Tensile damage in a symmetric [090 ]2s silicon-carbide fiber-reinforced titanium-matrix composite

PII:

SO266-353(97)00198-X

Composites Science and Technology 58 (1998) 9 15-93 I 0 1998 Elsevier Science Ltd. All rights reserved Printed in Northern Ireland 0266-3538/98 $-see front matter

ELSEVIER

TENSILE DAMAGE IN A SYMMETRIC [0/901zs SILICONCARBIDE FIBER-REINFORCED TITANIUM-MATRIX COMPOSITE W. 0. Soboyejo,“,* B. M. Rabeeh,” J. Zhangb & N. Katsubeb aDepartment of Material Science and Engineering, Ohio State University, 2041 College Road, Columbus, OH 4321&1179, USA bDepartment of Applied Mechanics, Ohio State University, Boyd Laboratory, 155 W. Woodruff Ave. Columbus, OH 43210-I 179, USA (Received 14 March 1997; revised 8 September 1997; accepted 7 October 1997)

Abstract The results of a detailed study of the eflects of composite microstructure on the micromechanisms of tensile damage in a symmetric [0/9O]z, Ti-15Al-3Cr-3AI-3Sn (Ti-I.53)ISiC (SCS-6) composite are presented. Matrix microstructure is controlled by heat treatment, which is used to produce metastable p or Widmanstatten u + p microstructures. The sequence of damage initiation and evolution at room temperature is identified using ex situ scanning electron microscopy (SEM) during incremental monotonic loading to failure. Damage mechanisms are also studied using non-destructive acoustic emission (AE) techniques, and matrix hardening is characterized using matrix micro-hardness measurements within individual plies. Idealized and actual microstructure-based jiniteelement moakls are used to rationalize the observed tendency of damage to propagate from the outer plies to the inner plies. The paper highlights the importance of simplified micromechanical models in the development of a fundamental understanding of the eflects of composite architecture on damage initiation phenomena. 0 1998 Elsevier Science Ltd. All rights reserved

and fracture15-‘7 in titanium-matrix composites with various architectures. However, our current understanding of the effects of composite microstructure on the micromechanisms of damage initiation and propagation in these composites is still very limited.14 Most of the previous studies have shown that a complex sequence of damage is associated with deformation and failure under monotonic/cyclic thermal or mechanical loading. Previous preliminary studiesi have shown that the sequence of damage is strongly affected by matrix and interfacial microstructure. Early damage nucleation has also been shown to occur in the layered interfacial regime that exists in the region in between the fiber and the matrix. This region has been shown to have a highly complex microstructure in recent high resolution transmission electron microscopy studies. 18,19 Nevertheless, there have been very few efforts designed to study the effects of composite microstructure on damage mechanism. l4 There is also a strong need for the incorporation of experimental observations of damage into micromechanics models that are used for the prediction of damage/crack growth phenomena in continuously reinforced titanium-matrix composites. The results of a systematic study of the effects of composite microstructure on the micromechanics and micromechanisms of damage in a Ti-15V-3Cr-3Al3Sn/SiC (SCS-6) composite are presented in this paper. These include the effects of matrix and interfacial microstructure on tensile damage phenomena at room temperature. The experimental observations of damage initiation are used in the formulation of microstructurebased finite-element models for the rationalization of the observed tendency of damage to propagate from the outer plies to the inner plies. The implications of the results are also discussed for potential thermal exposures of the composites in the intermediate-temperature regime. The paper is divided into six sections. The material processing parameters and microstructures are presented in Section 2, prior to a detailed description of the experimental procedures in Section 3. Damage mechanisms observed under monotonic loading at room-temperature

Keywords: A. MMCs, B. fracture, C. FEA, D. acoustic emission, E. heat treatment

1 INTRODUCTION

Titanium-matrix composites are currently being con(500-650°C) sidered for intermediate-temperature structural aerospace applications by virtue of their attractive combinations of high strength and stiffness, and mechanical/physical property retention at elevated temperature. 1,2 The interest in potential aerospace applications of these composites has led to a number of focused research efforts designed to improve our understanding of the micromechanisms of fatigue3-I4 *To whom correspondence 292 1537.

should be addressed. Fax: 001 614 915

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are then discussed in Section 4. Damage initiation phenomena are rationalized by careful interpretation of the results obtained from idealized and actual microstructure-based finite element stress analyses in Section 5. A summary of the salient conclusions arising from the current work is then presented in Section 6. The paper highlights the importance of the combined use of micromechanics and materials science principles in the development of relatively simple models for the prediction of composite behavior.

2 MATERIALS

AND MICROSTRUCTURES

The symmetric eight-ply [0/9012, composite that was used in this study was supplied by Textron Specialty Metals, Lowell, MA. It was produced by the foil/fiber/ foil technique via hot pressing at 982°C for 2 h. The resulting composite microstructure is presented in Fig. l(a). This shows the side view of a typical hot pressed composite. The nominal volume fraction of the fibers is approx. 35 ~01%. The carbon-coated Sic fibers have a diameter of approx. 145 Km. The thickness of the carbon coating was approx. 2.6mm after hot isostatic pressing at 982°C for 2 h. A uniform distribution of SIC fibers was observed in the [0/9012, composite lay-up that was used. Large (-500mm) metastable B grains are also observed in the composite, and the typical layered interfacial structureI is revealed clearly at higher magnification. This interfacial layer has a highly complex microstructure that consists primarily of titanium carbides (TIC and Ti$) and some titanium silicides.‘8,19 Two heat treatments were used to control the composite microstructure. The first heat treatment involved annealing for 50 h at 540°C (below the matrix B transus of approx. 800°C in20) before an air cool. This heat treatment resulted in the transformation of the metastable as-received @ structure (Fig. l(a)) to a Widmanstgtten colony microstructure with small acicular a! grains in a matrix of B (Fig. l(b)). The changes in the matrix microstructure occurred without significant coarsening of the interfacial structure, as shown in Fig. 1 and Table 1. Note that the matrix a! grains are the light (white) phase, while the /? phase is the dark (gray) matrix background phase in Fig. l(b). The large grains Table 1. Summary of interfacial dimensions

Condition/heat treatment

Interface thickness (wm) Carbon coating Titanium carbide thickness thickness

As-received 54O”C/lO h/AC 54O’CjSOh/AC 54O”C/lOOh/AC 8lY’C/lO h/AC 8 15OCjSOh/AC 815”C/100h/AC -

2.6 2.6 2.4 2.8 3.0 3.0 2.8 -. ___ ~~~-.~~

2.9 2.6 2.3 3.1 3.2 3.2 2.9 ~~_~

Fig. 1. (a) Effects of annealing on composite microstructure: (a) as-received (included for comparison), (b) 54O”C/50 h/AC, and (c) 8 15Tj50

_~

h/AC (AC = air-cooled).

in Fig. l(a) are the metastable /? grains which eventually transform to the WidmansMtten structure during thermal exposure below the /I transus (Fig. l(b)). The second heat treatment was carried out at 8 15°C. A heat treatment duration of 50 h was employed before

Tensile damage in a symmetric composite

air-cooling to room temperature. This heat treatment was designed to promote significant coarsening of the interfacial microstructure without causing significant alteration of the metastable b matrix microstructure. However, unlike the four-ply unidirectional Ti-153jSiC (SCS-9) composite examined in previous studiesI the coarsening of the interfacial microstructure was limited in the eight-ply [0/9O]z, Ti-153/SiC (SCS-6) composite that was used in this study. Nevertheless, the formation of smaller fi sub-grains was observed in the metastable beta matrix after annealing above the B transus. A typical SEM photomicrograph of the composite annealed at 8 15°C is presented in Fig. l(c).

3 EXPERIMENTAL

PROCEDURES

Triplicate tensile tests were performed at room- and elevated-temperature on smooth 152.4 mm long specimens with rectangular cross sections (1.2 mmx 12.7 mm). Note, however, that the damage sequence in the elevated-temperature tensile tests will not be discussed in detail here because of the instabilities of the composite microstructures at the test temperature of 650°C. The specimens were fabricated by water-jet cutting, and a servohydraulic test machine was employed in the mechanical testing. The first set of tensile specimens were loaded continuously to failure at a strain rate of 5x 10-4 s-1. Strain was measured with a contact extensometer with a gauge length of 25.4mm. Damage under monotonic loading was monitored continuously with an automated acoustic emission unit with two piezoelectric sensors that were located about 100mm apart on either end of the gauge sections. Relevant acoustic emission data such as amplitude, energy, rise time, frequency, counts and duration, were obtained for the three microstructural conditions (Fig. 1 and Table 1). The use of two acoustic emission sensors also made it possible to locate the sources of the acoustic emission signals. I4 The initial acoustic emission data was then analyzed to identify signals that are characteristic of particular failure modes. A second set of tensile tests were conducted on smooth specimens to study the deformation and cracking phenomena associated with damage under monotonic loading. The specimens were loaded in incremental steps of 0.1 a,, to various fractions of the ultimate tensile stresses, a,, determined from the first set of tests. Damage phenomena associated with the different incremental loading steps were then identified by ex situ scanning electron microscopy examination of the sides (of the gauge sections) of the deformed tensile specimens. The changes in tangential modulus associated with each incremental loading step were quantified at the maximum load using a contact extensometer with a gauge length of 25.4mm. In this way, the sequence of damage was identified for the two microstructural conditions that were employed.

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A third set of tensile tests were carried out to further explore the details of matrix yielding and damage phenomena. These involved conducting matrix microhardness measurements, scanning electron microscopy examination, and acoustic emission analyses of specimens that were loaded to regimes where distinct acoustic emission signals were observed during the first set of (continuous) tests. At least 15 microhardness measurements were made using a 25g indentation load and a Vickers indenter. The indentations were made in each ply of the composites after loading and unloading from these distinct acoustic emission regimes. These measurements were made from the matrix regions. The changes in the matrix microhardness and composite cracking phenomena were thus related to the measured acoustic emission data and the composite tensile properties. Fractographic analysis of the fracture surfaces of selected tensile specimens was also carried out using scanning electron microscopy techniques.

4 TENSILE

DEFORMATION

AND FRACTURE

4.1 Tensile properties Average room- and elevated-temperature (65O’C) tensile properties of the as-received and heat treated composites are summarized in Table 2. Characteristic stressstrain plots are also presented in Fig. 2(a) and (b). The stress-strain characteristics of the as-received and 8 15’CjSO h/AC samples were almost identical, consistent with the similar composite microstructures in these two conditions (Fig. 1 and Table 1). The composite strength was also higher after annealing at 540°C 50 h/AC which resulted in a transformed Q + /? Widmanstatten microstructure. Such annealing was associated with lower ductility at elevated temperature. However, the ductilities were similar in the two microstructural conditions at room-temperature. The stress/strain response was linear at room temperature until a critical stress was reached. Average values of this critical stress, cl, are presented in Table 2. Non-linear behavior ensued beyond this critical stress at room temperature. Similarly, non-linear stress/strain response was observed at 650°C in the two microstructural conditions that were examined. The specimens annealed at 815°C were also observed to have lower strength and greater ductility than those annealed at 540°C (Fig. 2(b) and Table 2). Furthermore, the composites annealed at 815°C exhibited a distinct yield point, above which significant strain hardening was observed (Fig. 2(b)). The deformation and cracking phenomena associated with the observed stress/strain behavior at room-temperature are discussed below. 4.2 Damage mechanisms The damage phenomena associated with the different regions of the stress/strain curves at room temperature

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Table 2. Summary of tensile properties at room and high temperatures (25”C, 650°C) Condition/heat

..____ As-received 54O”CjSO h/AC 8 1SC/SO h/AC

treatment

650°C

25°C UTS” Critical _ Critical (MPa) stress 1 stress 2 (MPa) (MPa) .~~__ ~~~~ ~~~ ~____ 856 472 616 1028 345 540 891 377 611

are summarized in Fig. 2(c) and (d). Note that this section will focus on damage at room temperature since the microstructures were unstable at elevated-temperature (650°C). Damage initiation is also described arbitrarily to include all the damage events (sub-grain or slip band formation and interfacial cracking) prior to matrix cracking, while damage propagation/evolution is considered here to involve all the subsequent stages of damage (multiple crack growth, fiber fracture and crack coalescence) after the onset of matrix cracking. Composite microstructure (Fig. l(a-c)) was also found to have an effect on the sequence of damage (Figs 3(aw(b)), although similarities were also observed in the overall damage sequence in the two microstructural conditions that were examined in detail, i.e. the microstructures produced by annealing at 540°C for 50 h or 8 15°C for 50 h.

Total strain W)

~~~_

1.1 I.2 1.1

Critical” stress 1 (MPa)

Critical stress 2 (MPa)

169 266

414

UTS (MPa)

Total strain (%)

568 556

0.79 2.00

Local evidence of plasticity was observed early in the deformation sequence, i.e. prior to bulk yielding across the gauge of the specimen annealed at 540°C. This form of local plasticity, which is referred to subsequently as microplasticity, occurred at very low stresses (approx. 0.1 cU), and is illustrated in Fig. 3(a). The microplasticity manifested itself in the form of slip bands, which were observed to nucleate from the fiber/matrix interface (Fig. 3(a)). Similar slip band initiation mechanisms have been observed in previous studies by Majumdar et af.15,16on 0 and 90” Ti-15-3/SCS-6 composites. Note that the stress/strain behavior of the composite is still linear in this microplasticity regime, as is typically observed in conventional monolithic materials. Beyond the initial regime of microplasticity, crack initiation was observed to occur in the outer 90” plies

cd) Stress

Strain

Strain

Fig. 2. Summary of stress/strain behavior and underlying damage mechanisms: (a) room temperature (25°C) stress/strain behavior, (b) elevated-temperature (650°C) stress/strain behavior: (c) damage mechanisms in specimen annealed at 54O’CjSO h/AC; and (d) damage mechanisms in specimen annealed at 8 I 5”C/50 h/AC (AC = air-cooled).

Tensile damage in a symmetric composite

(Fig. 3(b)). Slip band intersection was also observed to occur in the regions between the 0 and 90” plies (Fig. 3(c)). These slip bands appeared to have been initiated by the localization of strain in the vicinity of interfacial/reaction zone cracks (Fig. 3(d)), as reported by Majumdar et al. 15,16 for 0 and 90” composites. Slip band intersection and further debonding were also observed at 0.5 gU in the 0 and 90” plies (Fig. 3(d)). The debonding was observed to occur at the region between the titanium carbide interface and the Ti-15-3 matrix. Subsequent matrix crack initiation occurred by the extension of interfacial cracks into the matrix at higher stresses (Fig. 3(e)). Transgranular matrix cracks were nucleated by the extension of interfacial cracks into the matrix. This occurred initially in the outer plies (Fig. 3(e)), while slip band activity was still dominant in the inner plies (Fig. 3(f)). Intergranular matrix crack growth was also observed in the boundaries between the large fi grains. The initiation of matrix crack growth (Fig. 3(h)) preceded final fracture via the coalescence of matrix and fiber cracks in the Mode I direction. A similar sequence of damage was observed in the specimen annealed at 8 15°C for 50 h prior to monotonic

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loading (Fig. 4(a-h)). However, unlike the specimens annealed below the beta transus, slip bands were not observed in these specimens. Instead, matrix microplasticity occurred by the formation of a sub-grain structure (Fig. 4(W)) was also found to precede catastrophic failure, which occurred in the Mode I direction during monotonic loading at room temperature. Catastrophic failure occurred by ductile dimpled matrix fracture and cleavage/quasi-cleavage fracture of the SCS-6 fibers in the different microstructural conditions that were examined (Figs 5 and 6) There was also strong evidence of fiber pull-out in all the specimens (Figs 5(a) and 6(a)), and the interfaces between the fibers and the matrix were clearly degraded/damaged during the fiber pull-out process (Figs 5(a) and 6(a)). Interfacial damage was apparent in both the 0 and 90 plies, and a transgranular fracture mode was observed in the MO cross weave that was used to hold to composite together during fabrication. The MO cross weave has been shown to promote premature failure previous studies.14 However, fractographic evidence of premature crack nucleation from the MO cross weave was not observed in this study.

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Fig. 3. Damage mechanisms under monotonic loading in specimen annealed at 54OTj50 h/AC and tested at room temperature (25°C): (a) slip band initiation in the outer 90” ply with debonding at 0.2 0 “r (a = 181 MPa); (b) matrix crack nucleation from interface in outer ply at 0.30~~ (a = 271 MPa); (c) intersection of slip bands in 0” and 90” plies at 0.3aUr (a = 271 MPa); (d) further debonding in 0” and 90” plies at 0.5 ffur (a = 452 MPa); (e) nucleation of matrix crack from interface with debonding in inner ply at 0.6 (YUT(a = 542 MPa); (t) debonding and slip band formation in inner plies at 0.7 0 Ur (a = 633 MPa); (g) matrix crack coalescence in the outer ply at 0.9 (Tur (a = 814 MPa); and (h) evohttion of matrix cracks in inner ply at 0.90 our (a = 814MPa).

4.3 Acoustic emission Prior to a presentation of the acoustic emission data, it is important to identify the different acoustic emission parameters that were used in this study. Most of these are summarized in Fig. 7(a) in which a typical discrete acoustic emission burst signal/wave packet is presented. Note that the energy envelope is the area under the plot of the peak voltage amplitude vs time. Also, the discrete wave packet includes noise signals that can be filtered out to ensure that only acoustic signals due to deformation and microcracking phenomena are analyzed.r4 A 60dB amplitude threshold level was used to ensure

that the acoustic emission signals were above the background noise levels. Note that each discrete energy burst detected by the acoustic emission probes is referred to as a hit or event in subsequent discussion. The number of counts is the number of peaks above the threshold/trigger level, while the area under the amplitude vs time plot

is a measure of the energy of the discrete wave packet. The other acoustic emission parameters such as rise time and duration are defined in Fig. 7(a). Typical acoustic activity data obtained from specimens that were deformed continuously to failure at room temperature are presented in Fig. 7(b-h). The initial acoustic emission signals were detected during the first stage of deformation in the microplastic regime (Figs 24). Note that the time axes in Fig. 7 scale directly with the load axes in Fig. 2 since constant loading rates were employed in the tensile tests. The initial hits shown in Fig. 7(b) were associated with the onset of microplasticity and debonding (Figs 3(a)-4(b)). Microplasticity in the current study was associated with slip band or sub-grain formation (Figs 3 and 4). The number of hits increased to a maximum level at the onset of bulk yielding and matrix cracking. It then decreased during subsequent deformation by matrix

Tensile damage in a symmetric composite

cracking and fiber fracture. Similar correlations have been established in studies on monolithic alloys.*‘~** A typical plot of energy vs time is shown in Fig. 7(c) for material annealed at 540°C for 50 h. The initial energy signals are associated with debonding and the onset of microplasticity (Figs 3 and 4). The sudden jumps in energy during the first phase of damage are attributed to matrix cracking in the outer plies, while the larger jumps in the second phase of damage (Fig. 7(c)) are associated with matrix cracking in the inner plies. The very high energy signals observed just before failure are due to crack coalescence and fiber fracture (Fig. 3). It is interesting to note here that the trends in the energy vs time plots were similar to those in the plots of count rate vs time. Differences between the damage phenomena in the outer and inner plies were also revealed by the plots of cumulative counts vs time which have three distinct slopes, as shown in Fig. 7(d). The relatively low initial slopes are associated with damage in the inner plies, while the moderate slopes in the second phase of damage are correlated with damage in the inner plies. The very high slopes observed during the final damage regime are associated with fiber frac-

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ture and crack coalescence prior to catastrophic failure (Fig. 3). The plots of amplitude vs time are presented in Fig. 7(e) for materials annealed at 540°C for 50 h. Note that the data has been filtered at the 60dB level to eliminate extraneous noise signals. Hence, no variations in the amplitude data are observed below 60dB. The amplitudes of the events increase almost continuously after the onset of debonding and microplasticity (Fig. 3). Events with amplitudes between 70 and 80dB are associated with debonding and microplasticity, while those with amplitudes between 80 and 90dB are attributed with matrix cracking in the outer plies. Cracking in the inner plies resulted in signals with amplitudes between 90 and lOOdB, while fiber fracture resulted in events with amplitudes greater than 100 dB. It is also interesting to note here that most of the acoustic events had frequencies that were between 50 and 3OOkHz (Fig. 7(f)). However, the total range of frequencies that was detected was between 0 and 1000 kHz, although very few events had frequencies greater than 200 kHz (Fig. 7(f)). Finally in this section, it is of interest to compare the acoustic events in the specimens annealed at 540 and

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Fig. 4. Damage mechanisms under monotonic loading in specimen annealed at 8 1SC/SO h/AC and tested at room temperature (25°C): (a) debonding in inner and outer plies at 0.1 (TU-r (a = 104 MPa), (b) initiation of sub-grains at 0.2~~“~ (a = 208 MPa), (c) matrix crack initiation at 0.3 o “r (a = 3 13 MPa); (d) matrix cracking in 90” ply with further debonding at 0.4 (TUT(a = 417 MPa); (e) further sub-grain formation at 0.5 0 or (c = 521 MPa), (f) grain boundary cracking and further debonding at 0.6 ct,T (o. = 626 MPa), (g) fiber fracture in 90” plies at 0.9 (T“r (a = 939 MPa), and (h) fiber fracture in 0” plies at 0.9 (TUT((T= 939 MPa).

815°C. Typical plots of energy vs counts are presented for these two heat treatment conditions in Fig. 7(g) and

(h). Higher energy events were generally observed in the materials that was annealed at 540°C for 50 h. This material had a Widmanstatten structure, as shown in Fig. l(b). The other acoustic emission signals were also generally higher in the composites with the more brittle Widmanstatten structure (Tables 3(a) and (b)). However, the trends in the most of the acoustic emission data were generally similar in the two microstructural conditions that were examined in detail, i.e. matrix crack initiation was typically associated with low or moderate acoustic emission activities with shorter durations. An increase in acoustic emission activity was then observed during the matrix crack growth phase. This was followed by a sharp drop in rise time and a significant increase in acoustic activity upon initiation of

fiber cracking. The rise times and all the other acoustic emission signals were very high in the fiber fracture/ crack coalescence regime that preceded catastrophic failure. 4.4 Modulus and hardness The stress/strain behavior of 0 and 90” Ti-153/SCS-6 composites has been studied in detail by Majumdar et LZ~.‘~,‘~ They show two-stage and three-stage behavior in 0 and 90” composites, respectively. They also attribute the kinks in their stress/strain plots to the effects of slip bands/matrix yielding phenomena. It is therefore not surprising that [O/9012, composites show multi-stage stress/strain behavior due to matrix yielding and damage phenomena in the individual plies (Figs 2-6, and Table 2). The trends in the apparent/tangential composite moduli are shown in Fig. 8 for materials

Tensile damage in a symmetric composite

OJ)

923

@)

Fig. 5. Typical fracture surface morphologies of specimens annealed at 54O’CjSO h/AC and tested at room temperature (25°C): (a) evidence of fiber pull-out; and (b) ductile dimpled matrix fracture.

Fig. 6. Typical fracture surface morphologies of specimens annealed at 815’CjSO h/AC and tested at room temperature (25°C): (a) evidence of fiber pull-out; and (b) ductile dimpled matrix fracture.

annealed at 540 and 815°C. These show that composite modulus generally increases with increasing applied stress. This is contrary to the behavior that is generally observed in cracked materials which tend to exhibit a decrease in modulus with increasing applied stress. Some discussion on the possible causes of the apparent increase in modulus is therefore required here. The apparent increase in global composite modulus can be explained by considering the possible effects of matrix hardening that was observed during incremental loading to failure at room temperature. Plots of matrix hardness are presented in Fig. 9(ad) which show typical microhardness variations established as a function of ply number (Fig. 9(a,b)) and the fraction of the failure stress (Fig. 9(c,d)). A symmetric variation is observed across the thickness of the samples, and the outer plies are clearly harder than the inner plies (Fig. 9(a,b)). The annealed specimens are also much harder than the as-received samples (Fig. 9(a,b)), and

the matrix hardness increases monotonically with increasing applied stress (Fig. 9(c,d)). In many metal-matrix composites, the rate at which the composite modulus decreases due to damage is often greater than the rate at which hardening occurs in the metal matrix. However, it is possible in some cases to for the composite modulus to increase with increasing stress if the rate of matrix hardening is greater than the rate at which the composite modulus decreases as a result of cracking phenomena. The observed increase in composite modulus with increasing stress (Fig. 8) may therefore be rationalized within the framework by considering the combined effects of hardening and cracking. Since an increase in matrix hardness should promote an increase in the composite modulus, while cracking phenomena will tend to decrease the composite modulus, the current results suggest that the effects of matrix hardening phenomena dominate the apparent global composite moduli in the composites examined in this

924

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0. Soboyejo et al.

ENVEMPE THRESHOLD

-

0

14

20

42

56

70

Time (set) (b)

0

14

28

42 Time (SW)

:__..;__ : ..._.i.... .._..i .1’ i;.. 70

56

0

14

28

42

56

6ca

800

Tune (SC) (d)

(C)

1

.

.

;

j

. . .

f . . . . i

. .

; . . . . .

.

i . . . . . :

0

14

42

28 Time (SK)

0

70

200

4co

loo0

IS(K)- -/

-

x

50

64

78

1200-

:

92

Amplitude

LdBJ

(a

Amplitude

[dB]

(h)

Fig. 7. Summary of acoustic emission activities of specimen annealed at 54O“CjSO h/AC and tested at room temperature (25°C): schematic representation of acoustic emission parameters; (b) acoustic emission hits vs time; (c) energy vs time; (d) cumulative counts vs time; (e) amplitude vs time; (f) hits vs frequency; (g) energy vs amplitude; (h) energy vs amplitude in specimen annealed at 8 1ST/50 h/AC (included for comparison).

study. Global measures of composite modulus therefore do not provide a simple measure of composite damage, as postulated in conventional damage mechanics theories. Different models are therefore needed to assess

the relative contributions of matrix hardening and cracking phenomena to the apparent composite moduli. Models are also needed to assess the effects of multiple configurations of randomly oriented defects.

Tensile damage in a symmetric composite

E (81YW”nal

400

600

in the current study where crack initiation was typically observed at the fiber/matrix interface (Figs 3 and 4). Note that the local variations in the interfacial stresses were estimated for closely spaced fibers with high stress concentrations. The stress estimates were obtained using idealized and actual microstructure-based finite element models presented below. Prior to a presentation of these models, it is important to note here that more rigorous three dimensional analyses are needed to obtain very accurate estimates of the stress distributions within the laminae. However, it is very difficult to conduct such three dimensional analyses on systems with detailed microstructural information. Simplified, but effective two dimensional plane strain deformation configurations, were therefore employed in the finite element modeling of the composite structures. The potentially important effects of residual stress and plasticity phenomena were also neglected in the analyses. Nevertheless, the accuracy of the stress estimates obtained from the current analyses is thought to be sufficient for the qualitative characterization of average and local stresses in the individual plies.

L.) [GPal

800

1000

1200

AppliedStress[MPa] Fig. 8.

Summary of measured composite moduli.

5 MICROMECHANICAL

MODELING

5.1 Background Since it was experimentally observed that damage in the outer 0” plies always occurred before damage in the inner plies (Figs 3 and 4), attempts were made to calculate the stresses in the individual plies in this section. Inherent to this approach, was the presumption that crack initiation would occur in regions with the highest local stress concentrations. This was generally the case

5.2 Idealized microstructure-based model The idealized 8-layer symmetric [O/9012,cross-ply laminate plate (Fig. 10(a)) was subjected to uniform in-plane tension (90” direction) in the analysis, i.e. the deformation of the entire laminate was modeled using a plane strain (b) GQ

I -

ARMatixHv

-

540~504.0’1

- -

Ply

YO-500.3TF SSJ-OBTF

-

1

I -

ARMntnx

-

5155O-O.OTF

-

815504.35TF 815-50-0.65TF

hv

Number

..! . . . . . . . . . . . . . . .I . . . . . . . . . . . . . . . . . . . .

A

260

925

i -

-

i . . . .. . . . .

j

Ply Hiv(540.501 cue, Ply H<540-50] Inmr

I

Effects of monotonic loading on matrix micro-hardness: (a) hardness vs ply number (54O”C/SO h/AC sample); (b) hardness vs ply number (815Y750 h/AC sample); (c) hardness vs fraction of failure stress (54OCj50 h/AC sample); and (d) hardness vs fraction of failure stress (8 1SC/SO h/AC sample).

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Table 3(a). Corellation of acoustic emission activities and visual observations in O”/W composites under monotonic loading Damage

mechanism

Acoustic Rise time (P s)

Counts

Energy

(J)

emission Duration (CLs)

activity

~~~ ~~~~~

Amplitude (dW

Frequency (KHz)

Applied

stress

(M(TPa)

Matrix crack initiation I 54O”C/50 h/AC 2 8 15”C,i50 h/AC

43 51

133 152

498 690

1088 1704

99 100

12 8

546 607

Matrix crack growth 1 54OTjSO h/AC 2 8 1ST/SO h/AC

105 II

357 219

1348 750

3216 3044

100 99

II 7

580 676

Fiber cracking 1 54OTj50 h/AC 2 815”CiSO h/AC

10 2

543 445

1813 I568

4219 5028

100 100

12 8

1015 799

Crack coalescence 1 54O’CjSO h/AC 2 815”C/50 h/AC

92

967

233

232

3534 819

6957 3073

100 98

13 7

1021 876

Table 3(b). Range of acoustic activities in 540”C/50h/AC 8lS’C/SO h/AC specimens AE parameter Rise time (ps) Number of counts Energy (J) Duration (11s) Amplitude (dB)

Frequency (kHz)

54OTjSO h/AC l&l05 95-967 374-3534 1005S6957

97-100 8-13

and

815”C/50h,‘AC l-233 19445 137-1568 1119~5152”

97-m1ooh I-1 I

“Filtered 2 1000 (ws). hFiltered 296 (dB). model of deformation in the y-z plane. This is illustrated schematically in Fig. 10(b). In the plane strain model, the 0” plies are constrained to deform along a path that results in the retention of the inclusion geometry. The 90” plies are homogenized, i.e. each of the 90” plies is replaced by homogeneous transversely isotropic layers with effective elastic properties that are determined using composite theories which will be described later on in this section. The idealized laminate model assumes that there are clear boundaries between neighboring plies (Fig. lO(a,b)). A uniform distribution of fibers in the 0” plies is also assumed, i.e. the fibers are assumed to be perfectly spaced and periodic in their arrangement. By considering the geometric and loading symmetries in Fig. 10(b), a finite element mesh was generated for the idealized model. This is shown in Fig. 1 l(a). In this mesh, special fiber-embedded polygonal elements were employed in 0” plies. These special elements were created using the hybrid method proposed originally by Zhang and Katsube.23124 Quadratic displacement interpolation was used on the outer boundary of the hybrid elements to satisfy the conditions required for inter-element displacement compatibility. Traditional displacement-based S-node iso-

parametric anisotropic elements were used in the 90” plies, along with effective material properties which were calculated using an approach proposed by Christensen.25 Appropriate boundary displacement constraints were enforced at the nodal points along the lines of symmetry. The dimensions of the elements that were used in the model are given in Fig. 11(a). The volume fraction of fibers of the laminate is about 39%. The radius of fiber is 2.65x 10-j. The Young’s moduli and Poisson’s ratios of the fiber and matrix are, respectively: Ef = 428 GPa (62.00 Msi), Em = 85 GPa (12.30 Msi), vf = 0.30, and V 0.32, where the subscripts ‘f’ and ‘m’ denote the fi;eT and matrix, respectively. The effective moduli of the equivalent transversely isotropic material in the 90” plies were evaluated using a method proposed by Christensen.25 This method provides explicit expressions for the evaluation of five independent effective stiffness coefficients for transversely isotropic material. In the derivation of these expressions, the composite cylinders model introduced by Hashin and Rosen 26 is used. The resulting expressions are given in terms of the fiber volume fraction and the material properties of the fiber and matrix. In the present case, only four of the five stiffness moduli are required due to the assumption of plane strain conditions. The stress/strain relationships in the 90” plies are thus given by:

(1)

where CII, CIZ, C22, and CM are the four independent effective material stiffness coefficients. These four effective stiffness coefficients in the present case are: Cl1 = 270GPa (39.1 Msi), Cl2 = 85GPa (12.3 Msi), C22 = 188GPa (27.3 Msi), and Cu = 55GPa (7.9 Msi).

Tensile damage in a symmetric composite

927

Ply Number

A=

(b)

Fig. 10. (a) Idealized microstructure-based model: schematic of [O/9032, composite and load orientation; and (b) schematic of plane

strain deformation model.

It is important to note here that the above homogenization schemes provide only an approximate description of stiffness moduli of the O/90 composite that was examined in this study. Nevertheless, the methods that were used to compute the effective moduli25,26provide useful estimates of the stiffness matrices of transversely isotropic systems. The idealization of the 90” layers has also been introduced to simplify the analysis of the O/90 composite. One effect of this idealization is the increase in the apparent volume fraction of fibers from 35 to 55 ~01%. This is needed to ensure that

the computed effective composite moduli are equal to the actual composite moduli in the O/90 composite system. The radial, hoop, maximum principal, and maximum shear stresses along the fiber-matrix interfaces in the outer and inner 0” plies are shown in Fig. 1l(b,c). The values of the stresses are normalized with respect to the applied stress oo. Also shown in these figures, are the principal stress directions at the interfaces. The symbols ‘A’ and ‘B’ denote the fibers in the outer and inner 0” plies, as shown in Fig. 1l(a).

i

15

-

radial $tlass hoop stress prtnclpal stress maximum shear Stwss

maximum

I -1 5

0

50

130

150 200 Angle (degree)

(a)

250

300

350

(b)

D B ii 1.2.

*~~;?:=“h

*

J

p

1

k? .l m

0

: outer 0 degree ply

0

:

x

Inner0 degree pty 90 degree ply (layer 2)

1

I 150

200

Angie (degree) (C)

250

x0

350

1

0.5

15

(d)

Fig. 11. Idealized finite element model: (a) finite element mesh for idealized plane strain analysis; (b) normalized fiber/matrix

interface

2

Y coordtinate (mm)

in the outer 0” ply (fiber A); (c) normalized stresses along the fiber/matrix interface B); and (d) normalized element average stresses in the outer and inner 0” plies.

stresses along the in the inner 0” ply (fiber

928

W. 0. Sohoyejo et al.

There is strong evidence of stress concentrations at the fiber/matrix interfaces. The normalized peak value of the maximum principal stress of the outer 0” ply (fiber A) is 1.24 and that of the inner 0” ply (fiber B) is 1.16. Both peak values occur at the locations with the angles of 0 and 180”, respectively, consistent with the initial observations of debonding (Fig. 11(b,c)). The normalized average stresses oYY/ac,are presented in Fig. 1l(d) for each layer. It is clear that the outer 0” ply has larger layer average stress than the inner 0” ply in regions that is distant from the centers of the plies. These regions represent the load transfer regimes where the loads supported by the outer 0” ply is gradually transferred to the inner 90” plies. The average element stresses along region M-M are presented in Table 4 for each layer before and after failure of the 0” plies. As shown in Table 4, three different cases are considered in the analysis: (a) the configuration before fiber/matrix debonding; (b) the configuration after debonding in layer 1; and (c) the configuration after debonding in layers 1 and 3. The primary objective in conducting these analyses. was to examine the load transfer tendency from the 0” plies during monotonic loading to failure. In case (a). the analysis was performed based on an assumption of perfect bonding, and the finite element mesh shown in Fig. 1l(a) was used. In cases (b) and (c), the failure of 0” plies was modeled by assuming total debonding of fiber/ matrix interfaces. Thus, the fibers shown in Fig. 1I(a) are replaced by holes in the 0” plies. The assumed fiber/matrix debonding actually occurs in the present study (Figs 3-6), since the fiber/matrix interface is relatively weak. 27.28However, if the interface is strongly bonded, the initiated microcracks at the interface may not propagate along the interface, but into the matrix region. Under such circumstances, the simple failure model in cases (b) and (c) will not be a good assumption. The assumption of a weak fiber/matrix interface is clearly valid in the present case, as shown in Figs 3 and 4, in which the damage sequence at room temperature are presented. Note that weak interfacial strength data (34200MPa) have been reported in the literature for the Ti-I 5-3/SCS-6 composites.27,28 Before debonding, the 0” and 90” plies support about 42 and 58% of the total load, respectively (Table 4). Of the load that is supported by the 0“ plies, the outer 0” plies support about 4% more load than the inner 0” plies. After the outer 0” plies have debonded completely, the capacity of these plies to support load is substantially reduced. Table 4 indicates that the average stresses in the outer 0“ plies decreases by about 62% after fiber/matrix debonding. Of the load that is released by the failed outer 0” plies, about 60% is transferred to the neighboring layer, i.e. layer 2. Also, 23 and 17% of the released loads are transferred to layers 3 and 4, respectively. After the inner 0” plies have debonded totally, the 90” plies support about 83% of the total applied load. The results shown in Fig. 1l(b,c) and Table 4 provide qualitative explanations of some of the observed

damage initiation phenomena. It is important to note here that the reference stress, 00, corresponds to the remote elastic stress that is applied to the boundaries of the layered composites (Figs 1l(a) and 12(a)). Since the actual stress concentrations (due to the effects of composite geometry) are independent of the actual value of a~, the reference stress was prescribed as 1 in the analysis. The normalized stresses shown in Fig. 11 and Fig. 12 are therefore normalized by a reference stress of unity. The relatively large stress concentrations at the fiber! matrix interface are one the causes of the initiation of local failure in the 0” plies. The maximum stress concentration factors are approximately 1.5 for both the outer and inner 0” plies with respect to their respective layer average stresses. The observation of damage in the outer 0” plies before damage in the inner plies can also be explained by the calculated stresses which show that the outer 0” plies support more load than the inner plies in the load transfer region. The absolute stress concentration factor in the outer 0” ply is thus larger than that in the inner 0” ply (Fig. 1l(d) and Table 4). The higher levels of stress in the 90” plies are consistent with the initial occurrence of debonding in the 90” plies before debonding in the 0” plies (Figs 3 and 4). The stresses in the outer 90” plies were also generally higher than those in the inner 90” plies. The stress estimates in the 90” plies therefore appear to be qualitatively accurate, in spite of the homogenization that was used to simplify the stress analyses in these plies (Fig. 11(a)). 5.3 Actual microstructure-based model The results of a microstructure-based finite element analysis are presented in this section. This analysis, which includes the effects of actual fiber distributions, is based on the actual finite model of a representative cross-section photomicrograph of an as-received composite laminate. As shown in Fig. 12(a), the matrix regions in the 0” and 90” plies are connected to each other in the finite element model. Three types of elements were used in the finite element mesh. Hybrid inclusion-embedded elements were used in the 0” plies.23,24 Anisotropic elements were used in the 90” plies. The height of these anisotropic elements was equal to the fiber diameter in the model. Traditional 8node isoparametric isotropic elements were used in the Table 4. Normalized average stress in each layer at section MM before and after 0” plies debonding

Layer

1 2 3 4

Average stress cry&r,, Before debonding

After layer 1 debonding

After layer I and 3 debonding

0.862 I.184 0.822 1.131

0.324 I.504 0.949 1.223

0.389 I .833 0.292 I.486

Tensile damage in a symmetric composite

2.0

929

mm (a)

0 -0.5 z

. .

-1 I

9 : maximum 0

pwwpal

: maxmum shear

stress stress 150 Angle

200 (degree)

(b)

z

2

z -0.5 z

A 9 = 0

-1 .I.5

0

I

radial

stress stress maximum principal stress maxmum shear stress

: hoop

50

100

150 Angle

principal direction i 200 (degree)

250

300

350

Cc)

Fig. 12. Actual microstructure-based finite element model: (a) finite element mesh; (b) normalized stresses along the fiber/matrix interface in the outer 0” ply (fiber C), and (c) normalized stresses along the fiber/matrix interface in the inner 0” ply (fiber D).

narrow strip of the matrix region between the two anisotropic layers in the middle of the model. The geometric dimensions used in the actual model are shown in Fig. 12(a). The material properties of the fiber and matrix, the volume fraction of fibers in the laminate, and the fiber radius are the same as those in the previous idealized model. However, the volume fraction of fibers in the homogenized anisotropic elements in the 90” plies is changed to 55%, since the height of these elements is reduced to the diameter of the fiber. The four effective stiffness moduli of the elements are correspondingly re-evaluated to be: Ctr = 338 GPa (49-O Msi), Ct2 = 105GPa (15.2 Msi), C22 = 236 GPa (34.2 Msi), and CM = 70GPa (10.1 Msi). The normalized radial, hoop, maximum principal and maximum shear stresses along the fiber/matrix inter-

faces in the outer (fiber C) and inner (fiber D) 0” plies are shown in Fig. 12(b) and (c). The peak values of the maximum principal stress occur at the location of 180” for both fibers, and their values are l-55 and 1.19, respectively. Once again, it is apparent that the outer 0 plies are subject to more severe stress concentration than the inner plies. Moreover, the peak stress value estimated in the outer 0” plies is about 20% higher than that obtained from the idealized model. This implies that the effect of actual distribution of fibers on local stress concentration factors is not negligible in the crossply laminate that was examined in this study. Finally, it is important to note that the possible effects of residual stress distributions have not been included in the current study. These distributions are difficult to compute in the absence of proprietary information on

930

W. 0. Soboyejo et al.

plies. The early occurrence of debonding and microplasticity are also attributed to the high levels of stress concentration at the fiber/matrix interface. The actual microstructure-based model predicts higher level of stress than the idealized finite element model when the fibers are more closely spaced in a realistic random fiber array.

the temperature profiles that were used in the processing of the composites. Nevertheless, the current analysis provides some useful insights into the effects of composite geometry on the stress distributions in the O/90 composite that was examined in this study. Further work is clearly needed to improve our understanding of the effects of residual stresses and actual fiber distributions on composite damage phenomena.

6 CONCLUSIONS

1. The initiation of room-temperature tensile damage in Ti-153/SiC (SCS-6) composites occurs by debonding and microplasticity at stress levels between 0.1 and 0.2 of the ultimate tensile stress. Microplasticity in the specimens annealed at 540°C is associated with slip band formation, while microplasticity in the specimens annealed at 815°C is associated with sub-grain formation. Microplasticity and debonding do not promote nonlinear stress/strain behavior. Subsequent damage, which occurs by matrix and fiber cracking, is associated with non-linear stress/strain behavior. 2. Correlations have been established between the different stages of damage and the levels of acoustic emission signals obtained observed deformation at room temperature. Interfacial cracking and matrix yielding phenomena are associated with acoustic activities with amplitudes less than 96dB and durations less than 1000~s. Matrix crack initiation is associated with low/moderate acoustic emission activity, while fiber cracking and crack coalescence are both associated with high acoustic activity. However, a distinct drop in rise time is observed at the onset of fiber cracking. Microstructural characteristics do not appear to affect the above correlations. 3. The changes in the apparent modulus that are associated with increasing applied stress can be rationalized by considering the combined effects of matrix plasticity and cracking phenomena. The composite modulus will therefore apparent increase when matrix plasticity dominates, or decrease when cracking phenomena dominate. The matrix hardness in the Ti-153/SiC(SCS-6) composite also increases with applied stress, and this increase in hardness appears to be largely responsible for the increase in the apparent modulus that is generally observed during monotonic loading at room temperature. 4. Idealized and actual microstructure-based plane strain finite element models have been developed for the estimation of stresses in the different plies of the composites. The magnitudes of the estimated stresses can be used to explain the gradual progression of damage from the outer to the inner

ACKNOWLEDGEMENTS The authors are grateful to Mr David Harmon of McDonnell Douglas for supplying the composite material that was used in this study. The research was supported by a grant from The National Science Foundation (Grant No. MSS 9309520) with Dr Oscar Dillon, Jr, and Dr William Spitzig as Program Monitors.

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Research