NASA Technica!Memorandum 103095 _

" _ Fatigue Crack GrOwth in a Unidirectional

SCS-6/Ti- 15-3 Composite

Peter Kantzos and Jack Telesman

National Aeronautics and Space Administration

Lewis Research_ Center .............

Cleveland, Ohio

Louis Ghosn

-Sverdrup Technology, Inc. _ .....

NASA Lewis Research Center Growl r ......

Cleveland, Ohio

Prepared for the

Third Sympos_ on Composite Materials: Fatigue and Fracture

sponsored by the American Society for Testing Materials

Orlando, Florida, November 6-9, 1989

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(_!ASA-T_-103095) FATIGUE CRACK G_OWT_4 IN A

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FATIGUE CRACK GROWTH IN A UNIDIRECTIONAL SCS-6/Ti-15-3 COMPOSITE

Peter Kantzos and Jack Telesman

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135

Louis Ghosn

Sverdrup Technology, Inc.

NASA Lewis Research Center Group

Cleveland, Ohio 44135

Abstract

An investigation was conducted to characterize and model the

fatigue crack growth (FCG) behavior of a SCS_6/Ti-15-3 metal matrix

composite. Part of the study was conducted using a fatigue loading

stage mounted inside a scanning electron microscope (SEM). This

unique facility allowed high magnification viewing of the

composite fatigue processes and measurement of the near crack tip

displacements. The unidirectional composite was tested in the [0]8

(i.e longitudinal) and [90]8 (i.e. transverse) orientations. For

comparison purposes unreinforced matrix material produced by the

identical process as the reinforced material was also tested.

The results of the study reveal that the fatigue crack growth

behavior of the composite is a function of specimen geometry, fiber

orientation and the interaction of local stress fields with the

highly anisotropic composite. In the case of [0]8 oriented single

edge notch (SEN) specimens and [90]8 oriented compact tension (CT)

specimens, the crack growth was normal to the loading direction.

However, for the [0]_ CT specimens the crack grew mostly parallel

to the loading and the fiber direction.

The unusual fatigue behavior of the [0]8 CT specimens was

attributed to the specimen geometry and the associated high tensile
bending stresses perpendicular to the fiber direction. These

stresses resulted in preferential cracking in the weak interface

region perpendicular to the fiber direction. The interface region,

and in particular the carbon coating surrounding the fiber proved

to be the composites weakest link.

In the [0]8 SEN the crack growth was confined to the matrix

leaving behind unbroken fibers which bridged the cracked surfaces.

As the crack grew longer, more fibers bridged the crack resulting

in a progressive decrease in the crack growth rates and eventual

crack arrest. The actual near crack tip displacement measurements

were used in a proposed formulation for a bridging-corrected

effective crack driving force, AK.f:. This parameter was able to

account for most of the experienced bridging and correlated the

[0]8 SEN fatigue crack growth data reasonably well.

Key Words: Metal Matrix Composite, Fatigue Crack Growth, Crack

Bridging, Fatigue Mechanisms.

Introduction

The new generation of aerospace vehicles will require

materials capable of withstanding high temperatures while retaining

a high stiffness under relatively high loads. Continuous fiber,

metal matrix composites (MMC) are candidate materials for such

applications. One of these candidate materials is a Ti-15V-3Cr3AI-3Sn matrix reinforced by continuous SiC fibers. Before these

2
materials become widely utilized in aerospace applications, their

fatigue behavior have to be understood and reliable life prediction

methodology has to be developed.

Relatively little information is available concerning the

fatigue behavior of continuous fiber reinforced composites. The

objective of this study is to broaden this knowledge by

investigating in detail the fatigue behavior of the above mentioned

composite system. The emphasis of the study was twofold: i)

identification of the fatigue crack initiation and propagation

damage processes with special attention focused on identifying the

composite constituent most susceptible to fatigue damage; and 2)

identification of the appropriate crack driving force for this

composite. To achieve these goals, part of the study was conducted

using a fatigue loading stage mounted inside an SEM. This unique

facility, developed at NASA Lewis (i), allows real-time dynamic and

static viewing of the fatigue crack initiation and crack

propagation processes.

Material

The composite used in this study is a Ti-15V-3Cr-3AI-3Sn (Ti15-3) alloy matrix reinforced by 145 _an average diameter,

continuous SiC (SCS-6) fibers. All the composite specimens used in

this study were obtained from a single, eight ply, unidirectionally

reinforced panel manufactured by Textron Specialty Materials

Division. The panel size was 30x30x0.21 cm. Specimens were also

machined out of a 30x30xl cm panel of unreinforced matrix, using

the same Ti-15-3 matrix foil and same consolidation techniques
employed in composite production.

The matrix is a metastable beta phase alloy, chosen because

it can be cold rolled to very thin sheets. This enables a more cost

effective processing of the composite. Sample micrographs of the

reinforced composite are shown in Fig i. The SiC fiber is

surrounded by a complex multilayer structure (Fig ib) which

consists mainly of a multilayer carbon coating approximately 3 _m

thick and an approximately 0.5-i _m thick reaction zone consisting

mainly of brittle intermetallic phases (2). The combination of the

carbon layers and the reaction zone will henceforth be referred to

as the interface region.

Heat Treatment

Specimens were tested both in the as-received and heat treated

condition. The heat treatment consisted of 24 hours at 700 C in

vacuum. The purpose of the heat treatment was to precipitate out

the alpha phase. The heat treatment did not produce any additional

noticeable fiber/matrix interactions and the reaction zone remained

unchanged (3).

Mechanical Testing

Fatigue crack growth (FCG) testing was performed using single

edge notch (SEN) and compact tension (CT) specimens The SEN

specimens were tested in the [0]8 orientation only. The CT

composite specimens were tested in two orientations: i)

longitudinal, [0]8 , with fibers oriented parallel to the loading

direction; and 2) transverse, [90]8 , with fibers perpendicular to

the loading direction. Unreinforced matrix specimens were also

4
tested using CT specimens.

All machining was done using diamond tip tools. Prior to

testing, all specimens were polished to enhance crack detection.

The test matrix is shown in Table I. The SEN specimens were

tested in the fatigue loading stage mounted inside an SEM. The

testing was conducted at room temperature at a load ratio R of 0.1

with the maximum load being 2.24 KN. Test frequency was 5 Hz with

frequent holds at varied loads to observe the fatigue processes and

measure near crack tip displacements. An extensometer was mounted

in the crack mouth of the specimen to determine the changes in the

compliance. The vacuum inside the SEM was approximately ixl0 "6 torr.

Testing of the CT specimens was performed at room temperature and

ambient environment using a computer controlled, closed loop,

servo-hydraulic machine. The tests were performed using constant

load range per each test at a frequency of 5 Hz and load ratios

R of 0.5 and 0.1. A computerized data acquisition system was used

to determine the crack length by measuring the specimen compliance

through an extensometer mounted in the crack mouth of the specimen.

Data was also gathered through periodical optical readings to

ensure accurate results.

Post-Failure Analysis

Following each test, extensive metallography and fractography

was performed in order to evaluate fatigue failure processes. Each

sample fracture surface was examined in the scanning electron

microscope, with some samples also being sectioned to determine

the extent of damage.
Results

SEN Specimens, [0]8 Orientation

The crack growth behavior of the [0]a SEN specimens tested in

the insitu loading stage resulted in the crack growing primarily

in the direction normal to the applied load. The crack length

versus cycles plot (Fig 2) reveals a progressive decrease in the

fatigue crack growth rate with an increase in the crack length

followed by an eventual crack arrest.

The high magnification observations of the fatigue processes

revealed that crack initiation occurred in the first few cycles.

The initiation occurred by a failure of the partly machined fibers

at the notch tip (Fig 3). The crack front subsequently propagated

into the matrix ligament. Crack propagation in the matrix was

always associated with formation of slip bands in the cyclic

plastic zone ahead of the crack tip. The stable fatigue crack

growth was mainly confined to the matrix leaving in its wake

unbroken SCS-6 fibers to bridge the two cracked surfaces.

A sequence of high magnification micrographs revealing the

crack growth process observed for the SEN specimens is shown in Fig

4. As shown in the first micrograph in the sequence (Fig 4a) the

main crack front has already approached a fiber and has started

to grow around it. The reappearance of the crack front from behind

the fiber to the outer specimen surface was preceded by the

formation of slip bands (Fig 4a). As the crack front approached the

outer surface, the slip bands became more pronounced (Fig 4b). A

few thousand cycles later, the leading segments of the crack front

8
broke through the outer surface (Fig 4c). These crack segments were

visible at maximum load (Fig 4c), however as the load was released

the crack segments closed making it impossible to detect the crack

tip region, even at high magnifications (Fig 4d). After further

crack growth in the matrix ligament (Fig 4e) the entire ligament

failed leaving in its wake unbroken fibers to bridge the crack.

Once the entire matrix ligament between the two fibers was broken,

a significant increase in the crack opening displacements (COD) in

the matrix ligament occurred due to the residual tensile stresses

present in the matrix (Fig 4f). At that point the cracked matrix

ligament remained open during the entire loading cycle. The

behavior described above continued throughout the test until crack

arrest was reached, resulting in four unbroken fiber rows bridging

the two cracked surfaces (Fig 5).

In cases where the fibers were partly exposed during the

polishing process, some secondary cracking initiated at the

fiber/matrix interface ahead of the crack tip. This secondary

cracking was confined predominantly to the near surface region in

the immediate vicinity of the exposed fiber. There was very little

interaction observed between these small surface cracks and the

main crack. Overall their effect on the crack growth behavior of

the main crack was negligible.

Using high magnification micrographs of the near-crack tip

region, crack opening displacements were measured as a function of

the applied load and the distance behind the crack tip, t. One of

the many data sets obtained is shown in Fig 6. Shown in the figure
is the measured COD range for a bridged crack as a function of

distance behind the crack tip, Au(t)B , and analytically derived

COD range for an unbridged SEN specimen, _u,_, at the same applied

stress intensity range, AKap p The crack opening displacement

range, Au(t), s for an unbridged SEN specimen was calculated by the

use of weight functions (4) as a function of distance behind the

crack tip using the following integral equation:

Au(t).B=l_ fh(t,±0,C)_Kap p dc (i)

E

where E is the modulus of the composite, h is the Bueckner weight

function, c is the crack length and dc is the incremental crack

length. The difference between the measured COD's and the predicted

COD's for an unbridged SEN specimen is significant (Fig 6) and can

be used to quantify fiber bridging effect as will be shown later

on.

The extensometer mounted in the crack mouth of the specimen

showed no apparent change in the compliance throughout the entire

duration of the test. In comparison, according to the calculations

performed, an unbridged SEN specimen would have been expected to

experience an almost three-fold increase in the compliance for the

amount of crack growth that has taken place.

CT Specimens, [0]s Orientation

In contrast to the SEN specimens, the overall direction of

fatigue failure for [0]8 oriented CT specimens was parallel to the

loading direction and perpendicular to the machined notch, as shown

in Fig 7. The FCG data for the [0]8 specimens is shown in Fig 8 in
the form of crack length a vs number of cycles N (Fig 8a) and crack

growth rate da/dN vs a (Fig 8b). The data was not analyzed in

terms of the applied _K due to the lack of KI and Kn solutions

for a crack in a CT specimen growing parallel to the loading

direction. The FCG data, shown in Fig 8, does reveal an increase

of the FCG rate with an increase in the crack length indicating a

probable increase of the crack tip stress field with crack growth.

A fractographic evaluation was performed to reveal the fatigue

failure processes which resulted in this rather unusual behavior.

The first row of fibers at the machined notch which were partly

damaged by the machining process (Fig 9) failed preferentially.

The crack front continued for a short distance to grow in the

plane of the first fiber row parallel to the loading direction (Fig

7) before turning and growing perpendicular to the loading

direction through the matrix ligament. As further seen in Figs 7

and 9, after the crack front reached the undamaged second row of

fibers it again started to grow along the loading direction. After

the crack reached a critical length, the specimens failed through

debonding of the fibers and tearing of the matrix.

The stress field of the main crack resulted in the initiation

of number of secondary cracks ahead of the crack tip (Fig i0). As

shown in the figure, the cross section of the specimen revealed a

multitude of small secondary cracks in various stages of

coalescence. Careful examination of the figure showed that the as

yet uncoalesced cracks have all initiated in the interface region.

Upon further fatigue loading, these small interface cracks

g
propagated into the matrix ligaments and Joined.

The microcracking appears to have initiated typically in the

fiber/matrix interface region. A closer examination revealed a

somewhat greater tendency for the cracks to initiate in the carbon

coating of the fibers (Fig ii) rather than the reaction zone

adjacent to the coating. As shown in Fig lla, a crack presumably

initiated between the carbon coating layers, cracked through the

reaction zone and entered the matrix (crack width in the matrix is

exaggerated due to the preferential etching of the matrix material

adjacent to the crack). Irregardless of the precise crack

initiation location, the microcracks usually tended to encompass

the carbon coatings as well as the brittle intermetallic reaction

zone (i.e. the entire interface region). Post failure analysis

revealed carbon coating regions adhering to the matrix after the

fiber has broken away, again indicating that cracking must have

occurred between the fiber and the coating (Fig llb). These results

support recent findings of Gabb et al. (5) who also identified the

carbon coating as a site for preferential crack initiation. Johnson

et al. (6) while not differentiating between the carbon layer and

the reaction zone, also suggested that the weak interface region

is responsible for degradation of fatigue life of this composite.

Post failure fractographic analysis of the CT [0]8 composite

revealed that the interface region was not only the site for

preferential crack initiation but was also the controlling factor

which influenced crack propagation behavior. Throughout the test

the crack growth occurred preferentially in the interface region.

I0
The crack growth process consisted of the continuous cracking of

the interface region at or ahead of the crack tip, followed by

crack growth into the matrix ligaments and subsequent crack

coalescence. This phenomena is illustrated in Fig 12. In

particular, note in the fractograph (Fig 12b) of a high da/dN

region just prior to final failure, that stable crack growth has

occurred in the areas adjacent to the interface. However, the

remaining part of the matrix ligament failed through void

coalescence during the catastrophic failure. This fractograph

clearly shows the direction of localized crack growth to be from

the interface region to the matrix ligament.

The crack propagation within the matrix was transgranular in

nature (Fig 12). The active fatigue failure mechanisms in the

matrix included evidence of striation formation, slip band

deformation and some cleavage. Some secondary cracking was present

in the matrix foil laminate boundaries perpendicular to the main

crack front.

CT Specimens, [90]8 Composite and Unreinforced Matrix Material

Contrary to the behavior of the CT [0]8 specimens, for both

the [90]8 composite and the unreinforced matrix material CT

specimens the crack growth direction was perpendicular to the

applied load. The comparison of the FCG behavior for the two

materials is shown in Fig 13. As shown for the AK range tested,

the crack growth rates for the [90]8 oriented composite are always

considerably higher than those of the unreinforced material.

Possible reasons for this type of behavior will be discussed later.

ii
The FCG data in Fig 13 is shown for both the heat treated and

as-received conditions. For the limited number of tests conducted,

there appears to be no appreciable effect of heat treatment on the

FCG behavior of the unreinforced material. However, in the [90]8

composite the heat treated material exhibits higher FCG rates in

at least the high _K region.

There were many similarities between the FCG processes

observed for the [90]e CT specimens and those already described for

the [0]8 CT specimens. For the [90]s oriented composite, as in the

[0]8 , the crack initiation occurred mostly in the fiber/matrix

interface regions of fibers damaged by the machining process.

Figure 14 shows a notch area where the fibers have pulled away from

the matrix, leaving in its wake portions of the carbon coating.

This indicates that crack growth was accompanied by the cracking

of the carbon coating. In a manner very similar to the [0]8

oriented material, crack growth occurred preferentially in the

interface region, followed by growth into the matrix ligaments (Fig

15). Again, the matrix deformed mainly by striation forming

mechanisms and slippage. In the final stages of cracking, the same

mechanisms of debonding and matrix tearing were observed as in the

[0]8 CT specimens. There were no major differences in the failure

processes between the heat treated and as-received composite.

Discussion

The results obtained in this study pose a number of important

questions: i) Why is the behavior of the [0]8 oriented SEN and CT

specimens so radically different?; 2) What is the effective

12
driving force for crack growth in the SEN specimens ?; and 3) Why

are the fatigue mechanisms of the [0]8 and [90], oriented CT

specimens so similar?

As mentioned earlier, the crack growth direction was parallel

to the loading direction for the [0]8 CT specimens and normal to

the loading direction in the [0]8 SEN specimens. A comparison of

the simple bending and shear stresses at the notch root for the two

specimen geometries may explain the difference in behavior. As

shown in Fig 16 the bending and shear stresses were respectively

13 and 2 times greater for the CT specimens in comparison to the

SEN specimens. These higher bending stresses were able to

preferentially crack the interface region and drive the crack front

in the direction parallel to the loading direction. In the case of

the SEN specimens, the considerably lower bending stresses and

somewhat lower shear stresses were inadequate to crack the

fiber/matrix interface region. The result was that the crack growth

in the SEN specimens was confined to the matrix and bypassed the

fibers leaving them unbroken in the crack wake with little evidence

of failure in the fiber/matrix interface. As the crack grew longer

in the SEN specimens, fiber bridging acted to reduce the bending

stresses at the crack tip and thus the crack growth direction

remained normal to the loading crack.

Fiber bridging in the SEN specimens reduced the near tip crack

displacements which are the driving force for crack extension. As

the crack grew longer, more fibers bridged the crack wake. This

resulted in a progressive decrease in the fatigue crack growth

13
following:

AuNs _ _K_pp

or for a more general case:

rates and finally lead to full crack arrest (Fig 2). The very

efficient transfer of the load from the cracked matrix to the

unbroken fibers resulted in no apparent change in compliance as

measured by the extensometer mounted in the crack mouth. This

points to the inadequacy of the far field measuring devices to

characterize the fatigue crack growth behavior of a complex

material such as this composite.

The use of the fatigue loading stage inside the SEM allowed

precise measurements of the near tip crack opening displacements

to be obtained. These measurements were used to estimate the extent

of fiber bridging and to calculate the effective driving force in

the [0]8 SEN specimens. For a given crack length c and distance

behind the crack tip t, eq (i) can be shown to result in the

(2)

Au a AK (3)

Since according to eq.(3) crack tip displacements are proportional

to AK, and both the bridged and unbridged displacements are known,

the following equation for an effective crack driving force can be

applied:

In Fig 17

AK,_: = ___uB- AK,pp (4 )

_usB

the fatigue crack growth rate data for a SEN

specimen is plotted as a function of both the global AK_pp parameter

and the calculated _Ke:: parameter. For comparison purposes also

14
shown in the figure is the FCG rate data obtained from

unreinforced matrix CT specimens. The FCG rates for the [0]8 SEN

specimens decreased as a function of the AKap p. For the same AKapp,

the FCG rates were three orders of magnitude slower for the [0]_

composite than the unreinforced material. This large difference in

the measured FCG rates is due almost entirely to the crack bridging

experienced by the SEN specimens. Some of the SEN fatigue crack

growth data was recalculated in terms of the AK._ parameter to

account for crack bridging. As shown in Fig 17, the corrected [0]8

data is significantly closer to the data trends exhibited by the

unreinforced matrix material. For the data points for which the

_K,_ parameter was calculated, the stress intensity shifted from

approximately 30 MPa_m to 6-8 MPa_m. A stress intensity range of

3-4 MPa_m produced comparable crack growth rates in the

unreinforced matrix material. The remaining difference between the

actual FCG rates in the unreinforced matrix material and the FCG

rates of the [0]8 SEN specimens, measured in terms of _K,_:, may in

part be due to environmental effects. The unreinforced matrix

material was tested in the ambient environment while the [0]8

composite was tested in vacuum. In the low crack growth region

ambient environment has been shown to increase the FCG rates of

titanium based alloys (7). Also, the COD measurements were

obtained on the specimen surface and they may not fully represent

crack tip opening displacements through the specimen thickness.

The Au b measurements used to calculate the _K,:: parameter were

obtained at an approximate distance of 100 _an behind the crack tip.

15
In order to check the sensitivity of the calculated AK,:: parameter

to the location of the _u b measurements, the AKe:: parameter was

calculated using displacement measurements obtained at various

distances behind the crack tip. It was determined that as long as

the measurements where within a relatively small distance from the

crack tip, the AKe:: parameter remained essentially the same.

The difference between the _Kap p and aKef: is the equal to

the reduction in the applied stress intensity due to fiber

bridging. For these particular SEN tests, the bridging resulted in

a decrease of approximately 22-24 MPaJm in the stress intensity

range.

There have been many discussions in the past regarding the

applicability of linear elastic fracture mechanics to characterize

the FCG behavior of composites. However in the case of the [0]_

oriented SEN specimens, crack propagation was confined to the

isotropic matrix material leaving in its wake unbroken fibers. The

presence of the unbroken fibers in the crack wake can be thought

of as a highly rigid mechanical constraint acting to reduce the

crack tip opening displacement range in the isotropic matrix.

Linear elastic fracture mechanics is well suited to describe this

type of behavior.

One of the major findings of the study was the close

similarity of the fatigue failure processes for the two different

composite orientations tested using CT specimen geometry (compare

Figs 12 and 15). The most important event which lead to this

similarity in the fatigue behavior was the preferential cracking

16
of the interface region, in particular the carbon coatings

surrounding the fibers. The cracking of the interface region

occurred for both CT orientations tested, and promoted confinement

of the crack growth to the fiber direction. While crack growth in

the fiber direction is not surprising in the [90]8 oriented CT

specimens, it is somewhat surprising in the [0]8 CT specimens. For

the [90]a specimens, cracking of the weak interface region normal

to the applied tensile load apparently occurred first. These

microcracks later propagated into the matrix and linked together

resulting in a dominant crack growing in the fiber direction. For

the [0]8 CT specimens, with the main crack growing in the direction

of the applied load, one would expect the shear stresses to control

the FCG behavior. However, calculations of the shear and bending

stresses shown in Table II , reveal that the bending stress _xx

perpendicular to the fiber direction, is initially twice that of

the shear stress Txy parallel to the fibers, with the ratio

increasing as the crack grows further in the fiber direction. These

relatively high bending stresses must have been high enough to

initiate cracking of the interface region perpendicular to the

_xx direction. The cracking in the interface region was followed by

propagation and linkage of the cracks in the matrix ligaments. The

relatively high shear stresses in the [0]e CT specimens might have

further contributed to keeping the crack front in the fiber

direction by weakening the fiber/matrix interface. However, from

the fractographic evidence as well as the above mentioned analysis,

it is evident that the tensile loads perpendicular to the fiber

17
direction played the major role in controlling the FCG behavior of

the CT [0]8 specimens by preferentially cracking the interface

region.

A schematic illustration of the crack initiation and

propagation processes discussed above, for both composite

orientations, is proposed and shown in Fig 18. The process can be

summarized as follows: i) cracking of the interface region

perpendicular to the resulting _xx stresses; 2) local debonding

occurring in the failed interface region and growth of microcracks

into the matrix; 3) linking of the cracks in the matrix ligaments.

This is a continuing process requiring continuous cracking of the

interface region layer at or ahead of the crack front.

The FCG rate of the [90]8 CT specimens was considerably

greater than that of the unreinforced matrix for a given stress

intensity tested (Fig 13). The fatigue mechanisms within the matrix

region (i.e. formation of striations, slip deformation etc) were

very similar for both the [90]a and the unreinforced material and

thus probably did not contribute to the observed differences in the

FCG rates. However, the ease of cracking of the interface region,

in particular the carbon coating, is probably responsible for the

increase in the crack growth rates. This is consistent with the

hypothesis that the fiber/matrix interface has lower fatigue

resistance than the matrix. This weak composite constituent

therefore limited the fatigue cracking resistance of the [90]8

composite.

18
Conclusions

i. Fatigue crack growth behavior of the SCS-6/Ti-15-3 composite is

a function of specimen geometry, fiber orientation and the

interaction of local stress fields with the highly anisotropic

composite.

2. For the [0]8 SEN specimens the crack growth was normal to the

loading direction while for the [0]_ CT specimens crack growth was

along the fiber direction parallel to the loading direction. The

difference in the fatigue behavior for the two specimen geometries

was attributed to the significantly higher bending stresses in the

interface region for the CT specimens.

3. Near crack tip opening displacements were measured for the SEN

specimens. A formulation that uses both the actual bridged

displacements and analytically derived unbridged displacements was

proposed to determine the effective stress intensity parameter AK,::

to account for the effect of bridging on crack driving force. The

new parameter accounted for most of the bridging.

4. The fatigue crack growth processes for the CT specimen geometry

were controlled by the preferential failure of the fiber/matrix

interface for both composite orientations tested. The interface

proved to be the composites weakest link.

ig
References

i. Telesman, J., Kantzos P., and Brewer D., "Insitu Fatigue Loading

Stage Inside Scanning Electron Microscope" in Lewis Structures

Technology 1988, NASA CP 3003 Vol.3, 1988, pp 161-172.

2. Rhodes, C.G., and Spurling, R.A., "Fiber-Matrix Reaction Zone

Growth Kinetics is SiC-Reinforced Ti-6AI-4V as Studied by

Transmission Electron Microscopy" in Recent Advances in Composites

in the United States and Japan, ASTM STP-864, J.R. Vinson and M.

Taya,eds., ASTM, Philadelphia, 1985, pp. 585-599,.

3. Lerch, B.A . , Gabb, T.P . , and MacKay, R.A., "A Heat Treatment

Study of the SiC/Ti-15-3 Composite System", NASA TP 2970, 1990.

4. Bueckner, H.F., "Weight Functions for the Notched Bar", Z.

Angew. Math. Mech., Vol 51, 1971, pp. 97-109.

5. Gabb, T.P, Gayda, J., and Mackay R.A., "Isothermal and

Nonisothermal Fatigue Behavior of a Metal Matrix Composite".

Journal of Composite Materials, Voi.24, No.6, June 1990, to be

published.

6. Johnson, W.S., Lubowinski, S.J., and Highsmith, A.L,"Mechanical

Characterization of Unnotched SCS_/Ti-15-3 Metal Matrix Composite

at Room Temperature" in Thermal and Mechanical Behavior of Ceramic

and Metal Matrix Composites, eds., Kennedy, Moeller and Johnson,

ASTM STP 1080, 1990, to be published.

7. Shih, T.T and Wei, R.P., "Load and Environment Interactions in

Fatigue Crack Growth", in Prospects of Fracture Mechanics, eds.,

G.C. Sih, H.C. van Elst, and D. Broek, Noordhoff, 1974, pp. 231-

248.

G

20
TABLE I. - TEST MATRIX

Material tested

CT, [018

SEN, [0] 8

CT, [901_

CT, unreznforced

matrix

CT, unreinforced

matrix

R Ratio

0.I

.5

.i

.i

Specimens tested

As-received Heat treated

2 1

2

1 1

1 1

.5 2

TABLE If. - A COMPARISON OF BENDING AND SHEAR

STRESSES FOR THE CT, [0] 8 SPECIMENS

AT INITIATION AND FAILURE

Initiation

Failure

Maximum bending

stress,

MPa

90

220

Maximum shear

stress,

MPa

45

65

21
ORIGINAl.: PAGE'

BLACK &ix113._HITE EHO.T_OGRAP_H

(a) AS RECEIVED-ETCHED TO REVEAL MICROSTRUCTURE.

(b) DETAILED VIEW OF SCS-G FIBER-HEAT TREATED.

FIGURE 1. - SCS-6/Ti-15-3 COMPOSITE.

22
1.9

1.8

1.7

E 1.6 E

E 1.s

_J

"_ 1.4

1.2

1.1

1.0

-- _ T_O

_

__ O

SIDE

[] B

1 I

50 100 150 _0 250 300 350 q_

CYCLES x 103

FIGURE 2. - CRACK LENGTH VERSUS CYCLES FOR THE [018 ORIENTED

_N SPECI_N.

ORIGINAL PAGE"

BLACK A_ _N.HITC£HO_'ZOG__H

PRIOR TO TESTING AFTER 150 CYCLES

FIGURE 3. - CRACK ]NIHATION IN THE [0]8 ORIENTED SEN SPECIMEN.

23
ORIGINAL PAGE'

BLACK A___ ._'/iLTE. PJI.O_QGRJ_PJ-j

(=)

AFTER 100 000 CYCLES MAXIMUM LOAD AFTER 103 085 CYCLES MAXIMUM LOAD

20 pm

107 044 CYCLES MAXIMUM LOAD 107 044 CYCLES MINIMUM LOAD

so _m

110 605 CYCLES MAXIMUM LOAD 112 894 CYCLES MAXIMUM LOAD

FIGURE 4. - CRACK GROWTH SEQUENCE IN THE IOI 8 SEN SPECIMEN RESULTING IN FIBER BRIDGING OF

THE CRACK WAKE,

358 287CYCLES MAXIMUM LOAD

FIGURE 5. - GENERAL APPEARANCE OF THE FATIGUED [0]8 SEN SPECIMENS AS OBSERVED IN

LOADING STAGE INSIDE AN SEM.

24
DISPLACEMENTPROFILE VERSUS DISTANCE FRON CRACKTIP

C,

X

E

1,2 --

IJNBRIDGED _/

1.0 -- 0 DELT_

.8 --

.4"6--/ _ D D [] [:3 D

.2 / [] [] []

DD

I 1 ____J__l 1 I . I

0 1 2 5 4 5 6 7 8

DISTANCE FROM CRACKTIP, m x I0 -_

FI6tlRE6, - £ONPARISDN OF ThE NEAStlREDDEtT_ COD'S WITH

ANALYTICALLY CALCULATED COD'S FOR AN UNBRIDGED SEN SPECIMEN.

ORIGINAL PAGE

BLACK AAI_. _HLT£ P3tOIOGl_p3_

LOADING

DIRECTION

_0.2 n_

[0] 8 ORIENTATION.

FIGURE 7. - CRACKGROgTHPARALLEL TO TIlE LOADING DIRECTION FOR THE CT [0] 8 ORIERTATIOfl.

ORIGINAL PAGE IS

OF POOR QUAUTY

25
E

E

1-04

E

E. 1-05

2 _

1 m

1-06

SIC/T1-15-3

R,T.

5Hz, R=O.S

IX

],,

Z_ A

IX

Z_

I [ I

1O0 200 300

CYCLES, THOUSANDS

(a) CRACK LENGTH VERSUS CYCLES.

I

400

_- SiC/Ti-15-3

R.T.

" 5Hz, R=0.5

m

[]

[]

[]

[]

I-I

1 l I m I Il,J 1 J [ J I ,I,I

1 1,0 10

CR_,CKLENGTH, mm

(b) CRACK GROWTH RATE VERSUS CRACK LENGTH.

FIGURE 8. - FATIGUE CRACK GROWTH RATE DATA FOR CT [0]8

SPECIMENS.

ORIGINAL PAGE _

Of. POOR QUAUTf

26
ORIGINAL PAGE

BLACK_LO WH_[E P_HO]_OGRAR.H

FIGURE 9, - CRACKGROWTHIN THE [0] 8 CT SPECIREN.

FIBER DIRECTION.

0.2 mm

DIRECTION OF LOADING PARALLELTO

ORIGINAL PAGE f$

OF POOR QUALITY,

27
ORIGINAL PAGE

61_CK A.tNJ3,.YYHJEg EH.OE_QG&A£_H

m

± CUT 1

©

FIGURE 10. - MIJLTIPLE CRACK INITIATION AND COALESCENCE IN THE EARLY STAC, ES

OF FATIGUE LIFE IN THE CT [0] 8 CORPOSITE. ARROWS MARK CRACKS INITIATING

IN THE INTERFACE REGION.

OR.wGINAL PAGE IS

OF POOR QUALITY

28
OR1GINAE PAGE

BLACK A.I__ WHLT_F,P_HO_T_OG-RAP_

REACTION

ZONE

CARBON

COAT! NG

SiC

FIBER

(a) CRACKING ERANATING FROR THE CARBONLAYER.

FIBER C_OOVE_, ..--CARBON COATING

ORIGINAL PAGE IS

OE POOR QUALITY

(b) CARBONLAYER SEPARATEDFROIq THE FIBER DURING FATIGUE.

FIGURE 11. - PREFERENTIAL CRACKING OF THE CARBONLAYER.

29
ORIGrNAL PAGE

BLACKAND WHITE.P-HO_rOGP,_._

MICROSCOPIC GROWTffDIRECTION

(a) LO_da/dN [0] 8 .

I MACROSCOPIC GROWTH DIRECTION

MICROSCOPIC GROWTH DIRECTION

ORIGINAL PAGE IS

OF POOR QUAUTY

OVERLOAD

(b) HIGH da/dN [0]8,

FIGURE 12. - MICROSCOPIC CRACK GROWTH DIRECTION FROM THE INTERFACE INTO THE MATRIX

LIGAMENT IN TIIE[0]8 CT SPECIMENS,

30
)

1,0-06

1.0-07

1.0-08

1.0-09

3

{90] 8 • H.T, R = 0.5

[90] 8, A.R, R 0.5

"_ U.M.M., H.T, R = 0,5

U.M.M., A.R., R = 0.5

0 U.M.I4.,A.R., R = 0.I

$l.

I.

¢

$

i)

4. ÷÷

6 lO 20 30 40

DELTA K, MPa,_

FIGURE 13, - COMPARISON OF FC6 RATES FOR THE [9018 AND THE UNREINFORCED MATRIX

I_ATERIALIN THE AS-RECEIVED (A.R.) AND HEAT-TREATED (H.T.) CONDITION.

FIGURE lb. - SEPARATION OF CARBON COATING FRO@I TIIEFIBER IN [9018 CT SPECIMEN. IN THE

OENTER OF THE FRACTOGRA_ NOTE A CRACK BEIWEEN C/_RBONCOATING AND THE INTERFACE.

ORIGINAL PAGE _8

05 POOR QUALITY

ORIGINAl/ PAGE

BLACK A!_ _fll,T£ PHOTOGRARH

31
OR1GINA[ PAGE

BLACK AND WHITE PHOTOGRAPH

R|CRO_IC GROWTH DIRECTION

T OVERALL

GROWTH

DIRECTION

(a) LOWda/dN.

ORiG;NAL PAGE iS

POOR _UAUTY

t

32
CT SPECIMEN

INITIAL K = 18 MPa_

MAX BENDING STRESSOxx = 90 MPa

SHEAR STRESS TXy = 45 MPa

* CRACKINGPARALLELTO THE LOADING AXIS

A

A Oxx

"XX 1

SEN SPECIMEN

INITIAL K = 18 MPaq_--

MAX BENDING STRESS Oxx = 7 l_°a

SHEAR STRESS T = 23 MPa

xy

* CRACKING NORMAL TO THE LOADING AXIS

FIGURE 16. - COMPARISON OF BENDING AND SHEAR STRESSES FOR THE TWO SPECIMEN GEOMETRIES.

1.0 -06

1.0 -07

1,0 -08

1.0 -09

1.0 -10

0 Ti-15-3 MATRIX

SEN _Ka.^)w <>o

• SEN (_Kef f) ,jI

I

oy,

.,_..... Ii _-_-__-_._

IR=

®

B_

m_

o

0

0 0--._0

I I i I I I tl I m I J

zl 6 10 20 30 qO 50

&Kef f OR _K, MPa_

FIGURE i7. - SEN CRACKGROWTHDATA CO/4PAREDTO UNREINFORCEMATRIX MATERIAL IN TERMSOF ,}.Kapp AND THE PROPOSED,_,Keif PARAMETER.

ORIGINAL PAGE 18

OE POOR QUALITY

33
I Oyy FOR [90]8 SPECIMENS

2 TXy --_0 FOR [90]8 SPECIMENS

(a) CRACK INITIATION IN THE INTERFACE. (b) CRACK GRowTH AND LOCAL DEBONDING.

c

J -'ll"".-...J.., / OF APPLIED CY(_LEs

(c) CRACK COALESCENCE IN THE MATRIX,

FIGURE 18. - PROPOSED CRACK INITIATION AND CRACK PROPAGATION PROCESSES FOR THE CT SPECIREN 6EO_llIY.

ORIGINAL PAGE IS

O_. POOR QUALITY

34
NationalAeronauticsand Report Documentation Page

SpaceAdministration

1. Report No. 2. Government Accession No. 3. Recipient's Catalog No.

NASA TM- 103095

4. Title and Subtitle 5. Report Date

Fatigue Crack Growth in a Unidirectional SCS-6/Ti-15-3 Composite

7. Author(s)

Peter Kantzos, Jack Telesman, and Louis Ghosn

9. Performing Organization Name and Address

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135-3191

12. Sponsoring Agency Name and Address

National Aeronautics and Space Administration

Washington, D.C. 20546-0001

6. Performing Organization Code

8. Performing Organization Report No.

E-54 i 3

10. Work Unit No.

510-01-01

11. Contract or Grant No.

13. Type of Report and Period Covered

Technical Memorandum

14. Sponsoring Agency Code

15. Supplementary Notes

Prepared for the Third Symposium on Composite Materials: Fatigue and Fracture sponsored by the American

Society for Testing Materials, Orlando, Florida, November 6-9, 1989. Peter Kantzos and Jack Telesman, NASA

Lewis Research Center. Louis Ghosn, Sverdrup Technology, Inc., NASA Lewis Research Center Group,

Cleveland, Ohio 44135.

16. Abstract

An investigation was conducted to characterize and model the fatigue crack growth (FCG) behavior of a SCS-6/T[-15-3 metal

matrix composite. Part of the study was conducted using a fatigue loading stage mounted inside a scanning electron microscope (SEM). This unique facility allowed high magnification viewing of the composite fatigue processes and measurement of

the near crack tip displacements. The unidirectional composite was tested in the [018 (i.e., longitudinal) and [90] 8 (i.e.,

transverse) orientations. For comparison purposes unreint_rced matrix material produced by the identical process as the reinforced material was also tested. The results of the study reveal that the fatigue crack growth behavior of the comD_site is a

function of specimen geometry, fiber orientation and the interaction of local stress fields with the highly anisotropic composite.

In the case of [0]8 oriented single edge notch (SEN) specimens and [_90]8_oriented compact tension (CT) specimens, the crack

growth was no-rm-al to the loading direction. However, for the [0]s CT specimens the crack grew mostly parallel to the loading

and the fiber direction. The unusual fatigue behavior of the I0]8-CT specimens was attributed to the specimen geometry and

the associated high tensile bending stresses perpendicular to the fiber direction. These stresses resulted in preferential cracking

in the weak interface region perpendicular to the fiber direction. The interface region, and in particular the carbon coating

surrounding the fiber proved to be the composites weakest link. In the [0] 8 SEN the crack growth was confined to the matrix

leaving behind unbroken fibers which bridged the cracked surfaces. As the crack grew longer, more fibers bridged the crack

resulting in a progressive decrease in the crack growth rates and eventual crack arrest. The actual near crack tip displacement

measurements were used in a proposed formulation for a bridging-corrected effective crack driving force, _Ke|- f. This parameter was able to account for most of the experienced bridging and correlated the !0] s SEN fatigue crac]_ growth data

reasonably well.

1'7. Key Words (Suggested by Author(s}l

Metal matrix composite

Fatigue crack growth

Crack bridging

Fatigue mechanisms

18. Distribution Statement

Unclassified- Unlimited

Subject Category 26

19. Security Classif. (of this report) 20. Security Classif. (of this page) 21. No. of pages

Unclassified Unclassified 36

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