NASA Contractor Report 175051

Grain Boundary Oxidation and an Analysis

of the Effects of Pre-oxidation on

Subsequent Fatigue Life

i (NASA-CR-175051) GBAI N BOUNDAE Y OXIDATION N86-2165U

AND AS ANALYSIS Of THE EFIECTS OF

IPBE-OXIDATION ON SUBSEQUENT FATIGUE LIFE

i (Syracuse Univ., N. Y.) 38 p HC A03/MZ A01 Unclas

v CSCL 11F G3/26^ ' 05655

Yoshiki Oshida and H.W. Liu

Syracuse University

Syracuse, New York

February 1986

Prepared for

Lewis Research Center

Under Grant NAG 3-348

NASA

National Aeronautics and

Space Administration
Grain Boundary Oxidation and an Analysis of the Effects of

Pre-oxidation on Subsequent Fatigue Life

Yoshiki Oshida and H.W. Liu1

Syracuse University

Syracuse, New York 13210

Abstract

A mechanistic study on the effects of pre-oxidation on subsequent fatigue life

was conducted. The surface oxidation and grain boundary oxidation of a nickel-base

superalloy (TAZ-8A) were studied at 600 to 1000°C for 10 to 1000 hours in air. The

surface oxides were identified and the kinetics of surface oxidation was discussed.

The grain boundary oxide penetration and morphology were studied. The pancake

type grain boundary oxide penetrates deeper and its size is larger, therefore, it is

more detrimental to fatigue life than the cone-type grain boundary oxide. The oxide

penetration depth, a , is related to oxidation temperature, T, and exposure time, t,

by an empirical relation of the Arrhenius type. The effects of T and t on the statistical variation of a were analyzed according to the Weibull distribution function.

Once the oxide is cracked, it serves as a fatigue crack nucleus. The statistical

variation of the remaining fatigue life, after the formation of an oxide crack of a

critical length, is related directly to the statistical variation of grain boundary

oxide penetration depth.

Research Assistant Professor and Professor of Mechanical Engineering.
I. INTRODUCTION

Fatigue life is often shortened at elevated temperatures.

The shortened fatigue life is usually attributed to oxidation

and/or creep. Creep deformation creates grain boundary voids,

which is said to weaken the material. In conjunction with the

applied cyclic fatigue load, the grain boundary voids may accelerate

the initiation and the propagation of fatigue cracks, and thereby

shorten fatigue life. The subject of creep-fatigue interaction

has been studied extensively (1,2,3,4,5).

a

In a vacuum of 10" torr, the effects of cyclic frequency

and test temperature on fatigue lives of A286 at elevated temperatures are minimal in comparison with those in air (6,7).

Therefore oxidation is more detrimental, at least for some

materials, than creep.

This paper deals 'with surface and grain boundary oxidation

and the analysis of the effects of the .grain boundary oxidation on fatigue life.

Gibb's free energies of metal oxide formation are negative.

No metal or alloy is stable when exposed to an oxidizing environment. Hence a high oxidation resistance of an alloy implies a

low oxidation rate. A low oxidation rate is generally due to

the slow diffusion rate through the oxide film.

^-v

Because grain boundary is a path of rapid diffusion, grain

boundary oxidation rate is higher and grain boundary oxide

penetration is deeper than the surface oxidation. When an oxide

reaches a critical size at an applied stress level, the oxide
will fracture (8). The oxide crack may serve as a nucleus of

a fatigue crack, and will grow by the subsequent cyclic fatigue

load. Therefore, an oxide crack may shorten fatigue life, and

grain boundary oxide penetration depth can serve as a "measure"

of high temperature fatigue damage.

McMahon and Coffin (9) have shown ruptured oxide film to act

as a fatigue crack nucleus. Shida et al. (lO)andStott et al. (11)

have studied grain boundary oxidation. In this paper, a mechanistic

study on the effects of pre-oxidation on subsequent fatigue life

was conducted. The surface and grain boundary oxidations of a

nickel-base superalloy (TAZ-8A) were investigated. The surface oxides

were identified by an electron-probe x-ray energy dispersion spectrometry analysis, oxide weight gain and weight loss analysis, and

reflection electron diffraction method. The three-dimensional

grain boundary oxide morphology was studied. The grain boundary

oxidation kinetics was analyzed, and the statistical distribution

of the grain boundary oxide penetration depth was studied. The

effects of oxide penetration depth and oxide crack on subsequent

fatigue life are discussed.
II. SURFACE OXIDATION

A nickel-base superalloy (TAZ-8A) was subjected to high

temperature oxidation in air under the stress-free condition.

The chemical composition of TAZ-8A is listed in Table 1. The

oxidation temperatures were 600, 700, 800, 900, and 1000°C,

controlled within ± 8°C. The oxidation time varied from 10 to

1000 hours.

2

The average weight gain per unit surface area, AW+0 (mg/cm ),

due to both surface and grain boundary oxidations was obtained

for each disk coupon, 15 mm in diameter and 5 mm in thickness.

At a given temperature, a linear relationship between the square

of AW+Q and oxidation exposure time, t, was found, Figure 1.

The slopes of the lines in the figure are known as the parabolic

rate constant, K . The data in the figure indicate that the

oxidation of TAZ-8A obeys the parabolic rate law as many other

superalloys and refractory materials do. The parabolic oxidation

rate implies that the oxidation is controlled by diffusion.

Figure 2 shows the parabolic rate constant versus inverse

absolute temperature of TAZ-8A and several other superalloys.

The data of TAZ-8A seem to follow two line segments similar to

Mar-M200. Additional studies are needed to verify whether these

two line segments indicate two different oxidation kinetics in

the high and low temperature regimes.

The measured weight gain during oxidation, AW+Q, reflects

the amount of oxygen that reacted with the substrate to form an

oxide. On the other hand, the weight loss of the substrate due
to the oxidation can be measured by descaling. For rapid

descaling, the oxidized sample was immersed into an acid aq.

solution (50% HC1, 49% HgO, 1% H202) at 70°C (12) for three

minutes and subsequently in a solution of 30% HN03, 10% R2

Q

2'

10% H3P04 and 50% CHgCOOH at 90°C (12) for three additional

minutes. The descaling process was followed by an examination

with an optical microscope. If the amount of oxide removal was

not enough, this process was repeated. The final descaling

process is accomplished at a much slower rate by cathodic

descaling to avoid the removal of the substrate; oxidized sample

at cathode; lead plate at anode; 5% HgSO. electrolyte; and

o

1 mA/cm current density (13).

The measured weight loss, AW_M, reflects the amount of

metal loss due to oxidation. The weight gain to weight loss

ratio, AW+Q/AW ,., should be equal to the atomic weight ratio

of oxide if the oxide is stoichiometric. The measured ratios

are 0.32, 0.44, and 0.49 for oxides formed at 600°, 800°, and

1000°C, respectively. The atomic weight ratio for Cr203 is

0.462. Similar ratios of other oxides are 0.273 (NiO), 0.889

(A100,), and 0.393 (NiCr.,0,). « O & 7t

The reflection electron diffraction pattern identified that

the outermost layer of the surface oxide formed at 800°C for

960 hours consists mainly of Cr203 (corundum type of sesquioxide)

with a small amount of spinel type oxide possibly such as

NiO-Cr203 or NiO-(Cr,Al)203< Some of chromium ions in Cr203

might be substituted by aluminum ions to form (Cr,Al)203.
Figure 3a shows the picture of a cross-section, under an

optical microscope, of a test coupon oxidized at 1000°C for

500 hours. Figure 3b is the cross-section under a scanning

electron microscope. The microstructures of the substrate shows

the hardening precipitates (Y'-phase or Ni^Al) in y-phase matrix

with relatively large carbides. After the sample was oxidized,

the decomposition of y'-phase and the chromium depletion were

observed in a substrate layer next to the substrate-oxide interface. By increasing the oxidation temperature and prolonging

oxidation time, these microstructural changes due to high temperature oxidation became remarkable.(e.g., the thickness of

100 urn of Y'-phase decomposition zone was observed after

1000°C x 960 hours oxidation.)

Figure 4 shows the line profiles for NiK^, CrKa and AlKa

constructed from the electron-probe spot analyses by an x-ray

energy dispersion spectrometer. The spot analyses were made

at five locations on each one of the three test coupons oxidized

at 600, 800 and 1000°C. The five locations are: (1) in the

substrate far away from the substrate-oxide interface, (2) and

(3) next to the substrate-oxide interface, one in the substrate

and the other in the oxide, (4) at the mid-section of the surface

oxide layer, and (5) at the outermost layer of the surface oxide.

The oxide layer at 600°C is thin; the spots (4) and (5) merged

together.

The major findings from the spectrometric analysis are as

follows:
i) The oxides of aluminum, nickel and chromium are present

in the surface oxide layer. -However, at 600°C, nickel oxide

is the dominant component, at 800°C and 1000°C chromium oxide

is the dominant one.

ii) At 800 and 1000°C, the chromium concentration increases

continuously from the substrate to the outermost layer of the

surface oxide with a minor chromium depleted zone in the substrate

next to the substrate-oxide interface.

iii) At 800 and 1000°C, the concentrations of the chromium oxide

at the outermost layers of the surface oxides are much higher

than that of either the aluminum or nickel oxides.

iv) The concentration of nickel decreases continuously from

the substrate to the outermost layer of the surface oxide.

This spectrometric analysis indicates that the oxides formed

at 800°C and 1000°C in air are mainly Cr00o or (Cr,Al)0Or} with a £ w £t O

much less amount of NiO-Cr203 and/or NiO- (Cr, AlKOo. This is

consistent with the measured weight gain/weight loss ratio, if

the relative amounts of the various oxides are taken into consideration. For example, at 1000°C, the relative portions of

the chromium, nickel, and aluminum oxides are 2:1:1. This will

give a weight gain to weight loss ratio of 0.52, which is very

close to the measured ratio of 0.49. The activation energy in

Figure 2 in the high temperature regime is very close to the

activation energies of the diffusion of chromium in either Cr203

or NiO-Cr203 (14).
At 600°C, the spectrometric analysis indicates that nickel

oxide is the dominant component, and the ratio (AW+Q/AW . ) is

0.32, which is between the values of NiO (0.273) and NiO-Cr203

(0.393). Therefore, in the low temperature regime, it is speculated that the oxidation kinetic is controlled by the diffusion

of nickel atoms through the oxide layer.
III. GRAIN BOUNDARY OXIDATION

Figure 5 shows a picture of an oxidized surface. The

grain boundary oxide penetrated much deeper than the surface

oxide. A cracked grain boundary oxide can serve as a nucleus,

which will grow by the subsequent fatigue cycles, and will

shorten the fatigue life. Fatigue crack growth rate is related to AK, and AK is related to the size and shape of the

micro-crack. Therefore the shape and the depth of oxide penetration at grain boundaries are important factors for fatigue

life prediction.

The three-dimensional shapes and sizes of the grain

boundary oxides were traced optically. First a microhardness

indentation was placed close to the oxide to be studied, and the

size and shape of the oxide on the surface were traced. After

the initial measurements, the surface was mechanically polished.

The thickness of the surface layer polished off can be monitored

by the change in the size of the hardness indentation. After

polishing, the size and shape of the oxide were traced again.

This process was repeated to trace the three-dimensional morphology of the grain boundary oxides. As shown in Figures 6

and 7, two types of grain boundary oxides have been observed;

the cone type and the pancake type. The pancake type was

larger, seemed to penetrate deeper, and due to its shape and

size, would have a higher value of stress intensity factor and

would reduce fatigue life more than the cone type.
The aspect ratio a/b varies widely. On the. average, a/b

is close to one (15). However, the ratios for the deeper

oxides are much larger than one,often in the range of 10 to

20 or more (see Figure 5). The depth of penetration by grain

boundary diffusion depends on the misorientation of the two

neighboring grains (16, 17). Conceivably, the depth of grain

boundary oxide is also dependent on the misorientation.

Oxide depth is also a function of oxidation temperature

and the length of the time of exposure. At a given temperature

and exposure time, the oxide depth varied widely from one grain

boundary to another. The deepest grain boundary oxide crack

will control the amount of the reduction in fatigue life.

Therefore, the extreme value, i.e., the deepest oxide penetration should be analyzed. In this paper, the grain boundary

oxide depth as a function of oxidation temperature and exposure

time, grain boundary oxidation kinetics and the statistical

distribution and the extreme value of grain boundary oxide

depth were studied.

The statistical distribution of grain boundary oxide penetration was studied at exposure times from 100 to 1000 hours

and at temperatures of 600, 800 and 1000°C. Cylindrical test

coupons were oxidized in air. The oxidized disk coupons were

sectioned, each sectioned surface was examined under an optical

microscope, and the maximum grain boundary oxide penetration

depth, a ., of the sectioned surface "i" was measured. Then a

mi

thin layer of the coupon approximately 80 pm thick was ground
off, the new surface was polished, and another a . of the new

mi

surface was measured. This' process was repeated many times

(often 12 times) for each test coupon to collect enough data

for the statistical analysis'. The grain size of the nickel

alloy is between 50 to 80 um. After a removal of 30 ym, an

entirely new set of grain boundaries was exposed. Therefore

each sectioned surface can be considered as an independent

Sample.

Let N be the total number of sectioned surfaces examined.

The measured a .'s are ranked from the shallowest penetration

mi

depth to the deepest penetration. M^ is the ranking number

for ami'

For a large number of sectioned surfaces, the probability of

finding an oxide depth less than ami is P(ami) = P(a)i = Mi/(l +N)

The probability of having a depth equal to or deeper than a . is

Weibull assumes the distribution function

M m ^ 1 r Umi " au

)b

[1 - P(a).J =exp [ J 1 (1) 1 ao

and

ln ln

 [1 - P(a);LJ ' b ln (ami

In a Weibull plot of In In [1/{1 - PCa)^] versus

In (a . - a ), Equation (2) gives a straight line. au is the

horizontal shift for each of the data points, so that all the

points will be on a. straight line in the plot. The term a

10
sets the scale of the plot. The term b is the slope of the

line and is called shape parameter or Weibull modulus, a and

a are location parameter and scale parameter, respectively.

Figure 8 shows the Weibull plot for 480 data points at

800°C for 500 hours. Before the shift, the data points follow

a curve A in the right hand side of the figure. After the

shift of a (40 urn) the straight line B correlates well with

the data points. The vertical spread of the data points,

especially at the lower end of the line, is caused by a large

number of measurements having the same depth. The difference

in depths between neighboring points could not be resolved

easily with the microscope used. The three parameters were

observed as follows: b = 2.0, a = 40 urn and a = 740 \im.

The statistical distribution of a . at each of the test

conditions can be analyzed according to the Weibull plot

(Figure 8). Twelve plots are needed to analyze the combinations

of three oxidation temperatures and four exposure times. The

amount of analysis could become overwhelmingly large if wide

ranges of temperatures and exposure times are needed, and the

applications of such results could be complicated. Therefore,

it is desirable to develop a simplified data correlation method.

The oxide deptc h a . has been measured from 100 to 1000

 mi

hours at 600, 800 and 1000°C. It is assumed that the relation

between a . . t and T has the form

mi

ami

 = ai ^ exp (

- & (3)

11
The regression analysis of the data gives the following empirical relation.

ami(cm) = 1.34 * 10"3t°'25 exp (- 4.26/RT) (4)

t in seconds, the activation energy in kcal/mol, and T in °K.

The coefficient of correlation is 0.96. The plot of

0 25

In (a-mi/t ' ) versus (1/T) is shown in Figure 9. The slope

of the line is (4.26/R).

Grain boundary diffusion can be considered as a channeled

one-dimensional flow, with a constant concentration, C at the

interface between the surface oxide layer and the substrate (

see Figure 10). The oxygen concentration in the grain boundary

is (13)

C(x,t) - CQ [1 - erf ( - $ - )] (5)

where erf is the error function and D . is the grain boundary diffusion coefficient. The oxygen in the grain boundary will form

bulk oxide if the concentration reaches a -certain critical value

C . According to Equation (5),

X

C

- C0 [1 - erf(2 ± ) (6)

is the depth of the oxide penetration, where C = GC<

The quantity x /2 /D ,t is a constant and XG must be

proportional to /D~Tt. This is different from Eq. (4)

12
The orientation of the grain boundary oxide is determined by

the orientation of the substrate grain into which the oxide is

growing. The orientations of the two halves of the grain boundary

oxide growing into two "neighboring" substrate grains 'are different,

and an oxide grain boundary exists between these two halves. The

diffusion of oxygen will pass through the oxide grain boundary. The

oxygen in the grain boundary oxide will diffuse into the substrate

by bulk diffusion.

Equations (4) and (5) are true only for a one-dimensional

diffusion flow. The grain boundary oxidation is further complicated by the bulk diffusion of oxygen outward from the grain

boundary oxide into the substrate.

Figure 10 shows an element of the grain boundary, Ax x Ay.

The diffusion coefficient D ,, is much larger than the bulk

diffusion coefficient D,, often several orders of magnitude

larger. The large aspect ratio, a/b, in Figure 5, indicates that

the flux along the grain boundary is much larger than the flux

into the substrate. Whipple (19) has taken the bulk diffusion

into consideration in his analysis of grain boundary diffusion.

According to his study, the oxide penetration depth is affected

by the bulk diffusion. However, the extent of the effect has to

be assessed quantitatively for TAZ-8A. In addition, the chemical

process of oxidation and its effect on the gradient of chemical

potential have yet to be studied. For the moment, Equation (4)

can be treated as an empirical relation.

13
The empirical relation (4) correlates the data reasonably

well. However, sizable scatter is shown in Figure 9.

It is well known that grain boundary diffusion is a function

of the relative orientations of the two neighboring grains.

Couling and Smoluchowski (16) measured grain boundary diffusion

penetration of radioactive silver in copper as a function of

the angle of (100) tilt boundaries. Their results are shown in

Figure 11. Turnbul and Hoffman (17) have shown that the activation energy of grain boundary diffusion is a function of the

angle of misorientation, 8. Therefore, the activation energy

Q in Equation (3) is not a constant but dependent on 9. Since

we are measuring the deepest oxide penetration, perhaps Q is

close to the minimum value. However, a statistical variation

of Q from one cross-section to another is expected.

The dependence of a . on t is closely related to the oxide

morphology. For example the dependence on t for the cone type

oxide will be different from the pancake type. The detailed

oxidation process, the oxide structures and the effect of bulk

diffusion will affect the dependence of a . on both t and T.

mi

Equation (4) is an empirical equation that correlates

the measured oxide penetration. The deviation from the empirical relation of each measurement can be lumped into the term

°i

a. = amit-n exp () (7)

14
At any temperature T and exposure time t, with each of the measured a , the value of a. can be calculated with the empirically

determined constants n and Q. The Weibull plot of all of the 144

values of c^ of 12 data sets of 12 data points each is shown in

Figure 12. The values of b, a , a are 1.85, 0.53 x 10 and

— 6

4.8 x 10~ respectively.lt is interesting to note that the b-value

of 1.85 for a. for the 144 data points is close to the average

b-value of 1.8 of the 12 data sets for a .. a . is linearly promi mi

 J *

portional to a., and the statistical distribution of a. is directly related to-the statistical distribution of a ..

mi

One may consider the maximum value of a along the periphery of a sectioned surface "i" as the ai of an area of irDd

of the test coupon. D 'is the coupon diameter and d is the

grain size. Another sectioned surface at a distance one grain

size away can be considered a section containing an entirely

different set of grain boundaries and it is an independent

sample.

15
Let P(a) be the probability of finding the maximum a-value

less than a. on a unit surface area of a test coupon. P(a) is

related to the probability of a sectioned surface P(a).

[1 - P(a)t] = [1 - PCcc)] (8a)

Pc is the probability of finding an a-value less than a. on a

O J.

surface area S.

and

- P(a)s] = [1 - P(a)]S (8b)

[1 - P(a)s] = [1 - P(a).]S/1TDd

ln[l - P(a)g] = (S/rrDd) ln[l - P(a)i] (9)

We define R,

R = -(S/irDd) ln[l - PCa)^ (10)

Following Equation (1), we have

[1 - P(a)i] = exp[-(ai - au)

b

/ao] (11)

R =

P(o)g = 1 - exp(-R) (12)

The value of P(a)s thus obtained can be taken as the value

of P(a)s for the oxide depth a = atnexp(-Q/RT) . The P(a)g thus

determined is not the same as the probability of finding an oxide

depth less than "a" among a set of measurements determined at a

given oxidation temperature and exposure time. It incorporates

the additional "uncertainty" of T and t.

16
Zheng and Liu have studied the fatigue lives of precracked

cylindrical specimens. Both the smooth and precracked cylindrical

specimens are shown in Figure 13. The specimens were precracked

by a jeweler's saw (20).

The measured fatigue reversals, 2Np, are shown in Figure 14

2

function of deformation work density, which is (AJ /8E +

e ). AO is the applied fatigue stress range; E, the Young's

modulus; aYC> the cyclic yield strength; Aep, the applied

cyclic plastic strain range. The fatigue lives were measured

in the low cycle fatigue regime. The straight lines for the

smooth and the precracked specimens are parallel in the log-log

plot. Thus the ratio between the fatigue lives of the precracked

specimens, Np., and the fatigue lives of the smooth specimens, N_

is a constant for a given precrack size, and in general, it is a

function of the quantity (a /a.)

NN •— a ' Fo

 ai

a. is the precrack size, and a is a fictitious crack size for

the "smooth" specimens. The function F can be determined empirically. Once the function is determined, the fatigue life of the

preoxidized specimens can be predicted.

If one considers the grain boundary oxide as a precrack of

size a. , a. is a random variable and is also a function of the

time, t, and temperature, T, of oxidation. The remaining fatigue

life Np. of a pre-oxidized specimen must also be a function of t

and T, and N-. must also be a random variable.

r l

17
A mechanistic model to predict the fatigue life of a preoxidized specimen is outlined above. With additional information

on the effects of cyclic stress and cyclic strain on oxide crack

formation and a mechanistic model of fatigue crack growth at high

temperatures, the fatigue life at elevated temperatures can be

predicted quantitatively. A mechanistic model of grain boundary

oxidation accelerated fatigue crack growth has been proposed by

Liu and Oshida (21).

18
SUMMARY AND CONCLUSIONS

The effects of pre-oxidation on subsequent fatigue life was

studied mechanistically. Nickel-base superalloy (TAZ-8A) was

subjected to high temperature oxidation (at 600 to 1000°C for

10 to 1000 hours) in air. Surface oxidation kinetics and chemical

and structural studies on surface oxides were carried out. In the

high temperature regime (i.e., 800 to 1000°C), outward diffusion

of chrominum ions through the surface oxide layer appears to

control the oxidation kinetics. In the low temperature regime

(i.e., 600 to 800°C), nickel ion diffusion through the oxide

layer seems to govern the parabolic oxidation rate.

Through the three-dimensional morphological observations,

it is found that there exists two different types of grain

boundary oxides; cone type and pancake type. The pancake type

oxide will be more damaging because of its size and shape.

The grain boundary oxide penetration depth, a , is related

to oxidation temperature, T, and exposure time ,t, i.e.,

a = cttnexp(-Q/RT), with n - 0.25 and Q = 4.26 kcal/mol. m

A sufficiently large grain boundary oxide will be cracked

by an applied stress. The largest and deepest of the oxide cracks

is the most damaging to fatigue life. The statistical analysis

using the Weibull distribution function for the grain boundary

oxide penetration depth gives the probability of finding an

oxide crack of a certain size on a surface area exposed to

an oxidation temperature, T, for an exposure time, t. The size

of the oxide crack can be related to the remaining fatigue life.

19
REFERENCES

(1) "Manual on Low Cycle Fatigue Testing" ASTM STP No.465 (1969)

(2) "Symposium on Fatigue at Elevated Temperature" ASTM STP

No.520 (1972)

(3) Proc. International Conf. on Creep and Fatigue in Elevated

Temperature Applications (1973)

(4) ASME-MPC-3 Symposium on Creep-Fatigue Interaction (1976)

(5) R.M.Pelloux and N.S.Stoloff ed.;"Creep-Fatigue-Environment

Interactions" AIME (1980)

(6) L.F.Coffin, Jr.; Proc. of International Conf. on Fatigue:

Chemistry, Mechanics and Microstructure, pp.590/680 (1972)

(7) L.F.Coffin, Jr.; Met. Trans., 3, pp.1778/1788 (1972)

(8) S.D.Antolovich, P.Domas and J.L.Strudel; Met. Trans., 10A,

pp.1859/1868 (1979)

(9) C.J.McMahon and L.F.Coffin, Jr.; Met. Trans., I, pp.3443/3450

(1970)

(10) Y.Shida, G.C.Wood, F.H.Stott, D.P.Whittle and B.D.Bastow;

Corr. Sci., 2_1(8), pp.581/597 (1981)

(11) F.H.Stott, G.C.Wood, Y.Shida, D.P.Whittle and B.D.Bastow;

Corr. Sci.; 2_1(8), pp.599/624 (1981)

(12) R.Hales and A.C.Hill; Corr. Sci.; 12., pp.843/853 (1985)

(13) G.R.Rundell; Metal Progress, 127(4), pp.39/54 (1985)

(14) P Kofstad; "Nonstoichiometry, Diffusion, and Electrical Conductivity in Binary Metal Oxides" Wiley-Interscience, NY

pp.101/362 (1972)

20
(15) Y.Oshida and H.W.Liu: "Oxidation and Low Cycle Fatigue Life

Prediction", NASA Conf. Pub. 2339 on Turbine Engine Hot Section

Technology, pp.321/329 (Oct. 1984)

(16) L.Couling and R.Smoluchowski ; J. Appl .Phys. , 25_, p.1538 (1954)

(17) D.Turnbull and R.Hoffman; Acta Met., 2, p.419 (1954)

(18) P.G.Shewmon; "Diffusion in Solids" McGraw-Hill Book Com.,

NY (1963)

(19) R.T.P.Whipple; Phil.Mag. , 4J5, p.1225 (1954)

(20) M.Zheng and H.W.Liu; Technical Report NASA Grant No. NAG3-

348, "Crack Tip Field and Fatigue Crack Growth in General

Yielding and Low Cycle Fatigue", (1984)

(21) H.W.Liu and Y.Oshida; "Grain Boundary Oxidation and Fatigue

Crack Growth at Elevated Temperatures" to be published

(1985)

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25
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Figure 4 Line profiles of Ni, Cr, and Al on coupons oxidized at

three different oxidation conditions

26
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Figure 5 Cross-sectional micrograph, showing a deep penetration

by a grain boundary oxidation

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36
National Aeronautics and

Space Administration

Lewis Research Center

Cleveland. Ohio 44135

Official Business

Penalty for Private Use $300

SECOND CLASS MAIL

ADDRESS CORRECTION REQUESTED

HUH

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Space Administration

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