Low Cycle Fatigue Behavior of Rene 80

Low Cycle Fatigue Behavior of Ren6 80 at Elevated Temperature

STEPHEN D. ANTOLOVICH, S. LIU, AND R. BAUR

Low cycle fatigue of Ren6 80 was studied at 871 and 982 ~ It was found that when the data were represented on the basis of plastic strain, the life increased with decreasing frequency and imposition of a 90 s hold at maximum strain. Transmission electron microscopy studies showed that the 7' coarsened and an interfacial array of edge dislocations developed. The density of dislocations in the matrix was very low. Light optical microscopy revealed that cracks generally inititiated at oxide spikes in surface connected grain boundaries. A crack initiation criterion based on the maximum stress and oxide depth at the time of crack initiation was found to represent the data very well. Based on that representation, an expression for the initiation fatigue life was developed. That expression includes temperature, frequency and cyclic stress strain parameters as variables.

I. INTRODUCTION

LOW cycle fatigue (LCF) at elevated temperature is an important consideration in the design of turbine com- ponents such as disks and turbine blades. Various methodologies for predicting the life of such compo- nents have been proposed 1~ and one of these methodologies ~ has been the subject of an international conference? In addition, an exhaustive review of available data and life prediction schemes has recently been made. 6 It is striking that various approaches seem to work well for some alloy systems but not for others and that a universally applicable fatigue law has not been discovered. This is probably due to the fact that in different classes of materials, different forms of damage can occur. For example, grain boundary sliding and cavitation might occur relatively easily in those systems where there is no pinning of the boundaries by precipitate or carbide particles. On the other hand, some systems may be very prone to environmental attack. Still others will fail due to the accumulation of deformation debris. Not only will the failure mode depend on the system being considered, but also on the temperature, strain range and strain rate. In addition, interactions between the modes of damage accumulation can occur. Metallurgical factors have been discussed in review papers. 7,8

In this study, the Ni base superalloy Ren6 80 was investigated primarily at 871 ~ (1600 ~ and at 982 ~ (1800 ~ Ren6 80 is currently used as a turbine blade material in this temperature regime.

STEPHEN D. ANTOLOVICH is Professor of Materials Science, University of Cincinnati, Cincinnati OH 45221. S. LIU, formerly Research Associate at the University of Cincinnati, is presently Research Metallurgist, Special Metals, New Hartford, NY 13413. R. BAUR is Senior Engineer, General Electric Co., Aircraft Engine Group, Evendale OH 45215.

Manuscript submitted February 15, 1980.

METALLURGICAL TRANSACTIONS A

II. EXPERIMENTAL PROCEDURE

A. Materials and Heat Treatment

Test specimens of Ren6 80 were cast to shape. The composition (wt pct) is 3.0A1, 5.0Ti, 14.0Cr, 3.9W, 4.0Mo, 9.8Co, 0.17C, Ni bal. The heat treatment has been described elsewhere 9 and results in a mixture of small spherical 7' particles and large V' cubes. After heat treatment, the specimens were machined by low stress grinding.

B. Mechanical Testing

The LCF behavior was investigated using longitu- dinal specimens at 871 and 982 ~ Testing was done using an MTS in the strain control mode. The cycle character was generally zero-tension-zero (R, = %,,/%,x ~ 0.05) and the test frequency was adjusted to maintain loading and unloading strain rates of either 0.5 pct/min or 50 pct/min. Some testing was done with a 90 s hold time at the maximum strain.

Temperature was maintained at either 871 or 982 ~ using an RF generator and the maximum variation along the gage length was 3 ~ C.

C. Metallography

1. Optical. In order to examine microstructural features, as well as the relationship between the crack path and the microstructure, tested specimens were sectioned on a plane containing the stress axis and crack initiation site.

2. Scanning Electron Microscopy (SEM). The fracture surfaces were characterized by SEM. SEM was also used to investigate microstructural features at high magnifications.

3. Transmission Electron Microscopy (TEM). Thin wafers about 0.5 mm thick were cut perpendicular to the specimen axis about 2 to 3 mm from the fracture surface. The thickness of the wafer was reduced to 0.13 mm by mechanical polishing. TEM foils were then

ISSN 0360-2133/81/0311-0473500.75/0 9 1981 AMERICAN SOCIETY FOR METALS AND VOLUME 12A, MARCH 1981--473

THE METALLURGICAL SOCIETY OF AIME

0 10

o~

LU

Z 101 < r r

Z m < cc I-,- -2 O9 10 O m

I-- O3 < J

&. 10 ~ 2

10

871%

KEY T IME~D STRAIN HOLD RATE (SEC) %/MIN 0 90

5 9

50 ~} 9 J , L , l ,~L i , i i i , L i ]

10 a 10"

CYCLES TO

, t , , , , , , 5

10

INITIATION N i Fig. 1--Coffin-Manson behavior of Ren6 80 at 871 ~ Cycles to initiation.

uJ

z

z

I.-.- u? 0 I....-

_A

9 " \

\

KEY [] HOLD TIME

(SEC)

0 90

[] 9

STRAIN

RATE

%/MIN O5

1~32 so i

0 a

982b

o c

l ~ , I 1,11103 , I I I , i~l i i 1 0*

CYCLES TO IN IT IAT ION, N,

Fig. 2--Coff in-Manson behavior of Ren~ 80 at 982 ~ Cycles to initiation.

lO

UJ O Z r r

Z

r r p- 03

O F- o9 / 13-

ld

[] o 9 9 871~

KEY

STRAIN RATE HOLD TIME. SEC

%/MIN 0 [ 90

0.5 0 1 II ~o C] 9

[ I I I ] l i l { 1 t

ld

CYCLES TO

th

J l l , , i l i I I l l I I I I

10 1@ 1(~

FAILURE N F

Fig. 3--Coff in-Manson behavior of Rene 80 at 871 ~ Cycles to failure.

(5 Z

_z r r

CO

o3 < g.

10 ~

STRAIN

RATE 7~ /MIN

5

50 10 2

i

10'

. . . . - ' ' , . . . . , - . , i

BoA kko 982~;

E~

KEY ~ HOLD rIME n [] 0

(SEC) X 0 9O

0 9

i , , , , , , I . . . . . . . . I

ld 1r

I I p I 9 i me

1@

CYCLES TO FAILURE, Nf

Fig. 4--Coff in-Manson behavior of Ren~ 80 at 982 ~ Cycles to failure.

made by electropolishing using a Fischione twin-jet unit. They were subsequently examined in a JEM 200A transmission electron microscope operated at 200 Kv. An in situ stereographic projection of all foils was made to facilitate the excitation of desired reflections.

D. Oxidation Studies

Since environmental interactions have been previously demonstrated to be of prime importance for this alloy, 9 the oxidation kinetics were studied both by optical microscopy and by measuring the weight gained by a polished sample as it oxidized in air. The optical microscopy entailed sectioning a specimen that had been oxidized between 1 and 500 h and measuring the longest oxide spikes that were observed along grain boundaries. The other procedure was carried out using a microbalance connected to a strip chart recorder to provide a continuous record of weight gain vs time for temperatures in the range of interest.

III. RESULTS AND DISCUSSION

A. Mechanical Properties

1. Tensile Properties. Tensile test results have been reported previously. 9 In general the strength decreases and ductility increases with increasing test temperature. Strain rate effects are small at 760 ~ but become pronounced at higher temperatures. Thus to investigate the effects of thermally activated deformation (i.e., creep) most testing was done above 760 ~

2. Low Cycle Fatigue (LCF). The LCF test results are listed in Tables I and II and Coffin-Manson plots are shown in Figs. 1 to 4. Low rate and hold time tests exhibit a marked improvement in LCF compared to high rate continuous cycling. The reasons for this behavior are discussed in a subsequent section.

Cumulative glide plots are shown in Figs. 5 and 6. In general, the material shows initial rapid softening followed by a regime in which the stress falls off less rapidly. This softening can result from fatiguing of a

474- -VOLUME 12A, MARCH 1981 METALLURGICAL TRANSACTIONS A

cold worked structure, 1~ from shearing of precipitates I1 or from structural instabilities) 2 The reasons for the observed softening behavior will be discussed later in light of microstructural observations. Finally near the end of the test when cracks are present between the extensometer probes the stress again drops off very

rapidly. Crack initiation was defined by linearly ex- trapolating the midlife and end of life regimes of the cumulative glide plots and taking the point of inter- section as N,.

Cyclic stress strain curves are shown in Figs. 7 and 8. Figs. 7 and 8 show a reduction of stress amplitude with

Table I. LCF Oata for Ren6 80 at 871 ~

Plastic Total Strain Strain Stress Maximum Young's

Strain Range Range Range Stress Modulus Cycles to Cycles to Specimen Temperature Rate Aep Act Ao %~x EX10 -3 Initiation Failure Number (~ (~ pct/min) pet pet (MPa) (MPa) (MPa) Nj Nf Comments

16LV12 871 50 0.013 0.487 771 586 158 9780 11760 Failed 16LV3 871 45 0.020 0.50 755 731 163 - - 2980 Failed outside extensometer 16LV15 871 0.50 0.048 0.474 650 340 139 1945 - - Removed at N, 16LV9 871 50 0.048 0.781 1095 677 140 1244 1342 Failed 16V11 871 0.50 0.051 0.473 661 344 154 2682 - - Failed, temperature loss causing

failure, removed at 2742 16LVI4 871 50 0.051 0.736 1138 658 141 1417 - - Removed at N t 16LV13 871 50 0.051 0.800 1092 732 144 936 1226 Failed 16LV7 871 29 0.062 0.67 954 732 158 - - 710 Failed 16LV21 871 80 0.12 0.85 1172 627 154 710 980 Failed 16LV22 871 85 0.135 0.91 1241 662 161 470 - - Removed at 591 16LV23 871 91 0.14 0.97 1241 669 148 330 496 Failed 16LV17 871 50 0.23 1.25 1472 775 147 230 - - Removed at N i 16LVi9 871 0.50 0.25 0.815 827 416 147 455 - - Removed at N s 16LVI8 871 0.50 0.25 0.82 781 378 158 421 732 Stopped after 75 pet drop in load

from mid-life 16LVI6 871 50 0.25 1.06 1360 710 160 202 312 Failed 16LV24 871 0.50 0.255 0.82 965 428 172 485 541 Failed

loading 50

unloading 16HLV2 871 50 0.05 0.465 776 239 168 7440 - - Tensile hold. Removed at N~ 16HLV1 871 50 0.05 0.490 792 246 163 7008 12867 Tensile hold. Failed 16HLV6 871 0.50 0.25 0.73 607 258 136 580 - - Tensile hold. Removed at N t 16HLV5 871 0.50 0.25 0.70 709 317 134 662 929 Tensile hold. Failed 16HLV3 871 50 0.25 0.86 1015 309 151 631 1584 Tensile hold. Stopped after 75

pet drop in load from mid-life 16HLV4 871 50 0.25 0.885 1085 377 158 665 - - Tensile hold. Removed at Nt

Table II. LCF Data lot Ren6 at 982 *C

Plastic Total Strain Strain Stress Maximum Young's

Strain Range Range Range Stress Modulus Cycles to Cycles to Specimen Temperature Rate A% Act Ao Oraax EXI0 -3 Initiation Failure Number (~ (t pct/min) Pct Pet (MPa) (MPa) (MPa) N~ Nf Comments

18LV1 982 0.42 0.046 0.282 355 186 138 9720 - - Removed at Ni 18LV2 982 0.42 0.040 0.300 401 214 142 4772 5445 Failed 18LV3 982 53.1 0.040 0.468 622 341 144 2119 2511 Failed 18LV4 982 54.0 0.040 0.476 639 409 142 1525 - - Removed at 1810 OX 18LV8 982 54.0 0.136 0.624 731 374 146 752 - - Removed at 848 OX 18LV7 982 54.0 0.140 0.624 706 374 142 753 931 Failed 18LV5 982 0.46 0.184 0.532 485 252 143 523 877 Failed 18LV6 982 0.46 0.192 0.552 476 245 129 473 - - Removed at 503 18HLV1 982 0.46 0.084 0.332 381 168 137 1437 1620 Failed OX 18HLV9 982 0.47 0.090 0.342 399 177 139 1499 - - Removed at N, 18HLV4 982 54.0 0.104 0.476 611 182 139 835 1457 Failed OX 18HLV10 982 49.4 0.110 0.436 585 189 138 984 - - Removed at 984 18HLV5 982 0.46 0.196 0.532 495 226 145 526 573 Failed 18HLV6 982 0.46 0.212 0.532 496 234 139 331 - - Removed at N~ 18HLV11 982 54.5 0.222 0.628 654 241 143 405 - - Removed atN t 18HLV8 982 54.0 0.230 0.624 705 258 138 378 483 Failed

NOTE: OX indicates failure outside extensometer probes.

METALLURGICAL TRANSACTIONS A VOLUME 12A, MARCH 1981--475

increasing temperature and decreasing strain rate. This observation will also be discussed later in terms of structural changes.

3. Effect of HoM Time. Just as the life increased with decreasing frequency for the continuous cycling tests, the effect of introducing a 90 s hold was to further increase the life and this increase was all the more pronounced as the plastic strain decreased. This is a further manifestation of the fact that a negative creep/ fatigue interaction is not operative for this alloy.

The tensile hold tests showed longer lives than was observed for compressive holds. The tensile holds develop lower maximum stress than compressive holds and if maximum stress is an important factor in determining the life, such behavior would be expected.

CO W qs F- O3

500

250

o ~4~ 9 ~ t I t

0 400 800

CYCLE

Fig. 5---Cumulative glide behavior of Rene 80 at 871 ~ for continous cycling at 0.50 pct/min.

CO 03 W OZ I-- CO

250

0

i

o

I

= ; J

9 STRESS RANGE

9 MEAN STRESS

982 ~ C

I 4OO 8O0

CYCLE

Fig. 6---Cumulative glide behavior of Ren6 80 at 982 ~ for continuous cycling at 0.50 pct/min.

It has been suggested in a previous paper 9 that a combination of response stress and environmental damage interact to define crack initiation. This hypothesis is investigated further in subsequent sections.

B. Metallography

1. Optical and SEM. The surface region of an LCF specimen is shown in Fig. 9 which is typical of what was observed for all test conditions: oxide spikes in the grain boundaries and a uniform wall of oxidation elsewhere. Such behavior was observed in a previous study 9 and the oxide spikes were usually the source of crack formation. However, for high rate hold time specimens creep voiding was observed in addition to surface cracking and this can be seen in Fig 10. Even though this new failure mode began to appear under these conditions, the general trends in life were not changed. It would be reasonable to assume that if the test conditions were made more extreme in terms of hold times and strain rates that the creep voiding could become a dominant failure mode but this was not the case in the studies reported here.

Due to oxidation of the fracture surface features, detailed SEM fractography was not carried out. In

c~

uJ

::3 F--

,<

09 O3 !JJ

I-- CO

103

102

, i i , i J i J i

15 3 1~ 2

871~

Y%--S~ KEY

STRAIN RATE HOLD T IME

(SEC)

%IMIN O 90

5 O 9

50 D I I I

i , i l l , , i J J J i g i l l

161 10 (

STRAIN AMPLITUDE % Fig. 7--Cyclic stress/strain behavior of Rene 80 at 871 ~

. i t3

I--

_J 13_

03

or" I-- O3

10 a

102

982~

10 -3

, , i , i , I l i r i l l l l l

10 -2

STRAIN AMPLITUDE %

KEY STRAIN RATE HOLD T IME

ISEC}

%/MIN O 90

5 0 I

5O

I l i i g i l l

10 -1 100

Fig. 8--Cyclic stress/strain behavior of Ren6 80 at 982 ~

476--VOLUME 12A, MARCH 1981 METALLURGICAL TRANSACTIONS A

general, cracks initiated intergranularly with some evi- dence of transgranular propagation on the fracture surface. The final overload region had a characteristic dendritic appearance typical of these cast alloys.

2. TEM. A micrograph of the starting structure is shown in Fig. i 1. The structural changes as a function of test conditions were studied in detail, the results of which are summarized in Figs. 12 and 13. For all test conditions there was a tendency for the large 7' to coarsen and for the small 7' to disappear. For Ren6 77, in which the 7' is structurally similar to that in Ren6 80, it has been shown that coarsening of the 7' morphology takes place early in the fatigue process and requires plastic deformation to occur. ~3 Such behavior is likely to be the case for Ren6 80. A comparison of Figs. 12 and 13 shows that the morphological changes in V' are most pronounced for high temperatures, loading rates and imposed hold times.

At the temperatures and strain rates employed in this study, interfacial networks of near-edge dislocations are formed on the 7', similar to those observed elsewhere. 9,~4 The tendency for this to occur is again most developed for high temperatures, low loading rates and imposed hold times, Fig. 13(d). These observations are all

consistent with the observed cumulative glide and cyclic stress strain behavior. The continual drop off in stress with increasing cycles seen in Figs. 5 and 6 can be explained in terms of progressive coarsening. Similarly coarsening is more pronounced at high temperatures and low rates resulting in the cyclic stress/strain behavior seen in Figs. 7 and 8.

It is difficult to imagine that the structural changes and development of the interracial dislocations seen in Figs. 12 and 13 can be damaging. In fact, as has been previously pointed out, 9 the coarsened particles are associated with increased ducility and the interracial dislocations reduce the matrix/precipitate strain energy.

In view of the above observations, failure does not occur as a result of reaching a critical dislocation density and arrangement. Instead, the most damaging

Fig. 9---Oxide spike initiating surface crack. Specimen 16LV18, plastic strain 0.25 pct, life 732 cycles.

Fig. 10~Grain boundary voiding for high ramp rate hold time specimen. Specimen 16HLV1, plastic strain 0.05 pct, life 12867 cycles.

Fig. 11--Initial structure of Ren6 80 showing: (a) grain structure by light optical microscopy, (b) precipitate structure by dark field TEM, g = (100), (c) precipitate/dislocation structure by bright field TEM, g = (200).


event appears to be the development of oxide spikes along grain boundaries and interdendritic regions. Im- plications of this microstructural observation are used in the next section to suggest a possible model for high temperature LCF crack initiation.

C. Model Development

As a first order approximation, it will be assumed that as an oxide spike develops, at some point, under the action of the cyclic stress it loses coherency with the

base metal and becomes a well-defined crack.* Intui-

* Another possibility is penetration of oxygen or nitrogen along the boundaries which causes an embrittled region that eventually cracks. Such a mechanism is also qualitatively consistent with the following discussion.

tively, the point of cracking depends on the size of the oxide spike and the magnitude of the stress, which in turn depends on the underlying microstructure of the metal. The formation of an oxide spike is a thermally activated process and in many systems follows well-

(a) (b)

(~ rd) Fig. 12--Precipitate/dislocation microstructure of Ren~ 80 continuously cycled at 871 ~ (a) Specimen 16LVI2 (low strain, high rate) g

(200), (b) Specimen 16LVI6 (high strain, high rate) g ~ (200), (c) Specimen 16LV11 (low strain, low rate) g = (200), (d) Specimen 16LV18 (high strain, low rate) g = (111). See Table I for test conditions.

478--VOLUME 12A, MARCH 1981 METALLURGICAL TRANSACTIONS A

known kinetics which may be expressed mathematical ly as;

z, = ~ ~ [1] where l i = spike depth at initiation, D = diffusion constant, t, = time to crack initiation, and ct = geo- metric constant.

We have verified that Eq. [1] is applicable to grain boundary oxidation of Ren6 80 by the optical microscopy experiments cited previously. The results of those

observations can be represented by the following equation:

t, = ,~ (Dt ) '~ [2] where m = 0.49 at 982 ~ and 0.46 at 760 ~ The data were also used to determine an apparent activation energy for grain boundary oxidation of 9.6 kcal/mol. Further verification of parabolic oxidation kinetics was obtained from the microbalance experiments. The re- suits of these experiments showed that the weight gain

(a) (b)

~O fd)

Fig. 13--Precipitate/dislocation microstructure of Ren~ 80 subject to a 90 s hold at maximum strain during LCF at 982 ~ (a) Specimen 18HLV4 (low strain, high rate) g = (200), (b) Specimen 18HLV8 (high strain, high rate) g = (200), (c) Specimen 18HLV1 (low strain, low rate) g

(200), (d) Specimen 18HLV5 (high strain, low rate) g = (200).


per unit area could be described by the following equation:

AW __ = Ct]/2 A [31

where A W = weight gain, A = area, C constant, and t = time.

The constant depends on temperature and was measured as 7.6 X 10 -4 (gm/cm2).s 1/2 at 982 ~ and 4.9 X 10 -4 (gm/cm2)-s in at 928 ~

It has been shown elsewhere 15 that cyclic deformation significantly increases the oxidation rate in Ren6 77, a microstructurally similar alloy. The effect of cyclic deformation on the form of the oxidation equation is not known for these systems. Lacking this information it seems reasonable to assume that the form of the oxidation process can be represented by Eq. [1]. Using these assumptions, relative oxide spike depths were calculated for purposes of developing a correlation. To do this the spike depth was taken as unity for the test of shortest duration. The depth for any other test was simply found by taking the ratio of the test times as implied by Eq. [1]:

l , = ( t / to ) 1/2 [4 ]

where t = test time and t o = test time for reference specimen.

Next, a crack initiation criterion of the form shown below was adopted as an operating hypothesis:

o?.x . /e = Co [51

where o m"x = maximum stress at initiation, Co = material constant, and p = exponent.

Values of o, ~" were measured and plotted as a function of the computed l, values for both 871 and 982 ~ The results are shown in Figs. 14 and 15. The differences in constant reflect the fact that for each temperature the depth was taken as unity for the shortest test time. The results, to a high degree of confidence, reflect the applicability of Eq. [5] as an initiation criterion. It has been shown elsewhere 16J7 that Eq. [5] is also applicable to other materials and test conditions.

To develop an expression for the fatigue life in terms

871 ~ %Lo I SR: : : [ HO(L:EGTIME [ STRAIN RATE (SECI TIME

3~ CONTINUOUS CYCLING, 1%/MIN ~ lO ~[~ Omax = 758 I, -23 cf:.95)

Oma==757 91~ -29 If: 94)

10~ . . . . . . iI i i I i . . . . 10 0 101 10 2 1(

RELAT IVE OXIDE DEPTH AT IN IT IAT ION

Fig. 14- -Maximum stress at initiation as a function of oxide depth at initiation for Renr 80 tested at 871 ~ The equations of each line along with the coefficients of determination are given.

of more customary variables, Eq. [1] may be substituted into Eq. [5] and appropriate expressions may also be substituted for the maximum stress and time to initiation:

[61

t, = ( l / . + th)N, [7]

where d = mean stress, K -- fatigue strength coefficient, A% = plastic strain range, n' = cyclic strain hardening exponent, v = continuous cycle frequency, t h = hold time, and N, = cycles to initiation.

Finally, from Figs. 14 and 15 the exponent p in Eq. [6] may be taken as 1/4 with the result:

+ K Da2( 1 th) N , C [8]

Temperature effects are accounted for explicitly in the diffusion constant:

D = Doexp - Q /RT [9]

where the symbols have their usual meamng. Equation [8] represents an explicit evaluation of the correlations developed in Figs. 14 and 15.

The increases in life that were observed for low frequencies and hold times can be explained in terms of structural coarsening. The structure is unstable relative to plastic deformation and the stress drops off significantly. This necessitates a large value of l, in Eq. [5] and a correspondingly large number of cycles. Of course a prediction, in the true sense of that word, must include the ability to predict the stress response as a function of test times and temperatures. This means that the structural changes must also be predicted as a function of these variables. Obviously such studies must be undertaken if the ideas advanced here are to be put in more fundamental terms.

In all tests carried out for this study, the mean stress 8 was non-zero. However, in many instances stable systems are studied in fully reversed fatigue (# = 0). In this case, and taking a --~ 1 for convenience, Eq. [8] reduces to an expression of the form:

o ,10, N, = C2 exp~-Ac

It is noteworthy that for a given temperature and frequency Eq. [lO] has the same form as Coffin's frequency modified fatigue life with an explicit expression for the frequency and temperature dependence. Furthermore, the exponent fl in the Coffin-Manson equation is given by:

1 f l - 8n' [11]

It is interesting to compare the qualitative predictions of Eq. [10] with what is frequently observed for stable materials tested in fully reversed fatigue at high temperatures. For example, a decrease in life is predicted for low frequencies, long hold times and high temper-

480- -VOLUME 12A, MARCH 1981 METALLURGICAL TRANSACTIONS A

103 982 ~

D~ KEY STRAIN ] HOLD TIME

CON3 LNUOUS CYC l ING RATE (SEC)

O5

r ~ c tmax=385 9 I, -21 ( |= 9B5) %/MIN 0 90 ,.

ALL DATA

~:~ Omax:38291, -25 ~t : 93)

l , , t t l i , i I i I I I I I I [ I I I I I I l

10 ~ 101 102 103

RELATIVE OXIDE DEPTH AT INITIATION Fig. 15--Maximum stress at initiation as a function of oxide depth at initiation for Ren6 80 tested at 982 ~ The equations of each line along with the coefficients of determination are given.

atures and such behavior is, in fact, often observed, especially for the conditions enumerated above.

Finally, values of n' in the neighborhood of 0.15 to 0.20 are frequently observed and values of the Coffin- Manson exponent in the range 0.6 to 0.8 are predicted which again agrees with observations. At this point we would consider the model to be a promising initial approach to understanding high temperature LCF behavior in certain Ni base superalloys. Providing a more fundamental basis will require detailed studies of dy- namic oxidation and coarsening phenomena.

Based on the above discussion, it is clear one cannot conclude that accumulated deformation damage is the sole cause of LCF cracking merely because the data can be represented by a Coffin-Manson plot. Such a representation is also consistent with cracking by environmental damage. Identification of specific damage mechanisms would involve investigations of the temperature and frequency dependence of the fatigue life as well as detailed metallographic studies.

CONCLUSIONS

1) The fatigue life of Ren6 80 at both 871 and 982 ~ increases with decreased frequency and imposition of a 90 s hold time. Such behavior is at variance with the concept of a negative creep/fatigue interaction.

2) The ?' precipitate in Ren6 80 coarsened and developed an interracial array of edge dislocations with very few dislocations in the matrix. Such changes do not constitute damage in any fundamental sense.

3) Cracks initiated at oxide spikes in surface connected grain boundaries. Only in the case of hold time

specimens rapidly loaded to maximum strain was any internal void formation seen. Even in these cases there were also surface initiation sites.

4) The metallographic observations were used to develop an intuitive correlation between maximum stress and depth of oxidation at the time of crack initiation. All data were well-represented by this correlation.

5) The correlation demonstrated above was used to develop an expression for the crack initiation life. The expression has the form of the Coffin-Manson law for a given set of experimental conditions. In addition to cyclic stress/strain parameters frequency, hold time, and temperature are explicitly incorporated.

ACKNOWLEDGMENT

The authors wish to express their appreciation to the Air Force Office of Scientific Research for financial support of this project (AFOSR Grant No. 76-2952), and to Dr. A. Rosenstein, Program Manager, for his very useful assistance.

REFERENCES

1. S. S. Manson: ASTM STP 520, 1973, p. 744. 2. L. F. Coffin: Proc. Second lnternational Conf. on Mech. Behavior of

Materials, p. 866, ASM, Metals Park, OH, 1976. 3. S. Majumdar and P. S. Maiya: Proc. ASME-MPC Symposium on

Creep-Fatigue Interaction, MPC-3, p. 323, ASME, N.Y., 1976. 4. W. Ostergren: J. Test. Eval., 1976, vol. 4,p. 327. 5. A GARD Conference Proceedings, AGARD-CP-243, 1978. 6. L. F. Coffin, S. S. Manson, A. E. Carden, and L. K. Severud:

ORNL Report # 5073, Jan. 1977. 7. M. Gell and G. R. Leverant: ASTM STP 520, 1973, p. 37. 8. Stephen D. Antolovich: La Fatigue des materiaux et des struc-

tures, C. Bathias and J. P. Bailon, eds., p. 465, Maloine S. A., Paris, 1980.

9. Stephen D. Antolovich, P. Domas, and J. L. Strudel: Met. Trans A, 1979, vol. 10A, p. 1859.

10. C. E. Feltner and C. Laird: Acta. Metall., 1967, vol. 15, p. 1621. 11. R. E. Stoltz and A. G. Pineau: J. Mater. Sci. Eng., 1978, vol. 31, p.

275. 12. D. Fournier and A. Pineau: Met. TransA, 1977, vol. 8A, p. 1095. 13. E. Rosa: M. S. Thesis, University of Cincinnati, 1978. 14. A. Lasalmonie and J. L. Strudel: Philos. Mag., 1975, vol. 32, p.

937. 15. Stephen D. Antolovich, Eduardo Rosa, and Andr6 Pineau: Mater.

Sci. Eng., in press. 16. Stephen D. Antolovich and N. Jayaraman: Paper presented at the

27th Army Materials Research ConferEnce, Lake George, NY, July, 1980. Paper in press in the conference proceedings.

17. Stephen D. Antolovich, R. Baur, and S. Liu: Superalloys 1980, J. K. Tien, S. T. Wlodek, H. Morrow III, M. Gell, and G. E. Maurer, eds., p. 605, ASM, 1980.


Documents

Low Cycle Fatigue Behavior of Rene 80