NASA Technical Memorandum 102332 

. 

Fatigue Crack Growth Study of 

SCS6/Ti- 15-3 Composite 

Peter Kantzos and Jack Telesman 

Lewis Research Center 

Cleveland, Ohio 

August 1989 

(BASA-TB-fO2332) PATIGUE CRACK GBOYTB STUDY N89-26989 

Research Center) 20 p CSCL 11F 

i OF SCS6/Ti-15-3 COHPOSITE (NISIL. Levis 

Unclas 

63/26 0225949 
FATIGUE CRACK GROWTH STUDY OF SCSg/TI-I5-3 COMPOSITE 

Peter Kantzos and Jack Telesman 

National Aeronautics and Space Administration 

Lewis Research Center 

Cleveland, Ohio 44135 

d 

0 

UY 

I 

w 

SUMMARY 

A study was performed to determine the FCG behavior and the associated 

fatigue damage processes in a [Olg and [go18 oriented SCS6/Ti-15-3 composite. 

Companion testing was also done on identically processed Ti-15-3 unreinforced 

materi a1 . 

The active fatigue crack growth failure processes were very similar for 

both composite orientations tested. 

growth was along the fiber direction. It was found that the composite constituent most susceptible to fatigue damage was the interface region and in particular the carbon coating surrounding the fiber. 

region lead to crack initiation and also strongly influenced the fatigue crack 

growth (FCG) behavior in this composite. 

was apparently driven by normal stresses perpendicular to the fiber direction. 

For both orientations, fatigue crack 

The failure of the interface 

The failure of the interface region 

The FCG rates were considerably higher for the [go18 oriented CT specimens 

in comparison to the unreinforced material. 

nario in which the interface has lower fatigue resistance than the matrix, 

causing lower composite fatigue resistance. 

ite could not be directly compared to the [go18 composite but were shown to 

increase with an increase in the crack length. 

This is consistent with the sceThe FCG rates of the [Olg composINTRODUCTION 

The new generation of the 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 

consisting of a Ti-15V-3Cr-3A1-3Sn matrix reinforced by continuous, unidirectional Sic fibers, was investigated. 

The fatigue behavior and life prediction of MMC’s must be understood 

before they can be widely utilized in aerospace applications. 

tively little information is available concerning the fatigue behavior of continuous fiber reinforced composites. 

determine the microstructural processes controlling fatigue crack initiation 

and propagation in this system. Emphasis has been placed on the identification of the composite constituent most susceptible to fatigue damage. 

However, relaThe objective of this study has been to 

MATER I A L 

The composite used in this study is a Ti-15V-3Cr-3A1-3Sn (Ti-15-31 a1 

matrix reinforced by 145 pm average diameter, continuous Sic (sCs6) fibers 

All the composite specimens used in this study were obtained from a single 

OY 
eight-ply, unidirectional panel manufactured by Textron Specialty Materia 

Division. The panel size was 30 by 30 by 0.21 cm. Also tested were spec 

machined out of a 30- by 30- by 1-cm panel made without fibers, using the 

Ti-15-3 foil and same consolidation techniques employed in composite 

production. 

The matrix is a metastable beta phase alloy, chosen because it can b 

S 

mens 

same 

cold 

rolled to very thin sheets. 

the composite. Sample micrographs of both, the unreinforced and reinforced 

composite are shown in figures 1 and 2. The Sic fiber is surrounded by a complex multilayer structure (fig. 2(b)) which consists mainly of a multilayer 

carbon coating approximately 3 pm thick and an approximately 0.5-1 pm thick 

rcaction zone consisting mainly of brittle intermetallic phases (ref. 1). The 

combination of the carbon layers and the reaction zone will henceforth be 

referred to as the interface region. 

This enables a more cost effective processing of 

HEAT TREATMENT 

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

The heat treatment consisted of 24 hr at 700 "C in vacuum. The purpose of the 

heat treatment was to precipitate out most of the alpha phase. 

ment did not produce any additional noticeable fiberjmatrix interactions and 

the reaction zone remained unchanged (ref. 2). 

The heat treatMECHANICAL TESTING 

Testing was performed using compact tension (CT) specimens, such as the 

one shown in figure 3. The specimens were tested in two orientations: (1) 

longitudinal, [Ola, with fibers oriented parallel to the loading direction; 

and (2) transverse, [9018, with fibers perpendicular to the loading direction. 

All machining was done using diamond tip tools. Prior to testing, all 

specimens were polished to enhance crack detection. The polishing process was 

automated and developed at NASA Lewis Research Center (ref. 3). 

The test matrix is shown in table I. Testing of the CT specimens was performed at room temperature and ambient environment using a computer controlled, 

closed loop, servohydraulic machine. The tests were performed using constant 

load range at a frequency of 5 Hz and 

by measurjng compliance of the specimen through a clip gage mounted in the 

crack mouth of the specimen. Data was also gathered through periodical optical 

readings to ensure accurate results. 

R = 0.5. The crack !ength was obtained 

POST-FAILURE ANALYSIS 

Following each test, extensive metallography and fractography was performed in order to evaiuate fatigue failure orocesses. Each sample fracture 

surface was examined in the scanning Zlectron microscope, with some samples 

also being sectioned to determine the extent of damage. 

2 
RESULTS 

Longitudinal, Cola Spec mens 

The overall direction of fatigue failure for COI8 oriented specimens was 

parallel to the loading direction and perpendicular to the machined notch, as 

shown in figure 4. The fatigue crack growth (FCG) data for the [Olg specimens 

is shown in figure 5 in the form of crack length a versus number of cycles N 

(fig. 5(a>) and crack growth rate daldN versus a (fig. 5(b)>. The data was 

not analyzed in terms of the applied AK due to the lack of KI and KII 

solutions for a crack in a CT specimen growing parallel to the loading direction. The FCG data, shown in figure 5, 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 were partly damaged by the machining process 

(fig. 6). The damaged fibers at the notch produced multiple crack initiations 

at the fiberlmatrix interface regions. Upon further fatigue loading, these 

small interface cracks propagated into the matrix ligaments and joined 

(fig. 7(a)). When the coalesced cracks encountered the next row of fibers, 

the crack trajectory changed and the cracks started growing along the fiber 

direction as shown in figure 6. In these early stages of fatigue life there 

appears to be an extensive amount of microcracking and crack coalescence 

(fig. 7(a)). In the later stages of the fatigue life, the microcracks coalesced to form one dominant crack. Once this crack formed, the damage appears 

to be mostly confined to the main crack front with little evidence of secondary cracking (fig. 7(b>). After the dominant crack reached a critical length, 

the specimens failed through debonding of the fibers and tearing of the matrix. 

The microcracking appears to have initiated typically in the fiberlmatrix 

interface region. The crack initiation sites were not evenly distributed with 

an observed preference for initiation to occur in a thickness direction 

(fig. 7(a)). 

the cracks to initiate in the carbon coating of the fibers (fig. 8(a>) rather 

than the reaction zone adjacent to the coating. As shown in figure 8(a), crack 

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 both the carbon coating as well as the brittle intermetallic reaction zone (i.e., the entire interface region). Post failure analysis revealed 

suspected carbon coating regions adhering to the rnatrix after the flber has 

broken away, again indicating that cracking must have occurred between the 

fiber and the coating (fig. 8(b)>. These results support recent findings of 

Gabb et al. (4) who also identified the carbon coating as a site for preferential crack initiation. Johnson et al. (5) while not differentiating between 

the carbon layer and the reaction zone, also noted that the weak interface 

region is responsible for degradation of fatigue life of this composite. 

A closer examination revealed a somewhat greater tendency for 

Post failure fractographic analysis 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 entire 

test the crack growth occurred preferentially in the interface region. The 

3 
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 crack coalescence. This phenomena is illustrated in figure 9. 

In particular, note in the fractograph (fig. 9(b)) of a high da/dN region 

just prior to final failure, that stable crack growth has occurred in the areas 

adjacent to the interface. 

failed through void coalescence during the catastrophic failure. 

graph clearly shows the direction of localized crack growth to be from the 

interface region to the matrix ligament. 

However, the remaining part of the matrix ligament 

This fractoThe crack propagation within the matrix was transgranular in nature 

(fig. 9). The active fatigue failure mechanisms in the matrix were dependent 

on grain orientation, and included evidence of striation formation, slip 

aecohesion and/some cleavage. Some secondary cracking was present in the 

matrix foil laminate boundaries perpendicular to the main crack front. 

Transverse Composite, [901g, and Unreinforced Matrix Material 

Contrary to the behavior of the [Olg specimens, for both the [9018 composite and the unreinforced matrix material the crack growth direction was perpendicular to the applied load. The comparison of the FCG behavior for the two 

materials Is shown in figure 10. As shown for the AK range tested, the crack 

growth rates for the C9018 oriented composite are always considerably higher 

than those of the unreinforced material. Possible reasons for this type of 

behavior will be discussed later. 

The FCG data in figure 10 is shown for both the heat treated and asreceived 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. 

exhibits hlgher FCG rates in the high AK region. 

However, in the [go18 composite the heat treated material 

There were many similarities between the FCG processes observed for the 

c9018 specimens and those already described for the [Olg specimens. For the 

[go18 oriented composite, just as in the [OIg, the crack initiation occurred 

mostly in the fiber/matrix interface regions of fibers damaged by the machining process. Figure 11 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 [Olg oriented material, crack growth 

occurred preferentially in the interface region, followed by growth into the 

matrix ligaments (fig. 12). Again, the matrix deformed mainly by slip decohesion and striation forming mechanisms. In the final stages of cracking, the 

same mechanisms of debonding and matrix tearing were observed as in the COlg 

speclmens. There were no major differences in the failure processes between 

the heat treated and as-received composite. 

The fatigue failure in the unreinforced matrix material tended to be 

mostly transgranular and consisted mostly of slip deconesion and striation 

forming mechanisms in the stable crack growth region. 

ure and void coalescence occurred in the high AK and overload regions. The 

fatigue failure processes were thus usually similar to those encountered In 

the composite material. 

Some interqranular fail4 
The unreinforced matrix material exhibited some delamination at the foil 

boundaries (fig. 13). Delamination can be attributed to the poor bonding 

during consolidation caused by the presence of contaminants at the foil boundaries. Delamination increased with an increase in the applied AK. 

DISCUSSION 

One of the major findings of the study was the close similarity of the 

fatigue failure processes for the two different composite orientations (compare 

figs. 9 and 12). The most important event which lead to this similarity in the 

fatigue behavior was the preferential cracking of the interface region, in part’cular the carbon coatings surrounding the fibers. The cracking of the interface region occurred for both orientations tested, and promoted crack growth 

being confined to the fiber direction. While crack growth in the fiber direction is not to surprising in the [go18 oriented specimens, it is somewhat surprising in the EO18 specimens. For the [go18 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 up resulting in 

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

simple shear and bending stresses, shown in table I1 and figure 14, reveal that 

the bending stress uXx perpendicular to the fiber direction, is initially 

twice that of the shear stress T~ parallel to the fibers, with the ratio 

high bending stresses must have been high enough to initiate cracking of the 

interface region perpendicular to the uXx direction (fig. 7(a>). 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 

COlg 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 appears 

that tensile loads perpendicular to the fiber direction played the major role 

in controlling the FCG behavior of this composite system by preferentially 

cracking the interface region. 

discussed above, for both composite orientations, is proposed and shown in 

figure 13. The process can be summarized as follows: (1) cracking of the 

interface region perpendicular to the applied tensile loads; (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. 

increasing as the crack grows furt $: er in the fiber direction. These relatively 

A schematic illustration of the crack initiation and propagation processes 

The FCG rate of the [go18 composite was considerably greater than that of 

the unreinforced matrix for a given stress intensity tested (fig. 10). The 

fatigue mechanisms within the matrix region (i.e., slip decohesion, formation 

of striations, etc.) were ve y similar for both the [go18 and the unreinforced 

material and thus probably d d not contribute to the observed differences in 

the FCG rates. However, the ease of cracking of the interface region, in particular the carbon coating, s probably responsible for the increase in the 

crack growth rates. This is consistent with the hypothesis that the fiber/ 

5 
matrix interface has lower fatigue resistance than the matrix. This weak composite constituent therefore limited the fatigue cracking resistance of the 

compos i te. 

CONCLUSIONS 

A study was performed to determine the FCG behavior and the associated 

fatigue damage processes in a Cola and [9018 oriented SCS6/Ti-15-3 composite. 

Companion testing was also done on identically processed Ti-15-3 unreinforced 

material. The most important findings were: 

1. The fatigue crack growth processes were very similar for the two composite orientations tested, with fatigue crack growth being confined along the 

f i ber di rec t ion. 

2. The interface region (in particular the carbon coating) was the composite constituent most susceptible to fatigue damage influencing both the crack 

initiation and crack propagation behavior of the composite. 

3. The failure of the interface region was mostly the result of the 

applied bending stresses perpendicular to the fiber direction. 

4. The FCG rates were considerably higher for the C9018 oriented CT specimens in comparison to the unreinforced material, reaffirming the lower fatigue 

resistance of the fiber/matrix interface in comparison to the matrix. 

REFERENCES 

1. Rhodes 

ics is 

COPY * 

J.R. V 

2. Lerch, 

C.G.; and Spurling, R.A.: Fiber-Matrix Reaction Zone Growth KinetSIC-Reinforced Ti-6A1-4V as Studied by Transmission Electron MicrosRecent Advances in Composites in the United States and Japan, 

nson and M. Taya, eds., ASTM, Philadelphia, 1985, pp. 585-599. 

B.A.; Gabb, T.P.; and MacKay, R.A.: A Heat Treatment Study of the 

SiC/Ti-15-3 Composite System. NASA TM (to be published), 1989. 

3. Lerch’, B.A.; Hull D.R.; and Leonhardt, T.A.: As-Received Microstructure 

of a SiC/Ti-lS-3 Composite. 

4. Gabb, T.P.; Gayda, J.; and Mackay R.A.: Isothermal 

Fatigue Behavior of a Metal Matrix Composite. Subm 

Mater., 1989. 

5. Johnson, W.S., et ai.: Mechanical Characterization 

Matrix Composite at Room Temperature. NASP TM-1014 

NASA TM-100938, 1988. 

6 

and Nonisothermal 

tted to the J. Compos. 
TABLE I. - TEST MATRIX 

Mater i a1 tested 

Longitudinal , LO18 1 r Transverse, [90 18 

Unreinforced matrix I 

2 

1 

1 

TABLE 11. - A COMPARISON OF 

BENDING AND SHEAR STRESSES 

FOR THE [Olg SPECIMENS AT 

INITIATION AND FAILURE 

I I I 

Ini ti ati on 

Fai 1 ure 

FIGURE 1. - MICROSTRUCTURE OF THE UNREINFORCED MATRIX NATERIAL 

ORIGINAL PAGE 

BLACK AND WHITE PHOTOGRAPH 

7 
(a) GENERAL MICROSTRUCTURE. 

.\ 

I’ .. .. ’ 

I 

IO - mm 

(b) SCS6 FIBER. 

FIGURE 2. - AS RECEIVED COMPOSITE. 

8 GRIGINAL PAGE 

BLACK AND WHITE PHOTOGRAPH 
ORIGINAL PAGE 

BLACK AND WHITE PHOTOGRAPH 

FIGURE 3. - COMPACT TENSION (CT) SPECIMEN. 

LOADING 

DIRECTION 

[018 ORIENTATION. 

FIGURE 4. - CRACK GROWTH PARALLEL TO THE LOADING DIRECTION FOR THE [Olg ORIENTATION. 

9 
SiC/Ti-15-3 '1 R.T. 

4 5 Hz, R = 0.5 

3' 

2. - A 

a 

A 

I 

0 100 700 300 400 

CYCLES, THOUSANDS 

(a) CRACK LENGTH VERSUS CYCLES. 

1-05 '-OI 

SiC/Ti -15-3 

R.T. 

5 Hr. R = 0.5 

0 

0 

0 

0 

0 

0 

FIGURE 6. - CRACK GROWTH IN THE [OIg MATERIAL. DIRECT[ON OF LOADING PARALLEL TO FIBER 

DIRECTION. 

G2iClNAL PAGE 

lo BLACK AND WHITE PHOTOGRAPH 
E 

CUT 1 

(a) fiULTIPLE CRACK INITIATION AND COALESCENCE IN THE EARLY STAGES OF FATIGUE 

LIFE. 

CUT 2 

(b) CRACK GROWTH CONFINED MOSTLY TO ONE DOMINANT FRONT IN THE LATER 

STAGES OF FATIGUE LIFE. 

I 1 

FIGURE 7. 

FAT I GUE 

- COMPARISON OF EXTENT OF HICROCRACKING IN DIFFERENT 

LIFE IN THE IOlg ORIENTED COMPOSITE. 

STAGES OF 

11 
REACTION 

ZONE 

CARBON , 

COAT I NG 

si C 

FIBER 4 

(a) WALKING EIANATING F(C?l Tilt LARBON LAYER. 

FIBER GROOVE - ~ ,CARBON COATING 

(b) CARBON LAYER SEPARATED FROM THE FIBER DURING FATIGUE. 

FIGURE 8. - PREFERENTIAL CRACKING OF THE CARBON LAYER. 

12 
MICROSCOPIC -- GROWTH DIRECTION 

(a) LOW da/dN [Olg. 

MACROSCOPIC t GROWTH DIRECTION 

MICROSCOPIC GROWTH DIRECT I ON 

c------ 

OVE 

(b) HIGH da/dN [Olg. 

FIGURE 9. - MICROSCOPIC CRACK GROWTH DIRECTION FROM THE INTERFACE INTO THE MATRIX 

LIGAMENT IN THE [Olg SPECIMENS. 

13 
t 

8 A.R. I 

OA 

AA 

H.T. 

R.T. 

5 Hz. R = 0.5 

I I 

20 30 

DELTA K. MPA Jm 

FIGURE 10. - COMPARISON OF FCG RATES FOR THE [9018 AND THE 

UNREINFORCED MATRIX MATERIAL IN THE AS-RECEIVED (A.R.) 

AND HEAT-TREATED (H.T. 1 CONDITION. 

14 
ORIGINAL PAGE 

BLACK AND WHITE PHOTOGRAPH 

(a) REMNANTS OF THE CARBON COATING ADHERING TO THE MATRIX AFTER ITS SEPARATION FROM THE 

FIBER IN THE [9Olg COMPOSITE. 

(b) CARBON COATING REMAINING IN THE FIBER GROOVE. IN THE CENTER OF THE FRACTOGRAPH NOTE A 

CRACK BETWEEN CARBON COATING AND THE INTERFACE. 

FIGURE 11. - SEPARATION OF THE CARBON COATING FROM THE FIBER IN THE [901g SPECIMENS. 

15 
MICROSCOPIC GROWTH DIRECTION LI 

(a) LOW da/aN 

OVERALL I GROWTH 

DIRECTION 

(b) HIGH da/dN. 

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

["g SPECIRENS. 

ORJGINAL PAGE 

BLACK AND WHITE PHOTOG RAP^^ 

16 
LOU AK 

(a) SOME PLY DELAMINATION. 

OVERLOAD - HIGHAK 

(b) EXTENSIVE PLY DELAMINATION. 

FIGURE 13. - DELAPlINATION IN THE UNREINFORCED MATRIX MATERIAL. 

17 
A 

OYY 

XY 2T 

INITIAL K = 18 PlP~fi 

MAXIMUM BENDING STRESS Oxx = 90 WA 

SHEAR STRESS Txy = 45 MPA 

CRACKING IN THE SPECIPIEN PARALLEL TO THE LOADING DIRECTION 

FOR [9Olg SPECIMENS 

3 0 FOR [9013 

7 

(a) CRACK INITIATION IN THE INTERFACE. 

(b) CRACK GROWTH AND LOCAL DEBONDING. 

FIGURE 14. - DIRECTION AND MAGNITUDE OF THE BENDING AND SHEAR 

STRESSES FOR THE [Olg SPECIMENS AT THE NOTCH. 

tc) CRACK COALESCENCE IN THE MATRIX. 

FIGURE 15. - POSTULATED CRACK INITIATION AND GROWTH SEQUENCE FOR 

SiC/Ti -15-3 COMPOSITE. 

18 
m National Aeronautics and Report Documentation Page 

Space Administration 

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

NASA TM-102332 

4. Title and Subtitle 

Fatigue Crack Growth Study of SCS6/Ti-15-3 Composite 

5. Report Date 

August 1989 

7. Author(s) 

Peter Kantzos and Jack Telesman 

9. Performing Organization Name and Address 

National Aeronautics and Space Administration 

Lewis Research Center 

Cleveland, Ohio 44135-3191 

8. Performing Organization Report No. 

E-5041 

IO. Work Unit No. 

- 

I 13. Type of Report and Period Covered 

5 10-01 -01 

11. Contract or Grant No. 

12. Sponsoring Agency Name and Address 

National Aeronautics and Space Administration 

Washington, D.C. 20546-0001 

Technical Memorandum 

14. Sponsoring Agency Code 

7. Key Words (Suggested by Author@)) 

Metal matrix composite 

Fatigue 

Crack initiation 

Crack propagation 

15. Supplementary Notes 

18. Distribution Statement 

Unclassified - Unlimited 

Subject Category 26 

\ 

16. Abstract 

A study was performed to determine the FCG behavior and the associated fatigue damage processes in a [OI8 and 

[go], oriented SCS6/Ti- 15-3 composite. Companion testing was also done on identically processed Ti-15-3 

unreinforced material. The active fatigue crack growth failure processes were very similar for both composite 

orientations tested. For both orientations, fatigue crack growth was along the fiber direction. It was found that 

the composite constituent most susceptible to fatigue damage was the interface region and in particular the carbon 

coating surrounding the fiber. The failure of the interface region lead to crack initiation and also strongly 

influenced the fatigue crack growth (FCG) behavior in this composite. The failure of the interface region was 

apparently driven by normal stresses perpendicular to the fiber direction. The FCG rates were considerably 

higher for the [go], oriented CT specimens in comparison to the unreinforced matqial. This is consistent with the 

scenario in which the interface has lower fatigue resistance than the niatrix;Lcausing' lower composite fatigue 

resistance. The FCG rates of the [0l8 composite could not be directly compared to the [90], composite but were 

shown to increase with an increase in the crack length. 

9. Security Classif. (of this report) 

Unclassified 

20. Security Classif. (of this page) 

Unclassified 

1 21. No o:pg89es 22. Price' I A03 

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