained only 44 ~ C o , but the Mn content was 23.6 ~ . This composition is very close to that of stoichiometric MnCo204. This spinel phase was not observed metallographically in the scales formed after 2 and 5 h exposure, but there were weak, additional lines in the X-ray patterns which undoubtedly corresponded to this phase. At 950°C the light, dispersed phase was observed in the scale even after 5 h exposure. Furthermore, after 15 h, a practically continuous layer was formed, approximately half-way through the scale as shown in Fig. 2. The Mn content in the cubic oxide at the outer part of the scale was only 1-2 ~o, whilst in the inner part, on the alloy-oxide interface of the spinel particles, it was 4-5 ~o: similar to that measured at lower temperatures. Figure 3 presents the concentration profiles through the scale. The analysis of the spinel is similar to that at 850°C: however, in analogy

P o r o u s Void :i = zone ~

Oxide 100

i bd 3" hi .--I LI.I **

Alloy

i

Co

so

100

200

300

400

500

DISTANCE, pm

FIG. 3. Concentration profiles for Co and Mn across the scale formed on Co-6.5 %Mn oxidized for 15 h at 950°C.

682

F. GESMUNDO)P. NANNIand D. P. WHrI'rLE

with Ni-Mn, 7 Fe-Ni 13 and Fe-Co 18 systems, the spinel composition could range from Co304 to Mn304 .x9 There is only a limited degree of internal oxidation with this alloy. Co-23.72 ~oMn Scale spallation after 2 or 5 h exposure at both 750 and 850°C precluded metallographic examination. The scales formed after 15 h exposure at either temperature were essentially identical to those formed on the lower Mn alloy, with a discontinuous precipitate of the spinel phase embedded in the cubic oxide solid solution. The cubic oxide contained significantly more Mn in comparison to that formed on the Co-6.5~oMn alloy: 20-23~oMn in comparison to 5-6~o respectively. In the cubic oxide at the outer part of the scale, beyond the spinel precipitates, the Mn content was reduced to less than 5 ~o. The spinel itself contained 27-36 ~oMn and 53-42 ~Co. As with the lower Mn alloy, the inner region of the scale tended to be porous. Although in general, the spinel particles were confined to a relatively narrow band in the scale (Figs. 1 and 2) the concentration profiles through the scale occasionally showed some compositional fluctuations in what appeared to be the cubic oxide phase, and it is possible that very fine spinel particles were sometimes present almost throughout the scale cross-section. At 950°C and after 2 h exposure the concentration profiles in the scale were quite uniform, with a slight increase in Co content from 55 ~o near the alloy-scale interface to 62.5 ~o at the outside of the scale: the spinel phase was absent. However, after 5 and 15 h exposure, a continuous spinel layer was very pronounced and appeared about three-quarters of the way across the scale. Figure 4 shows a typical cross-section together with the Co and Mn X-ray images. The change in composition of the cubic oxide layer from one side of the spinel region to the other is also obvious: analysis gave a Mn content of 20 ~o in the inner cubic oxide and about 2 ~o in the outer part. The spinel layer contains about 46 ~oMn and 28 ~oCo. This alloy shows a more pronounced zone of internal oxidation than the lower Mn alloy, particularly at the lower temperatures. In many cases, sample preparation damage precludes detailed analysis. However, Fig. 5 shows an area where the internal oxide has been retained and microanalysis indicates a composition very rich in Mn, with only about 2-3 ~oCo. Co-46.19 ~oMn X-ray diffraction of the scales formed in 2 h showed only the solid solution of CoO and MnO. After 15 h oxidation, islands of the light coloured Mn-rich phase (Mn ~ 70 ~o, Co ~ 5 ~ ) were observed, buried between two darker layers of different concentration, the inner one (Mn ~ 65 ~o, Co ~ 20~o) being more Mn-rich than the outer (Mn ~ 35~o, Co ~ 50~o). This is similar to the lower Mn content alloys, except that the Mn contents of all the scale layers are generally higher. At 850°C the scale composition changed with oxidation time. For short periods (2 h) only one layer of virtually pure Mn oxide with small amounts of Co ( ~ 0.5 ~ ) was observed. For longer times (5 and 15 11) an inner grey layer of cubic oxide of constant composition ( ~ 50 ~oMn, ~ 30 ~oCo for the 5 h runs), was surmounted by a continuous lighter grey layer of variable composition extending up to the scale surface

High temperature oxidation of Co-Mn alloys

683

(Fig. 6). In this outer layer, the Mn content decreases from about 72 % at the inner interface to about 10% at the scale surface as shown in the concentration profiles presented in Fig. 7. This layer is identified as the spinel, with a strong Mn depletion in its outer portion.

Alloy

Oxygen

Mn " ' 5C

I 20

I 40

DISTANCE, pm

FIo. 7. Concentration profilesfor Co and Mn across the scale formed on Co-46.19 %Mn oxidized for 5 h at 850°C (scale only). Arrows indicate the positions of scale--oxygenand scale-alloy phase boundaries.

At 950°C, the samples oxidized for 2 h showed compact, single layer scales. Figure 8(a) shows a typical cross-section after 5 h oxidation. Again there is apparently only a single phase although the inner region is slightly darker in colour. The concentration profiles through this scale, shown in Fig. 8(b), are rather anamolous. Mn is enriched towards the outer part of the scale but then the extreme outer part of the scale shows a substantial enrichment in Co. If a spinel particle or layer were present between these two regions, then the observations would be consistent with those presented earlier; it was, however, not observed. Figure 9(a) shows a cross-section of the scale formed after 15 h oxidation and 9(b) the corresponding scale concentration profiles. Discrete islands of the Mn-rich spinel are present at the surface of the scale. For all the alloys of this group a large zone of internal oxidation was observed, with a morphology similar to that of the 23.72 %Mn alloys. At 950°C the presence of crevices was also observed, protruding deeply into the alloy (Fig. 8). A severe Mn depletion in the alloy towards the alloy-scale interface was observed, whilst the Mn]Co ratio in the surface scale at the interface was approximately the same as in the bulk alloy (Fig. 9(b)).

684

F. GESMUNDO,P. NANNIand D. P. Wnn-rLE

Oxygen

Alloy

LO LiJ ...J tad

50

I--1-tD LIJ

I

I

100

200

FIG. 8 (b). Co and Mn profiles across the scale only. Arrows indicate the positions of scale-oxygen and scale-alloy phase boundaries.

Oxide 100

Void

Alloy

J

I-Z ILl LU laJ Mn (_~ LU

50

3:

DISTANCE,

600 pm

FIG. 9 (b). Co and Mn profiles across the scale.

DISCUSSION A peculiar feature of the results obtained in the present investigation is the time dependence of both the types of phases formed in the scale, and the cation concentration profiles. Generally, for alloys in which the component oxides are completely miscible, steady-state, diffusion-controlled scaling behaviour is established relatively quickly, and for a given set of exposure conditions, concentration profiles in the scale

High temperature oxidation of Co-Mn alloys

685

can be expressed as a single-valued function of fractional position in the scale. 1 Formation of the spinel phase may be the complicating factor, although the time scale of the exposures is, in most instances, rather long to suggest that nucleation effects of this phase are important. Certainly, the emerging behavioural scaling pattern of these alloys is rather complex, and Table 2 is included to summarize the various types of scales which have been observed. In addition, there is usually a zone of internal oxidation, which is more extended with the Mn-rich alloys. TABLE 2.

CLASSIFICATION OF THE OBSERVED SCALE MORPHOLOGIES ON ALLOYS OXIDIZED AT 750, 850 AND950°C

Alloy Co-6.5 ~Mn Co-23.72 ~oMn Co-46.19 ~Mn

Oxidation time (h)

750°C

2 5 15 2 5 15 2 5 15

~ ----B A B B

Scale structure* 850°C A A B --B A D D

Co-Mn

950°C A B C A C C A A E

*A: Single-phase scale of cubic oxide solid solution (CoMn)O. B: Two-phase scale comprising an inner layer of solid solution of cubic oxides (CoMn)O, an intermediate layer of spinel (CoMn)sO4, precipitates dispersed in the cubic solid solution, and an outer cubic solid solution layer enriched in cobalt. C: A s B but the intermediate layer contains a continuous spinel phase. D: A layered scale comprising an inner layer of the cubic solid solution and an outer, thick layer of spinel which extends to the gas-scale interface. E : A cubic solid solution with discrete spinel islands only at the outer scale surface. The scale structures and compositions can best be analysed by reference to the diagram presented in Fig. l0 which shows oxygen activity as a function of composition in the Co--Mn-O system at 950°C: similar shaped phase fields exist at 750 and 850°C. This diagram has been calculated by assuming thermodynamically ideal behaviour of the C o - M n and C o O - M n O solid solutions which represents somewhat of an approximation for the alloy 2° but is reasonably correct for the oxide system, sl The spinel solid solution shows consistent deviations from ideality is and these have been taken into account in constructing the phase diagram. The expected phases on this basis are: (1) alloy, (2) the cubic solid solution C o O MnO, and (3) the spinel solid solution CosO4-MnsO4. Both types of oxide solid solutions have a restricted composition field for a particular range of oxygen activities. Nevertheless, since the oxygen activity in the gas phase is higher than that required for the cubic oxide--spinel equilibrium, the outer part of the scale should be composed of spinel. In the inner regions of the scale there is the possibility of a mixed cubicspinel oxide region, and then only the cubic oxide is expected adjacent to the alloyscale interface. The presence of a measurable degree of internal oxidation is also anticipated.

686

F. G~SMU~rDO,P. N A z i and D. P. WmTr~

Spinel

c~

\ -8

'.(Co,Mn)O

...............

=_

. f'[}

-1E (Co,Mn)0 +Alloy

-a4 ~

-

IB

.....................

c~

i (Alloy)

Q0

05 NHn NHn÷Nc=

1.0

FIG. 10. Oxygen activity/composition diagram for the C o - M n - O system at 950°C: schematic diffusion paths are also included and are described i n the text.

The main difference between the predicted and observed scale structures is the presence in some cases of an outer cubic oxide layer surmounting a spinel-containing region, either as a continuous layer or as discrete particles, which suggests at first sight a reversal of the oxygen activity gradient through the scale. This peculiar scale structure, already found for other systems, 1°,~ehas not been well understood. Moreover, the absence of an outer spinel layer as well as the presence of an internal oxidation region have to be considered. These different aspects will be dealt with separately. Internal oxidation The formation of subscalc in the oxidation of alloys forming solid solution scales has already been observed in a number of systems ~-10,12,is, le but has not been discussed in any great depth. Existing theories for the oxidation of these alloys either neglect the inward diffusion of oxygen from the alloy-scale interface1 or assume that the oxygen concentration inside the alloy is insufficient to form internal oxide. ~ The production of internal oxide particles of variable composition is essentially different from the formation of particles of only one of the alloy components, which has been extensively analysedfl~, ~3 since in the former case, the compositions of both alloy and oxide phases are functions of the oxygen activity in the alloy. Nevertheless, if the stabilities of the two oxide components of the solid solution are widely different the concentration of the less noble element in the alloy in equilibrium with almost the pure oxide of that component is very small. The formation of internal oxide in these systems cannot be considered in isolation from the growth of the surface scale and this is best carried out in relation to the oxygen activity/composition diagram (Fig. 10) for the C o - M n - O system at 950°C: behaviour at other temperatures is qualitatively similar.

High temperature oxidation of Co-Mn alloys

687

As indicated in Fig. 10, pure MnO is more stable than pure CoO, and thus, according to thermodynamic considerations, there is always a tendency for manganese to be selectively oxidized from these alloys. However, according to a recent detailed analysis, e the diffusivities or transport rates of the various species are often more" important than the thermodynamic properties in determining the actual scaling behaviour. Thus, the rate of diffusion in the alloy relative to the rate of diffusion in the scale is critical, and this is best expressed as the ratio of alloy interdiffusion coefficient, Da,, to the self diffusion coefficient of one of the cations in its pure oxide at unit oxygen activity, DOM,: that of Mn 2+ in pure MnO will be used here, although as will be seen later, the two cations, Co 2+ and Mn 2+ have similar diffusivities in the solid solution oxide (CoMn)O. One limiting case occurs when diffusion in the alloy is very fast, Dall/D°Mn --> oo, and this corresponds to negligible depletion of Mn in the alloy, although the surface scale would be considerably enriched in manganese. The scale composition in equilibrium with any particular alloy would be indicated by a horizontal tie line in Fig. 10 such as BC (constant oxygen activity) running from the point on the (CoMn)/ (CoMn)O + (CoMn) phase boundary at the particular alloy composition. Clearly for most alloy compositions the oxide would be virtually pure MnO. The other limiting case corresponds to Da,/D°Mn ---> 0 which implies that there will be no overall enrichment of the less noble element in the scale. The present system approximates more closely to these conditions. Recent measurements 24 of the alloy interdiffusion coefficient in the range 5-40 at ~oMn at 950°C gave a value of 7.39 × 10-xz cm 2 s -x, and using a value of D°M, obtained by extrapolation of the data of Price et aL, 25 DOMn = 5.25 × I0 -s cm z s -1 at 950°C, the value DalI/DOMn " 10-4 . Accordingly, under these conditions, when there is no overall enrichment of the less noble element in the scale, the composition of the scale immediately adjacent to the alloy depends on the kinetic parameters, 6 and in particular, the ratio of the two cation diffusivities in the scale, designated p, and corresponding to the ratio Dco/DM ., that is the diffusivity of noble cation to that of the less noble cation. Dco/DMn cannot be obtained directly, since measurements of Dco and DM, in the solid solution (CoMn)O have not been made so far. It is, however, possible to obtain an approximate value for p by an indirect procedure. The cation self diffusion coefficients depend on • the concentration of cation vacancies in the form 4 N~ Dco = D°co ~ and N~coo

DMn =

No

D°Mn - -

NaMno

,

(1)

where D0M is the self-diffusion coefficient of M in the pure MO oxide at unit oxygen activity, N~MO is the mole fraction of cation vacancies of charge ~ relative to the lattice in pure MO at unit oxygen activity and N ° the mole fraction of vacancies in the solid solution oxide and which is dependent on both oxygen activity and oxide composition. Thus, dividing equations (1) leads to Dc° D'c°° N"ra"° -- ~ D°c° P -- DM~ - - D°MnO N~co-~o D°M.

(2)

688

F. GESMUNDO,P. NANNIand D. P. WHITTLE

where 13 = NaM,o/N~coO.4The value of 13is obtained from published data: the deviation in stoichiometry in CoO 2e gives N"coo -- 6.92 × 10-3 at 1000°C as an average value from the results of several authors 2e-a° and N~MnO : 0.1376 at the same temperature, 31 giving 13 -~ 20. The self-diffusion coefficient of cobalt in CoO has been measured as a function ofpc~ and temperature and at 1000°C and 10sPa 02 is given by D°co = 2 . 6 8 × 10-gem 2s -1 ~2whilst D°Mn = 4 . 6 × 10- s c m 2s -a is obtained for the same conditions by extrapolation of reported data on manganese diffusion in MnO. 25Thus, p = 1.15. Using the 23 ~oMn alloy as an example, the mole fraction of MnO, NM,o, in the solid solution corresponding to a value of p of 1.15 would be 0.2564 which is close to the measured value. Thus, according to Fig. 10 the alloy composition (point E) in equilibrium with this oxide (point D) is virtually pure cobalt, and corresponds to a substantial depletion of manganese in the alloy. However, this would only be very near to the alloy-scale interface since, as indicated above, the ratio of alloy interdiffusion rate to scale growth rate (as measured by D°M.) is very small. In essence, then, the alloy composition would change from its value in the bulk alloy, point A in Fig. 10, to the value at the alloy-scale interface, point E, essentially along the line ABE. However, it must be remembered that in plotting, albeit schematically, a diffusion path on the ternary isotherm in this manner, all indication of distances is iost, and the physical separation of points B and E can be very small. Nevertheless, it seems more likely, and is indeed observed, that internal formation of oxide takes place within the alloy. Again any degree of overall manganese enrichment in the internal oxidation zone is unlikely, since in addition to D°a~l/D°Mnbeing small, so too is the ratio N°MnDalj/DoNSo where N°Mn is the bulk mole fraction of manganese in the alloy, NSo the solubility of oxygen in the alloy at the alloy-scale interface and Do the diffusivity of oxygen in the alloy. This latter ratio essentially represents the ratio of fluxes of manganese towards the alloy-scale interface to oxygen away from the interface. If it is small, then there can be no enrichment of manganese in the zone of internal oxidation. Thus, the diffusion path through the two-phase internal oxidation zone follows essentially the line BD, indicating that the Co/Mn ratio is identical to that in the bulk alloy. The composition of the internal oxide and that of the alloy in contact with it will vary along the lines CD and BE respectively, with corresponding compositions being connected by horizontal tie-lines. The amount of oxide will also vary but this cannot be established directly from oxygen activity-composition plots such as Fig. 10, since the mass of oxygen present is not indicated. However, it is clear in a qualitative sense that the position of the diffusion path in the two-phase region, such as BD, relative to the phase boundary lines gives some indication of the amounts of oxide and alloy phases. Thus, near to the alloy-surface scale interface the mole fraction of precipitated oxide will be high, and this will decrease in a direction normal to the interface. More detailed quantitative analysis requires precise solution of the coupled transport equations in the two-phase region, and this is presently being attempted. As indicated in the results section, internal oxidation proceeds to different extents depending on the bulk alloy composition: internal oxidation is almost absent in the more dilute alloys, but rather intensive in the more concentrated alloys. Again, this is consistent with the predictions from Fig. 10, in that for low manganese content alloys,

High temperature oxidationof Co-Mn alloys

689

the relative amount of oxide will be small in the greater part of the internal oxidation zone. External scale The inner region of the external scale is composed of a cubic oxide solid solution, and at constant temperature and pressure this single-phase region has two degrees of freedom, one of which is fixed at any coordinate point by the oxygen pressure. The extra degree of freedom thus cannot be fixed by a pure thermodynamic parameter and hence a kinetic parameter comes into play and determines the distribution of cations in these types of solid solution oxides as has been discussed recently, s The principal factor is the ratio of the cation diffusivities in the solid solution oxide, and as was seen earlier Do, = 1.15 DMn indicating a tendency for the faster diffusing cobalt to segregate to the outside of the solid solution layer, as indeed is observed. There is a further constraint: since diffusion in the alloy is much slower than that in the oxide, D~al/D°~ --~ 0 as shown earlier, the average ratio of Co/Mn in the scale must be the same as in the alloy. Thus, the compositional path through the cubic oxide inner scale layer will be similar to that depicted as DF in Fig. 10. However, there is a further complication, as illustrated in Fig. 10, in that in the outer portion of the scale where the oxygen activity is higher, a spinel phase is formed. This arises because as the oxygen potential is increased both cations, Mn 2+ and Co 2+, show an increasing tendency to pass into the trivalent state, which requires the creation of additional cation vacancies. As a result, the spinel phase is more stable, and the manganese, which has the greater tendency to exist in the trivalent state, concentrates in this phase. Two possibilities exist: (a) the spinel may form as a discontinuous phase within the oubic oxide, or (b) it may form as a continuous layer. According to the phase diagram, spinel should always be present at the extreme outer surface of the scale, yet this is seldom observed, and it may be that the oxygen potential falls very rapidly in that region of the scale and thus any outer complete spinel layer is very thin. Presumably, the critical factor which determines which of these two possibilities occurs is the flux of manganese out through the cubic oxide. If this is high, then a continuous layer of manganese-rich spinel could form; if it is below some critical value then only a discontinuous phase would be possible. It is clear that the formation of a continuous layer of spinel would be favoured on the more manganese-rich alloys. It is instructive, however, to consider a little further into the likely diffusion paths, as shown schematically in Fig. 10. Consider first the case of a discontinuous layer of spinel. The spinel will begin to form at a position in the scale where the composition curve for the cubic oxide crosses the two-phase region, and will be much richer in the more oxidizable component, manganese. From this position outwards, the compositions of the two phases in the mixed region are fixed and correspond to the phase boundaries (CoMn)O/(CoMn)O + spinel and spinel + (CoMn)O/spinel respectively. The relative amount of each phase depends on the way in which the oxygen activity changes through tlie scale, and the transport properties of each phase. However, it can be anticipated that transport through the spinel is somewhat slower than through the cubic oxide, and as a consequence the latter phase would predominate.

690

F. GESMUNDO,P. NANNIand D. P. WHrrri.E

Thus, there would be a tendency for the diffusion path to bend over to the cobalt-rich side of the diagram, F G H , and it is entirely conceivable that the diffusion path could re-cross the spinel + (CoMn)O/(CoMn)O boundary, and the cubic oxide, now enriched in cobalt, be stable in the outer part of the scale. Figures 2 and 5 show just this type of behaviour with manganese-rich spinel particles buried within the (CoMn)O layer. Turning now to the second possibility and the formation of a continuous layer of spinel, this would have a composition corresponding to point I at its interface with the cubic oxide inner layer (point F). The relative diffusion rates of the various cations in the spinel layer must now be considered. According to Goodenough, aa these spinels have what is called the normal structure, and for Mn-rich compositions, intermediate between CoMnzO4 and MnsO4, the Co is exclusively divalent in the A sites (tetrahedral) while Mn is partly divalent in the A sites and partly trivalent in the octahedral B sites. For manganese contents between 0 and CoMn2Oo Mn is exclusively trivalent in the B sites whilst Co is partly divalent in the A sites and partly trivalent in the B sites. Diffusion rates depend on the site occupied and also on the atomic path followed during diffusion, which in turn is affected by the preference for A or B sites of the individual ions, as measured by the so-called site preference energy, s4 In the composition range CoaO4-CoMn204, according to Armijo, s4 the diffusion of Mn 3+ ions should be difficult since Mn s+ ions have a strong preference for B sites and should make only B - B jumps, with a large strain energy for transport. On the contrary, Co 2+ ions (and probably also Co s+ ions) have a small site preference energy, so they should diffuse via A - B jumps along paths of smaller strain energy. These arguments are therefore in favour of a faster diffusion rate of Co than Mn ions. However, the evidence is not fully conclusive. In the corresponding spinel CoCr204, Cr s+ has a larger preference energy for B sites than does Mn s+ in the CoMn~O4. However, the limited diffusion data available suggest a higher diffusivity for Cr s+ ions. ~ Nevertheless, if the diffusion rate of cobalt ions in the spinel is higher, then the diffusion path in the spinel phase could again bend towards the cobalt-rich side of the diagram (IJ) and indeed cross back into the cubic oxide region. Phase morphologies such as this, with a continuous spinel layer buried in the cubic oxide have also been observed (Fig. 4). It is equally conceivable that the diffusion path could cross backwards and forwards from spinel to cubic oxide phases in some systems, resulting in a more complex multilayering of the scale. The factors which determine whether or not a continuous layer of one of the product phases is formed cannot be established precisely until a more rigorous analysis of the complex, coupled diffusion in these multi-phase scales has been carried out. In the present case, whether the spinel layer is continuous or not is of little consequence, other than that from a fundamental understanding of the behaviour, since the overall growth rate of the scale is virtually unaffected. REFERENCES 1. C. WAGNER,Corros. SeL 9, 91 (1969). 2. D. E. COATESand A. D. DALVI,Oxid. Metals 2, 331 0970). 3. D. P. WHITTLE,B. D. BASTOWand G. C. WOOD,Ox/d. Metals 9, 215 0975). 4. B. D. BASTOW,D. P. WHITTLEand G. C. WOOD,Corros. Sci. 16, 57 (1976). 5. D. P. WinTry, B. D. BASTOWand G. C. WOOD,Trans. Japan Inst. Metals 18, 257 (1977).

High temperature oxidation of Co-Mn alloys

691

6. B. D. BASTOW, D. P. WHITTLE and G. C. WOOD, Proc. R. Sac. 356, 177 (1977). 7. E. B. EVANS, C. A. PHALN1KARand W. M. BALDWIN,JR., J. electrochem. Sac. 103, 367 (1956). 8. G. C. WOOD, I. G. WRIGHT and J. M. FERGUSON,Corros. Sci. 5, 645 (1965). 9. G. C. WOOD and L G. WRIGHT,Corros. Sci. 5, 841 (1965). 10. A. D. DALVI and W. W. SMELTZER,J. electrochem. Sac. 118, 1978 (1971). 11. A. D. DALVI and D. E. COATES,Oxid. Metals 3, 203 (1971). 12. P. MAYERand W. W. SMELTZER,J. electrochem. Sac. 119, 626 (1972). 13. A. D. DALvi and W. W. SMELTZER,J. electrochem. Sac. 121,386 (1974). 14. P. MAYERand W. W. SMELTZFR,J. electrochem. Sac. 121,538 (1974). 15. P. MAYERand W. W. SMELTZER,J. electrochem. Sac. 123, 661 (1976). 16. P. MAYERand W. W. SMELTZER,Oxid. Metals 10, 329 (1976). 17. T. NARITA and K. NISHIDA, Trans. Japan Inst. Metals 17, 849 (1976). 18. A. SCHNEIDERand W. WUNDERLICH,Z. Metallk. 40, 260 (1949). 19. E. AUKRUSTand A. MUAN,Trans. metall. Sac. A.1.M.E. 230, 378 (1964). 20. g . TSIOPLAKISand T. G6DECKE, Z. Metallk. 62, 680 (1971). 21. E. AUKRUSTand A. MUAN, Trans. metall. Sac. A.LM.E. 227, 1378 (1963). 22. C. WAGNER,Z. Elektrochem. 63, 772 (1959). 23. R. A. RAPP, Corrosion 21,382 (1965). 24. Y. IIJIMA, O. TAGUCHIand K. I. HIRANO, Metall. Trans. 8, 991 (1977). 25. J. B. PRICE and Jr. B. WAGNER, JR., J. electrochem. Sac. 117, 242 (1970). 26. E. FRY'r, Oxid. Metals 10, 311 (1976). 27. N. G. EROR and J. B. WAGNER,JR., Phys. Chem. Solids 29, 1597 (1968). 28. B. FISHERand D. S. TANNHAUSER,J. chem. Phys. 44, 1663 (1966). 29. G. J. KOEL and P. J. GELLINGS, Oxid. Metals 5, 185 0972). 30. I. BP,ANSKYand J. M. WXMMER,Phys. Chem. Solids 33, 801 0972). 31. I. BRAlqSKYand N. M. TALLAN,J. electrochem. Sac. 118, 788 (1971). 32. E. FRYT, S. MROWEC and T. WALEC, Oxid. Metals 7, 117 (1973). 33. J. B. GOODENGOGH,Magnetism and the Chemical Bond. Wiley-Interscience, New York 0963). 34. R. ARMIJO, Oxid. Metals 1, 171 (1969).