Vol. 131,No. 4 S I L I C O N IN D R Y O X Y G E N 923

Tiller, This Journal, 127, 2243 (1980). 43. S. M. Hu, Appl. Phys. Lett., 43, 449 (1983). 44. G. Lucovsky, Philos. Mag. B, 39, 513 (1979). 45. A. G. Revesz and G. E. Walrafen, J. Non-Cryst. Solids,

54, 323 (1983). 46. A. G. Revesz and G. V. Gibbs, in "Proceedings of the

Conference on the Physics of MOS Insulator," G. Lucovsky, S. T. Pantelides, and F. L. Galeener, Edi- tors, p. 92, Pergamon Press, New York (1980).

47. P. Rabinzohn, G. Gautherin, B. Agias, F. Rochet, and S. Rigo, in Extended Abstracts of the 5th Interna-

tional Symposium on Passivity," p. 184, Bordeaux, France, May 30-June 3, 1983, Society de Chemic Physique, Paris (1984).

48. V..Q. Ho and T. Sugano, IEEE Trans. Electron. De- vzces, ed-27, 1436 (1980).

49. M. Kh. Karapet'yants, and M. L. Karapet'yants, "Ther- modynamic Constants of Organic and Inorganic Compounds," p. 193, Ann Arbor-Humphrey Science Publishers, London (1970).

50. M. Garric, "Cours de Chimie," p. 220, Dunod, Paris (1971).

The Characteristics of Alumina Scales Formed on Fe-Based Yttria- Dispersed Alloys

T. A. Ramanarayanan,* M. Raghavan, and R. Petkovic-Luton Exxon Research Center, Linden, New Jersey 07036

ABSTRACT

To provide resistance to aggressive environments at elevated temperatures, especially in excess of -1000~ alloys or coatings which develop a-A1203 scales are the best choice. It has been pointed out that the presence of highly stable rare earth oxide dispersoids in high temperature alloys leads to improvements in the corrosion-resistant properties of A1203 scales formed on such alloys. The present study is directed toward developing an understanding of how the properties of A1203 scales formed on Fe-based alloys are influenced by yttr ium oxide dispersoids in the alloy. The Fe-based alloy system selected for the current study consists of -20% Cr, -4.5% A1, -0.5% Ti, and -0.5% Y20~. The oxidation kinetics of the alloy have been established at various oxygen partial pressures in the temperature range 1000~176 The a-A1203 scales which result upon oxidation are observed to be columnar, ultrafine grained, and extremely adherent when thermally stressed. Plat inum markers initially placed on the alloy surface are found at the oxide/gas interface at the completion of oxidation, suggesting that scale growth occurs by exclusive inward oxygen migration. The ultrafine grain size (0.5-1 ~m) suggests that grain boundaries in the oxide scale are the preferred path for oxygen migration. The fine dispersoid particles in the alloy (200-500A) transform to coarse (-0.5 ~m) yttrium a luminum garnet upon incorporation into the AI~O.~ scale, leading to a garnet-saturated scale. It is suggested that the remarkable adherence of the ~-A1203 scales is a consequence of a combination of factors. First, yttr ium doping promotes the development of a fine-grained a-A1203 scale which can effectively relieve oxide growth stresses by diffusional plastic flow. Second, because the alumina scale grows by exclusive inward oxygen transport, growth stresses arising from A1203 nucleation within an existing scale are avoided.

High temperature alloys or coatings designed to resist ag- gressive environments at elevated temperatures should be capable of developing a surface oxide layer which is ther- modynamically stable, slow growing, and adherent. The three oxides which fit the requirement of slow growth are Cr203, SiO~, and A1~O3. Of these, chromia is the fastest growing and alumina is the slowest growing. The thermo- dynamic stability of these oxides is in the order A12Oz > SiO2 > Cr20~. At temperatures exceeding -1000~ chromium oxide scales tend to become unstable; in environments of relatively high oxygen partial pressure, the oxide can va- porize as CrO3 (1), while under highly reducing conditions the oxide can transform to other more thermodynamically stable phases (2).

For applications in hostile environments at temperatures in excess of -1000~ A1203 and SiO2 scales should be preferred to provide corrosion resistance. While the kinetic stages leading to the development of a continuous-surface oxide layer are certainly of importance, once such an oxide layer is established, the most important consideration is how well the oxide layer adheres to the alloy surface. The factors which influence the adhesion of oxide scales are not well understood. ~It has been pointed out by several in- vestigators that small amounts of rare earth oxide disper- soids present in the alloy lead to the formation of a more ad- herent surface-oxide scale (3-17). The literature information in this area was reviewed recently (18-19).

Several mechanisms have been proposed in the literature to explain the dispersoid effect:

1. The oxide scale develops numerous protrusions/pegs which form around rare earth oxide particles in the alloy. These pegs anchor the oxide scale to the alloy.

2. The rare earth oxide particles react with the main oxide scale and form a mixed-oxide layer between the main scale and the alloy. The mixed-oxide layer acts as a "graded seal" improving the stability of the oxide scale under thermal cy- cling conditions.

*Electrochemical Society Active Member. Key words: oxidation, defects, kinetics.

3. The rare earth oxide particles in the alloy act as sites on which vacancies generated by the oxidation process con- dense. The vacanc ies thus do not coalesce on the scale/metal interface, a process which could lead to oxide spallation.

4. The rare earth element dissolves into the scale and sup- presses Al transport in the scale. Thus, growth stresses, oth- erwise developed by the formation of alumina within an ex- isting scale, are decreased.

Sufficient experimental evidence does not exist to confirm the above mechanisms. Also, most of the pub- lished studies point to a combination of mechanisms rather than a single mechanism being operative. Thus, a consist- ent picture of the dispersoid effect has not emerged.

The purpose of the present investigation is to improve our understanding of the influence of rare earth oxide dispersoids on oxide-scale properties. As the first of a pro- posed series of investigations, it was decided to examine the properties of alumina scales developed on an Fe-based alloy system containing yttrium oxide as the dispersoid. The results of this study are contained in the present paper,

Materials and Experimental Procedure The experimental alloy selected for investigation was al-

loy MA956, made by the International Nickel Company. Two types of samples were used; one had a fine-grained (-0.2 ~m) microstructure typical of the as-received condi- tion, while the other had a coarse-grained structure which results upon heat-treating (Fig. 1). The distribution of dispersoid particles is shown in Fig. 2. The particles have been identified by microdiffraction to be YA103. The speci- mens were rectangular, 1 • 1/2 • 1/16 in.; the faces were g r o u n d to 600 grit SiC and c leaned u l t r a son ica l l y in acetone.

The kinetics of oxidation was measured by thermogravi- merry using a Cahn 1000 electrobalance. The elec~robal- ance was attached to a vertically placed quartz reactor tube; the sample to be oxidized was hung from the balance by

924 J. Electrochem. Soc.: SOLID-STATE SCIENCE AND TECHNOLOGY April 1984

Fig. 1. Fine-grained Gnd coarse-grained MA956 samples used in the study

Fig. 2. Dispersoid distribution in alloy MA956

means of a platinum suspension wire. A platinum resist- ance furnace surrounded the quartz tube; the furnace could be moved up and down hydraulically. The oxidizing gas en- tered the quartz tube at the bottom and escaped at the top, where it mixed with an argon stream used to provide an in- ert atmosphere within the electrobalance assembly.

At the start of an experimental run, the sample to be oxi- dized was hung from the electrobalance and the quartz re- actor tube slid into place, the furnace being in the "down" position. The reactor tube was kept under a purging argon stream, and the furnace brought to the oxidation tempera- ture, the sample still being held at room temperature. The argon flow was then replaced by the oxidizing gas or gas mixture, the flow rate being high enough so that mass transport in the gas phase would not be rate controlling. After an interval of -30 min, the furnace was hydraulically lifted up around the sample so that the sample was posi- tioned within the constant temperature zone (-I~ of the furnace. It took approximately 10 rain for the sample to at- tain the experimental temperature. An HP 9845T computer was used to collect and analyze weight-gain data.

The oxidation kinetics were investigated in the tempera- ture range 900~176 The e n v i r o n m e n t s used were undried air, dry oxygen, and CO/CO2 ga s mixtures of vary- ing oxygen partial pressures. The adherence of the oxide scales developed on the alloy surface was inferred by weight changes upon rapidly cooling the sample to room temperature by lowering the furnace and/or temperature cycling by raising and lowering the furnace at specific intervals.

Upon completion of the kinetic runs, the morphology and structure of the oxide scales formed on the alloy surface were examined using scanning electron microscopy, trans-

mission electron microscopy, microdiffraction, and sec- ondary ion mass spectrometry.

Results Most samples were oxidized for periods ranging from 25

to 100h in the thermogravimetric unit. A few samples were oxidized up to 1000h in a separate furnace; kinetic data are not available for these longer-term experiments. The oxide scales were characterized by their excellent adherence; no oxide spallation could be detected upon quenching the sample to room temperature from oxidation temperatures in the range 900~176 even after oxidation times as long as 100h.

Most of the experimental data were taken in the tempera- ture range 1000~176 A parabolic rate was obeyed for ox- idation times as long as 100h. The slope of a plot of the square of the oxidation weight gain per unit area vs. t ime is the parabolic rate constant, kp (g2 cm-4 s-l). The rate con- stant is plotted as a function of temperature in Fig. 3. The data for the fine- and coarse-grained samples fall on the same regression line, the grain size of the alloy having little influence on the oxidation rate. The rate constant can be expressed by the relationship

kp (g2 cm-4 s-~) = 4.2 • l0 s exp (-92,800/RT) [1]

"o

E o

10-11

10-1=

10.~3 u

1 0 -14 6

O

I !

,~ Coarse Grained

o Fine Grained

O

. = .2 x 10 = exp (-92,8001RT)

I I I 7 8 9

l IT x 10 ~

Fig. 3. Parabolic rate constant for ~-AI203 growth as a function of temperature.

Vol. 131, No. 4 A L U M I N A S C A L E S 925

The oxidized samples were fractured under liquid nitro- gen and the fracture cross sections examined by SEM. A co- lumnar fine-grained scale morphology was revealed (Fig. 4). The oxide grain sizes were temperature independent when oxidized in the temperature range 1000~176 for times ranging from 25 to 100h, with grains measuring ap- proximately 1-2 ~m in length and -0.5 ~m across. At 1200~ upon increasing the oxidation time to 300h, the average grain width increased to 0.75 ~m, while after 1000h, the width was approximately 1.25 ~m. Thus, the oxide scale re- mained essentially fine grained even after long oxidation times (Fig. 5). X-ray diffraction analysis showed the scales to be exclusively a-Al~O3.

Fig. 4. Columnar fine-grained morphology of c~-Al~O3 scales

In order to characterize the oxide films further, thin films were prepared and examined using a Philips EM 400 transmission electron microscope. The films were pre- pared by first electropolishing the metal away and then ion thinning the oxide scale. In some cases, only ion thinning was used. The TEM investigations again revealed the fine- grained nature of the oxide scale. A thinned section, viewed perpendicular to the oxide/metal interface, reveals the pres- ence of fine microvoids, -500A in size; the voids occur pre- dominantly at oxide grain boundaries (Fig. 6). Microdiffrac- tion analysis in the TEM confirmed the oxide scale to be a-A1203.

Further detailed examination of the scale was carried out m order to detect the YA103 particles originally present in the alloy. The only yttrium containing phase found in the oxide scale was yttr ium aluminum garnet, Y3AI~O,~. This phase was typically 0.2-0.5 ~m in size, as revealed by the TEM mic rog raph to the left in Fig. 7. The phase was identified by microdiffraction. The [001] diffraction pat- tern is shown on the right in Fig. 7. The zero layer in this pat- tern consists of a square grid of {200} reflections, and the third dimension was calculated from the diameter of the ring to be 12.05~, consistent with the lattice parameter of garnet. The garnet phase was located principally in the vi- cinity of a-A1203 grain boundaries. It is interesting to note the coarseness of the garnet phase when compared with the fine YA10~ particles present in the alloy (Fig. 2). The size difference is further revealed in the transmission electron micrograph in the vicinity of an oxide/metal interface (Fig. 8).

Fig. 6. TEM micrograph of a-AI203 revealing grain boundary microvoids.

Fig. 5. Dependence of c~-Al~O3 grain width on oxidation time. Top: 24h in air, 1200~ Center: 300h in air, 1200~ Bottom: 1000h in air, 1200~

Fig. 7. TEM micrograph and electron diffraction pattern showing the YsAI5Ol~ phase in ~-AI203.

926 J. Electrochem. Soc.: S O L I D - S T A T E S C I E N C E A N D T E C H N O L O G Y April 1984

I I I I I I

T = 9 0 0 ~

0.4 :'-

~ 0.3

0.1

I I I I I I 10 20 30 40 50 60

Time (hours)

Fig. 10. Oxidation kinetics of MA9S6 at 900~

Fig. 8. TEM micrograph in the vicinity of the oxide/metal inter- face showing coarse Y3AIsO12 in the oxide and fine YAIO~ in the metal.

In order to determine the distribution of the Y3A15012 phase in a thick a-A120~ scale, a MA 956 alloy specimen was oxidized for 1000h at 1200~ in air and examined in cross section using a Cameca secondary ion mass spectrometer. Ion Maps for A1 § and Y§ are shown in Fig. 9. The field of view in Fig. 9 is 150 t~m; the positive secondary ions were monitored, sputtered, and excited with a 15 keV beam of 02 +. Because the spacial resolution of SIMS is -0.5 ~m, the fine yttr ium containing oxides in the metal are not re- solved. The Y3A15Ol~ phase in the oxide scale is clearly shown; there is some degree of charging in the Y+ ion map which has exaggerated the size of the garnet phase. It is clearly seen that the fine YA103 particles have dissolved upon entering the oxide scale and reprecipitate as garnet.

A limited number of oxidation experiments were carried out at a lower temperature of 900~ at high oxygen partial pressures . A d i f fe rence in k ine t ics was no t i ced u p o n oxidizing fine-grained MA 956 samples in undried air and dry oxygen. The kinetics were faster in undried air (Fig. 10). Examination of the oxide scales by TEM revealed that in the air-oxidation case, the alumina scale was made up of a mixture of a- and 8-phases, while in the case of dry oxygen, the scale was exclusively the a-phase. A TEM micrograph of the 8-phase and the corresponding [100] microdiffraction pattern are shown in Fig. 11. The microdiffraction pattern shows a rectangular grid of (010) and (002) reflections. The 8-phase is seen to have a faulted structure and probably al-

Fig. 11. TEM micrograph and [100] electron diffraction pattern for ~-AI203 phase.

lows more rapid transport of reactant species, as suggested by the kinetic data shown in Fig. 10.

Discussion Anion migration through Al~03 scale.--Parabolic kinetics

of a-A1203 scale growth suggests that the oxide scale growth is limited by diffusional transport through it. To establish whether oxygen or a luminum is the predominant mobile species in the oxide scale, Pt-wire markers were positioned on the alloy surface prior to oxidation. Upon completion of oxidation, the Pt markers remained at the oxide/gas inter- face, suggesting that scale growth occurred by oxygen transport (Fig. 12).

The d i f fus iv i ty of oxygen in s ing le -c rys ta l l ine and polycrystalline (20-30 tLm grain size) a-A1203 at 1200~176 has been measured by Oishi and Kingery (20). Between 1400 ~ and 1800~ their data for oxygen diffusivity in single crystalline and polycrystalline a-A1203 are given, respec- tively, by

Fig. 9. SIMS ion maps for AI § and Y§ in MA9S6 at 900~

Vol. 131, No. 4 A L U M I N A S C A L E S 927

Fig. 12. Platinum marker study showing the inward growth of c~-Al~03 scale by oxygen diffusion

Do I = 1.9 x 102 exp (-152,000/RT) cm2/s [2]

Do e~r = 2 exp (- l l0 ,000/RT) cm2/s [3]

where Do ~ denotes the lattice diffusivity of oxygen in the single crystal and Do ~ denotes the effective diffusivity (made up of both lattice and grain-boundary contributions) of oxygen in the polycrystalline sample. Experimental data below 1400~ deviated from the above expressions presum- ably b e c a u s e of i m p u r i t i e s c o n t a i n e d in the a-A120~ samples.

If the data according to Eq. [2] and [3] are extrapolated to 1200~ then Do' - 5.3 x 10 -2~ cm 2 s -~ and Do ~f ~ 9.5 x 10 -'7 cm 2 s -~. The effective diffusivity can be expressed in terms of lattice and grain boundary components as (21)

Do ~" = (1 - J)D0 ' + fD0 gb [4]

whe re Do gb r ep re sen t s the grain b o u n d a r y o x y g e n diffusivity and f i s the fractional volume for grain boundary diffusion. The fac to r f can be equated to k6/d where k is a geometrical constant (k was taken to be 2 in this study), 6 is the grain boundary width, and d is the grain width. Thus

k8 k6 Do ~" = Do' + ~- (Do ~b - Do') ~ Do' + ~ Do gb [5]

since Do gb >> Do ~ Using the extrapolated values of Do ef~ and Do' at 1200~ it

is easily seen that for grain sizes in the 20-30 t~m range, the effective oxygen diffusivity in a-A1203 is predominantly de- termined by grain boundary diffusion. Also, in view of Eq. [5], the contribution to grain boundary transport would be even higher as the grain width, d, decreases. It can be con- cluded that in the a-A1203 scales formed on Alloy MA 956, in which d ~ 1 t~m, oxygen transport occurs almost exclu- sively through the grain boundary.

Incorporation of YAI03 particles into inward growing a-Al~03 scale: yttrium doping.--As discussed earlier, the YA10~ particles in the alloy are in the size range 200-500A. Particles in this size range or of this composit ion have not been detected in the a-A1203 scale. Instead, only coarse yt- trium aluminum garnet particles in the size range 0.2-0.5 tLm have been detected. This phase forms predominantly in the vicinity of a-A1203 grain boundaries. Energy dispersive x-ray analysis of the alumina in the grain boundary and bulk regions reveals a strong tendency for yttrium to segre- gate to the grain boundary (Fig. 13). The analysis in the grain-boundary region is averaged over a width of ~200A. Grain boundary segregation of yttrium in alumina has also been reported in a previous study (22).

In view of the above findings, it is suggested that the YA103 particles from the alloy dissolve upon incorporation into the oxide/metal interface region. The yttrium rapidly diffuses along the metal/oxide interface and segregates to the A1~O3 grain boundary region as the oxide grows into the metal. At some point, the solubility product for Y3Al~O,2 is exceeded in the vicinity of some of the A120~ grain bounda-

ries, whereupon this phase is precipitated

6Y + 5A1203 + 9 0 = 2Y~Al~O,2 [6]

The oxygen supply for reaction [6] is provided by grain boundary transport. The suggested mechanism is shown schematically in Fig. 14.

The a-A1203 scale formed on the alloy may be considered to be saturated with Y3A150,2, and the scale will have a cer- tain doped concentration of yttrium, which is predomi- nantly present in the vicinity of grain boundaries. However, the bulk grain is also probably doped with yttrium at lower concentration levels. It has been proposed that doped yt- trium tends to reduce A1 transport through bulk ~-A1203 grains (23, 24). Such an effect cannot be expected if yt tr ium occupies the A1 lattice site in a-A1203 as a trivalent dopant with zero effective charge. However, E1-Aiat and KrSger (25) argue that yttrium might act as a donor YAI", a yttrium defect with a positive effective charge. In the present dis- cussion, it is considered that yttrium might act as a donor or an accep to r in a cco rdance wi.th the fo l lowing defec t equilibria

3 2Y3A150,2 + --~ O5 = 6YA1 + 3 O{' + 9 Oo x + 5A1203 [7]

3 2Y3A15012 = 6YAI' + 3Vo'" + 6 Oo x + -~- O~ + 5A1~O3 [8]

The equilibria [7] and [8] imply that defects generally pro- posed for bulk grains also prevail in the vicinity of grain boundaries. This is appropriate because the high diffu-

Fig. 13. Grain boundary segregation of yttrium in a-AI203

928 J. Electrochem. Soc.: SOLID-STATE SCIENCE AND TECHNOLOGY April 1984

OXIDIZING ENVIRONMENT

OXIDIZING ENVIRONMENT

Fig. Y~AI~O~2 a-A1203.

14. Mechanism for precipitation in

O

c~-Al=O~ Grain

-'6- e r r ~

YAIO~

O O

1 @ @

@ @

O O O

c~-AI=O~ c~-Al=O= Oxide Grain Grain

@ i ~Y,AI~01=

o o o o ~ \ 7 ~ Alloy

YAIO~

sivi ty path for rapid t ransport in the v ic in i ty of an ~-AI~O~ grain bounda ry is ~100-150A wide (26). Wang and Kroger (27) have also used defect descr ip t ions general ly used for bu lk grains for the grain boundary region.

Alpha A120~ has a high bandgap of ~9.9 eV (28, 29); the concent ra t ions of nat ive defects fo rmed by electronic or ionic disorders is ex t r eme ly small. Thus, for equ i l ib r ium [7], one can use the e lec t roneutra l i ty condi t ion

[YA,'] = 2[O~'] [9]

where [ ] denotes concentrat ion. Combin ing the equil ib- r ium constant , K~, for react ion [7] wi th Eq. [9],

K, = 64 [O/'] 9 Po2 -3/2 [10]

Thus [O/'] ~ Po2 '~g. S ince the diffusivi ty of oxygen will have the same oxygen partial p ressure dependence as the major oxygen defect (30)

D o gb = Do *gb Po~ 1/6 [11]

where Do gb is the oxygen diffusivi ty in the v ic in i ty of a-AlzO3 grain boundary and Do *gb is the va lue at 1 a tm oxygen par- tial pressure.

Use of similar a rguments to react ion [8] a long with the neutra l i ty condit ion, [YA,'] = 2[Vo"] leads to

K , = 64 [Vo"]9 Po2~/~ [10.1]

Do g~ = Do *gb Po2 -'/~ [11.1]

Where KH is the equ i l ib r ium cons tant for react ion [8]. Quant i ta t i ve descript ion o f ox idat ion k ine t ics . - - In the

fo l lowing analysis, the a-AI~O~ scale is cons idered to consis t of co lumnar grains having a square cross sect ion of side d. The oxide scale is p r e s u m e d to grow exc lus ive ly by grain- boundary t ranspor t of oxygen. The wid th of the bounda ry reg ion where rapid t ransport occurs is des igna ted 8. The ox- ygen flux dur ing oxide scale g rowth is cons idered to occur exc lus ive ly in a di rect ion parallel to the grain boundary . The oxygen flux per uni t area of the grain bounda ry is then g iven by

Jo (tool cm -~ s - ' ) = - B o C o ( - ~ - )

1 B o R T C o ( O l n ) 1 ( O l n p = - -2 \--~x Po~ "= - -~ Dog'Co ~--~x 02] [12]

where Bo is the mobi l i ty of oxygen (velocity per uni t gradi- en t of chemica l potential), Co is the average oxygen concen- t ra t ion in the oxide in mol cm -~, ~o is the chemica l potent ia l of o x y g e n in the oxide, and x the d is tance coord ina te per- pend icu la r to the oxide /meta l interface. The oxygen part ial p ressure at the oxide/gas interface is des igna ted as Po2o and at the oxide /meta l interface as Po21. Per uni t area of the ox- ide scale, the area occupied by grain bounda ry equals ~28/d. Thus, the flux, Jo ' of o x y g e n per uni t area of the ox- ide scale is g iven by

Jo' = - DogbCo a d -~ ( ~ ln P \--~-x ~] [13]

In tegra t ing across thickness , hx, of ox ide scale

f Ro20

Jo'hX = Co 8 d -1 Dogbd In Po2 [14] J Po2 i

U s i n g t h e r e l a t i o n s , Jo ' = 1/16A ( d h w / d t ) a n d hx = Aw/(Poxid~foA), where poxide is the dens i ty of ~-A1203, Aw is the oxida t ion weight gain measu red thermogravimet r ica l ly , A is the cross-sect ional contac t area be tween ox ide and metal , and fo is the weigh t f ract ion of oxygen, Eq. [14] can be rewri t ten as

A2 AwdAw = [16 Poxid~foCo 8 d -I Dogbd In Po~] dt J PO2 i

[15] In tegra t ing and rear ranging

- - = [59.78 Co ~ d -1 Dogbd In Po2]t [16] o2 i

w h e r e b y the parabol ic rate cons tant (Fig. 3), k,, is g iven by

7 ~ O20 kp = 59.78 Co 8 d -1 Do gb d In Po~ [17]

o2 i

In in tegra t ing Eq. [17], the var ia t ion of Do gb wi th oxygen partial pressure mus t be taken into account . I f y t t r ium acts as a donor and leads to the format ion of O/' defects, then Eq. [11] holds. Co~mbining Eq. [11] and [17]

kp = 358 .68 (~d-lCoDo*gb(Po201/0 - Po211/6) [18]

At l l00~ po2 ~ ~ 5 x 10 -28 atm. The lowest env i ronmen ta l oxygen part ial pressure used in the present inves t iga t ion is ~10 -~4 atm. Thus Po2 ~ > > Po~ i~18 so that

kp = 358.688d-lCoDo*SbPo2 ~ [19]

Thus kp is s t rongly d e p e n d e n t on the env i ronmen ta l oxy- gen part ial pressure.

I f y t t r ium acts as an acceptor and creates oxygen vacan- cies as the p r e d o m i n a n t defects, then, combin ing Eq. [11.1] and [17]

kp = 358.688d-lCoDo*Sb(Po2 i-lIe - Po~ ~ [20]

In this case, s ince Po2 ~-l~e > > Po2 ~176

kp = 358.688d-lCoDo*gbPo2 i-~/~ [21]

Thus, kp should be i n d e p e n d e n t of the env i ronmen ta l oxy- gen part ial pressure.

To test whe the r Eq. [19] or [21] applies, the var ia t ion ofk~ wi th o x y g e n part ial pressure has been measu red at 1100 ~ and 1200~ Within l imits of expe r imen ta l error, kp is found to be i n d e p e n d e n t of the o x y g e n partial pressure. Typica l data at l l00~ are p resen ted in Fig. 15. It is, therefore , sug- ges ted that oxygen vacanc ies ra ther than interst i t ia ls are re- spons ib le for grain bounda ry oxygen transport . Us ing ex- pe r imen ta l va lues for kp and the average grain width , d, the

Vol. 131, No. 4 929

..--- tl

o %

10-14 10-13

10-11 , , i ' , 1 , , , I

10-12

10-13

�9 �9

�9 �9

10-12

I I I I I I I I I

e a �9

�9 �9

I I I I l l l l l I I I I I l l l l I I I I I l l l l

10-13 10-12 10-11

OXYGEN PARTIAL PRESSURE. ATM.

A L U M I N A S C A L E S

10-11 10-2

T = 1100~ AIIoy-MA 956

, , , , I I i i i I

Fig. 15. Variation of kp with oxygen partial pressure.

following expression has been evaluated for the product 8Do*~b

8Do*gb(cm 3 S -I) = 4.92 x 10 -2 exp (-133,940/RT) [22]

Equation [22] holds for 1 atm oxygen partial pressure. For other oxygen partial pressures

8Do sb = 8Do*gbPo2 -1/3 [23]

The activation energy term in Eq. [22] is made up of contri- butions from the temperature dependence of 8 and the tem- perature dependence of Do *gb.

The present results for 8Do *Sb may be compared with values for 8Do ~b calculated from Oishi and Kingery (20)

8I)o gb = 1.4 x 10 -3 exp (-llO,O00/RT) [24]

and values measured by Reddy (31)

8/:)0 gb = 6.15 exp (-142,000/RT) [25]

However, no oxygen partial pressure dependencies are available for these values.

A mechanism for microvoid formation in a-Al20, grain boundaries.--The production of voids in oxide scales growing under an oxygen partial pressure gradient is by no means a new phenomenon. Pores have been observed, for example, in oxidized scales of Fe304 (32), F e e (33), NiO (34), and A1203 (35, 36). In Fe304, Fee , and NiO, the pores were large enough to be detected by optical metallography. Evi- dence of pore formation in CoO placed under an oxygen partial pressure gradient was obtained by Yurek and Schmalzried (37). A pellet of CoO, initially equilibrated at Po~ = 0.2 atm at 1200~ was placed in a Po~ gradient so that one face was at 0.2 atm and the other face was at 10-9 atm. After 1.5h at 1200~ the crystal developed - 5 volume per- cent (v/o) porosity. The authors suggested that the pores are formed as a result of cobalt vacancy diffusion from the high Po~ side and subsequent precipitation at the low Po2 side.

The above examples suggest that the formation of pores in oxide scales placed under a Po~ gradient is related to atomic defects in the crystal and occurs by vacancy conden- sation. In their work on A1203 scales formed on Ni-Cr-A1 al- loys, Smialek and Gibala (35) tentatively proposed a model based on oxygen-vacancy supersaturation and condensa- tion to account for void formation. The voids were present both within the grain and in the grain boundary. The model they proposed is schematically indicated in Fig. 16. Assum- ing oxygen vacancies to be the predominant migrating de- fect, the solid line in Fig. 16 represents the oxygen vacancy gradient within the scale at a scale thickness, hx. Local e q u i l i b r i u m is a s sumed at the meta l /ox ide and the oxide/gas interfaces. When the~oxide has grown to a thick- ness, ~ ' , the new oxygen vacancy gradient is represented by the dotted line. This implies a supersaturation of oxygen vacancies in the growing scale, which condenses out as voids in establishing the new equilibrium oxygen vacancy gradient.

Elegant though this mechanism is, it suffers from a flaw; if the oxygen vacancy concentration at the oxide/metal in- terface exceeds the equilibrium value, the interface cannot move, since this would imply an oxygen concentration be- low that required for oxide/metal equilibrium. To maintain

the oxygen concentration (and, therefore, the oxygen va- cancy concentration) at the oxide/metal interface at the equilibrium value, the arriving oxygen flux must over- come any oxygen vacancy supersaturation within the oxide.

The following mechanism is tentatively proposed to ac- count for grain boundary voids in a-A1203 scales of the pres- ent study. As yttrium dissolves into the a-A1303 scale as the scale grows inward, the oxygen vacancy concentration in the vicinity of oxide grain boundaries increases by virtue of reaction [8]. If Schottky defects are assumed in a-A1203, the increase in oxygen vacancy c o n c e n t r a t i o n m u s t be accompanied by a decrease in the a luminum vacancy con- centration in the vicinity of grain boundaries by virtue of the Schottky defect equilibrium

2VAI" ' + 3Vo'" = null [26]

Ks = [Vo"] 3 [VAI" ,]2 [27]

where Ks is the Schottky defect equilibrium constant. Thus, in attaining the lower a luminum vacancy concentra- tion required by the increase in the oxygen vacancy con- centration, the excess a luminum vacancies must condense out 'in the grain boundary region; this process generates grain boundary microvoids.

Adherence of a-Al303 scales.--Of the effects produced by oxide dispersions on the properties of oxide scales on high temperature alloys, the most significant from a techno- logical point of view is the adherence of the oxide scale. In the sense in which the term "adherence" is used in most of oxidation literature, it refers to the resistance of the scale to spallation under the influence of stresses. In isothermal oxidation, one is mainly concerned with growth stresses. In alloy components used in various high temperature pro- cesses, additional stresses arise which are thermally in-

F

O x y g e n O x i d e

= AX =

I I I I I

A l l o y

A X ' =

Fig. 16. Schematic oxygen vacancy concentration profile in growing cz-AI203.

930 J. E lec t rochem. Soc.: S O L I D - S T A T E S C I E N C E A N D T E C H N O L O G Y A p r i l 1984

duced (due to temperature oscillations) and/or mechani- cally induced.

It was pointed out earlier that several mechanisms have been proposed in the literature as contributing to rare earth oxide dispersoid-induced oxide-scale adherence. The most widely accepted mechanism is that the oxide scale devel- ops protrusions (pegs), which form around dispersoid particles in the alloy; these protrusions anchor the scale to the alloy. Giggins and Pettit have suggested the existence of pegs in Co-Cr-A1-Y and Ni-Cr-A1-Y systems (38). Pegs have been observed especially in Hf containing alloys fo rming a lumina , in which i n t e rna l l y oxidized HfO~ particles were present in the alloy subsurface region. Whit- tle et al. (39) have presented evidence for peg formation in Co-Cr-A1-Hf alloys in which the Hf was oxidized to HfO2 particles. A more dramatic case of pegging has been re- cently observed by Hindam and Whittle (40) in Fe-10A1-Hf alloys in which subsurface HfO2 particles were present. They have suggested a mechanism for peg formation ac- cording to which HfO2 particles in the A12Oz scale act as short-circuit paths for oxygen transport. In regions of the scale where HfO2 particles are present, especially close to the alloy/scale interface, more rapid AI~O~ growth occurs, leading to protrusions which envelop HfO2.

A detailed examination of the oxide/metal interface was made in the present study to look for any evidence of pegs around dispersoid particles. A TEM micrograph in the vi- cinity of the oxide/metal interface is shown in Fig. 17. The dispersoid particles are clearly seen in the metal in the vi- cinity of the oxide/metal interface; there are no pegs around the particles. The absence of pegging is also clearly re- vealed by the ion maps for A1 § and Y+ generated by second- ary ion mass spectrometry (Fig. 9). Thus, pegging must be ruled out as a mechanism for adherence in the case ofyttria- dispersed alumina forming alloys.

The second mechanism proposed in the literature is that a mixed-oxide phase forms by reaction of the main oxide scale with the rare earth oxide, and this phase is present as a thin layer between the main scale and the alloy. The mixed- oxide phase is supposed to have a thermal expansion coefficient which is intermediate between that for the main scale and the alloy so that a"graded-seal" effect is pro- duced. There is no sound evidence of such a mixed-oxide layer in any reported investigation. In a study of the oxida- tion of a Cr-5 Wo Y~O3 alloy, Seybolt (12) provides tentative evidence for a thin YCrO~ intervening layer. However, the existence of such a layer has never been confirmed. No mixed-oxide layer could be detected between the a-A1203 scale and the alloy in the present investigation.

The third mechanism is based on the argument that va- cancies generated in the alloy when an element such as A1 is preferentially oxidized can condense at metal/dispersoid interfaces. These vacancies, therefore, are not available to coalesce on the metal/scale interface, a process which can

Fig. 17. TEM micrograph in the vicinity of oxide/metal interface revealing absence of micropegging.

lead to loss of scale adhesion. In their work on the mecha- nism of oxide-scale adherence on Fe-25Cr-4Al-(Y or Sc) al- loys, Tien and Pettit (41) suggested that internally oxidized Sc203 and Y203 particles act as sinks for vacancies gener- ated during oxidation. In a comparative study of the oxida- tion of Ni3A1 and NizA1-0.5Y alloys, Kuenzly and Douglass (42) obse rved c o n d e n s a t i o n of vacanc ies at the alumina/alloy interface in the case of the yttrium-free alloy. The vacancies were thought to form as the result of a "Kirkendall effect" in the alloy subsurface region upon preferential remoVal of A1 by oxidation, Voids were absent in ~he yttrium containing alloy and are thought to have con- densed around internally oxidized Y203 particles. However, no microstructural evidence for void condensation around Y203 particles was presented. Also, the present authors are not aware of published literature data which confirm the presence of voids around dispersoids. Examination Of dispersoid particles by TEM in the present study did not show any evidence of void condensation.

In the opinion of the present authors, the improved ad- herence of ~-A120~ scales on Y20~-dispersed alloys is caused by the chemical doping of yttrium into the oxide scale and consequent effects on the relief of growth stresses. It is well documented that during the oxidation process of a metal growth stresses arise. The subject has been reviewed by Stringer (43). It is considered that yttr ium doping leads to growth stress relief in two ways: (i) by effecting the forma- tion of a fine-grained oxide scale, and (ii) by retarding A1 transport in the scale.

It has been shown that in polycrystalline alumina having grain sizes less than ~15 tLm, diffusional creep is the most important mechanism for plastic deformation (44). The generalized expression for the creep strain rate, ~, is given by (45, 46)

14a f +_ SDACb/d!_(Do'_+ ] ~ = kTd 2' L3(DA, ' + V6DAlgb/d) + 2(Do' + ~ d ) J ~ [28]

where ~ is the volume of an alumina molecule, the super- scripts on D represent the diffusion pa th (lattice or grain boundary), the subscripts indicate the diffusing species, and o- is the stress under which deformation occurs. It is seen from Eq. [28] that ~ is extremely sensitive to the grain size. Thus, in the ultrafine-grained a-A1203 scale obtained in the present study, plastic deformation can occur with greater ease. The grain size of a-AI~O~ scales grown on dispersoid-free Fe-25Cr-4A1 alloys for 1000h at 1200~ is ap- proximately a factor of four larger (47). Thus long-term sta- bility of the oxide scale can be improved by having a fine grain size. Yttrium doping can lead to the retardation of grain growth kinetics in two ways: by forming grain bound- ary microvoids which, analogous to particles on a grain boundary, can pin the boundary, and by segregating to the grain boundary and decreasing grain-boundary velocity (48).

That relief of growth stresses by plastic flow even when thermally induced stresses are present can improve scale adherence has been shown in an oxidation study of copper carried out some years ago by Sartell et al. (49). They showed that when copper is oxidized at a temperature just below 700~ and cooled to room temperature, the oxide scale readily flakes off. However, when the oxidation is carried out at 870~ the oxide scale remains quite adherent upon cooling to room temperature. The authors argued that in the latter case, growth stresses are relieved by plastic flow at 870~ so that even at a higher level of thermal stresses, the oxide scale formed at the higher temperature is more adherent. Tylecote has shown (50) that above 700~ copper oxide, Cu~O, undergoes a transition from elastic to plastic behavior.

The second effect of yttrium doping is to suppress /tl transport through a-Al~O3 scales. In their work on Fe-Cr-A1 and Fe-Cr-A1-Y alloys, Golightly et al. (23, 24) suggested that Y in some unknown fashion decreased the transport o f A1 in A120~. In the yttrium-free alloy, according to these au- thors, oxygen diffusing inward through ~-A12Oz grain boundaries reacts with A1 diffusing outward through ~-A1~O3 grains, forming new A12Oz within an existing scale and generating compressive growth stresses. In an effort to relieve these compressive stresses, the oxide tends to grow

Vol. 131, No. 4 ALUMINA SCALES 931

laterally, resulting in a convoluted morphology, and be- comes detached from the alloy at the convolutions. Such detachment leads to loss of oxide scale adhesion.

Our preliminary work on Fe-25Cr-5A1 alloys also shows a convoluted oxide-scale morphology (47). Such convolu- tions are absent in the oxide scale on MA956. Thus, in agree- ment with Golightly et al. (23, 24), the present authors also feel that yttrium doping in a-A1203 suppresses A1 transport and avoids the build up of compressive stresses otherwise generated by Al~O3 formation within existing scale.

If it is assumed that Schottky disorder prevails in a-A1203, then a mechanism can be proposed for the suppression of A1 transport as a result of yttrium doping. As explained pre- viously, in fine-grained ~-A1203 grain-boundary oxygen transport predominates, so that one is dealing with a rather low flux of aluminum. In accordance with our proposed defect reaction [8] for yttrium doping, the concentration [Vo] is increased by yttrium, especially in the vicinity of A1203 grain boundaries. In accordance with Eq. [27], the concentration of A1 vacancies, [VA(' '] is correspondingly reduced. Since oxygen transport occurs through grain boundary, any reaction between A1 and O to form Al~O3 within an existing oxide scale must occur in the grain boundary region. Such A1203 formation is prevented by de- creasing [VAI" '] in the grain boundary region, which corre- spondingly decreases the flux of A1 to this region.

Concluding Remarks In the present study, the influence of yttrium oxide

dispersoids in Fe-base alloys on the microstructure and ad- herence of a-A1203 scales formed on such alloys has been emphasized. The oxide scale grows by inward oxygen grain boundary transport; the yttr ium oxide dispersoid particles dissolve upon incorporation into the scale, and yttrium seg- regates to the a-A1203 grain boundary; in consequence, it is suggested that grain growth and Al transport in a-A1203 are suppressed. These effects in turn can relieve and/or de- crease growth stresses in a-A1203 and improve oxide-scale adherence.

In view of the immense technological importance of a~A1203 scales, it is suggested that more work should be carried out both on the defect structure and doping charac- teristics of AI~O3. Some papers in the literature suggest that whatever the chemistry of the dispersoid used, there is an improvement in the adherence of the ~-A1203 scale formed on the alloy. While this may indeed be so, the do,ping effects of elements from different groups of the periodic table are clearly different. Thus, there is a need to understand more clearly the role of dispersoid chemistry in the alloy on oxide-scale properties.

Acknowledgment The authors gratefully acknowledge the efforts of J.

Mumford throughout this experimental study. They wish to thank D. P. Leta for performing the SIMS work and Lydia Massey for her skillful assistance in preparing the manuscript.

Manuscript submitted Jan. 9, 1983; revised manuscript received Nov. 22, 1983.

Exxon Research Center assisted in meeting the publica- tion costs of this article.

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