Sung R. Choi

Ohio Aerospace Institute, Brook Park, Ohio

Narottam P. Bansal and John P. Gyekenyesi

Glenn Research Center, Cleveland, Ohio

Dependency of Shear Strength on Test Rate

in SiC/BSAS Ceramic Matrix Composite

at Elevated Temperature

NASA/TM—2003-212182

April 2003
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Sung R. Choi

Ohio Aerospace Institute, Brook Park, Ohio

Narottam P. Bansal and John P. Gyekenyesi

Glenn Research Center, Cleveland, Ohio

Dependency of Shear Strength on Test Rate

in SiC/BSAS Ceramic Matrix Composite

at Elevated Temperature

NASA/TM—2003-212182

April 2003

National Aeronautics and

Space Administration

Glenn Research Center
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NASA/TM—2003-212182 1 

Dependency of Shear Strength on Test Rate in SiC/BSAS Ceramic Matrix 

Composite at Elevated Temperature

Sung R. Choi 

Ohio Aerospace Institute 

Brook Park, Ohio 44142 

 Narottam P. Bansal and John P. Gyekenyesi 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

ABSTRACT

 Both interlaminar and in-plane shear strengths of a unidirectional Hi-Nicalon fiber-reinforced 

barium strontium aluminosilicate (SiC/BSAS) composite were determined at 1100 °C in air as a function 

of test rate using double notch shear test specimens. The composite exhibited a significant effect of test 

rate on shear strength, regardless of orientation which was either in interlaminar or in in-plane direction, 

resulting in an appreciable shear-strength degradation of about 50 percent as test rate decreased from 

3.3×10–1 mm/s to 3.3×10–5 mm/s. The rate dependency of composite’s shear strength was very similar to 

that of ultimate tensile strength at 1100 °C observed in a similar composite (2-D SiC/BSAS) in which 

tensile strength decreased by about 60 percent when test rate varied from the highest (5 MPa/s) to the 

lowest (0.005 MPa/s). A phenomenological, power-law slow crack growth formulation was proposed and 

formulated to account for the rate dependency of shear strength of the composite. 

INTRODUCTION

 The successful development and design of continuous fiber-reinforced ceramic matrix composites 

(CFCCs) are dependent on understanding of their basic properties such as deformation, fracture and 

delayed failure (fatigue, slow crack growth, or damage accumulation) behavior. Particularly, accurate 

evaluation of delayed failure behavior under specified loading/environment conditions is prerequisite to 

ensure accurate life prediction of CFCC structural components at elevated temperatures. 

 Although CFCCs have shown improved resistance to fracture and increased damage tolerant as 

compared to monolithic counterparts, inherent material/processing defects or cracks in the matrix-rich 

interlaminar regions can still cause delamination by the presence of interlaminar normal or shear stress, 

leading to loss of stiffness or in some cases possible structural failure. Strength behavior of CFCCs in 

shear has been characterized in view of their unique interfacial architectures and its importance in 

structural applications [1–5]. Due to inherent nature of ceramic matrix composites, it would be highly 

feasible that interlaminar defects or cracks are susceptible to slow crack growth or damage accumulation 

in certain environments (mostly air) particularly at elevated temperatures, resulting in strength 

degradation or time-dependent failure. Although slow crack growth is one of important life-limiting 

phenomena, little studies have been done on this subject of slow crack growth of CFCCs under shear at 

elevated temperatures. 

 In the previous studies [6,7], ultimate tensile strength of several different CFCCs was determined as a 

function of test rates at elevated temperatures in air ranging from 1100 to 1200 °C. It has shown that the 

ultimate tensile strength of those CFCCs depended significantly on test rate: the greater strength at the 

greater test rate, and vice versa. The leading mechanism has been understood as slow crack growth or 

damage accumulation, evidenced additionally by the results of both stress rupture and preloading testing. 
NASA/TM—2003-212182 2 

This has suggested that life prediction parameters of CFCCs could be estimated at least for a short range 

of lifetimes by a relationship between ultimate tensile strength and test rate, termed “dynamic fatigue” 

that has been used in many monolithic brittle materials including glasses, glass-ceramics, and advanced 

ceramics. The current work, extending the previous studies to include shear, was to determine the rate 

dependency of shear strength at 1100 °C in air using a unidirectional Hi-Nicalon fiber-reinforced 

barium strontium aluminosilicate (SiC/BSAS) composite. Both interlaminar and in-plane shear strengths 

of the SiC/BSAS composite were determined with double-notch shear test specimens as a function of test 

rate ranging from 3.3×10–5 to 3.3×10–1 mm/s. The effect of test rate on shear strength of the 1-D 

SiC/BSAS composite was compared with that on tensile strength of a similar 2-D SiC/BSAS composite. 

A phenomenological, slow crack growth model was proposed to describe the rate-dependency effect of 

shear strength of the composite. 

EXPERIMENTAL PROCEDURES

 The processing of SiC/BSAS composite can be found elsewhere [5,8]. Hi-Nicalon fibers with an 

0average diameter of 14 μm were used as reinforcement. The fiber surfaces were coated with 0.4 μm BN 

followed by 0.1 μm SiC by chemical vapor deposition. The BN interfacial layer acts as a weak, crack 

deflection phase, while the SiC overcoat acts as a barrier to diffusion of boron from BN into the oxide 

matrix and also prevents diffusion of matrix elements into the fiber [8]. The precursor to the celsian 

matrix of 0.75BaO-0.25SrO-Al2O3-2SiO2 (BSAS) was made by solid-state reaction. The precursor 

powder consisted of mainly SiO2 and BaAl2O4 with small amounts of Ba2SiO4, α-Al2O3, and 

Ba2Sr2Al2O7. This powder was made into a slurry with an organic solvent with various additives. Tows 

of BN/SiC-coated fibers were impregnated with the matrix precursor by passing them through the slurry. 

The resulting prepreg tape was dried and cut into pieces. Unidirectional fiber-reinforced composite was 

prepared by tape lay up (20 plies) followed by warm pressing at 150 °C to form a “green” composite. 

Finally, dense composites were obtained by hot pressing under vacuum in a graphite die. X-ray 

diffraction showed that the precursor was fully converted into the desired monoclinic celsian phase 

through solid-state reaction. The fully dense composite laminate was about 4.2 mm thick and had a fiber 

volume fraction of about 0.42. Typical examples of microstructure of the SiC/BSAS composite in planes 

parallel and perpendicular to the fiber direction are shown in Figure 1. 

 The double-notch shear (DNS) test specimens measuring 4 mm (width) × 4 mm (depth) × 20 mm 

(length), as shown in Figure 2, was machined from the composite laminate. The DNS test specimens 

were used previously for the determination of interlaminar and in-plane shear strength of the composite at 

both ambient and elevated temperatures [5]. The square cross-section was intentionally made to make a 

one-to-one comparison between the interlaminar and in-plane shear strength at a given test condition 

without the interference of possible size effect. Dimensions of DNS test specimens used in this work 

were different from those recommended by ASTM C 1425 [4] since the test specimens in the standard has 

a rectangular cross section and would not be appropriate to determine both interlaminar and in-plane 

shear strength. Two notches were 0.3 mm wide, 5 mm away from each other, and situated in equal 

distance from both ends. The two notches were extended exactly to the middle of each specimen so that 

failure occurs on the plane between the notch tips. Detailed descriptions and stress analysis of the DNS 

test specimens can be found elsewhere [5]. Monotonic shear testing was conducted in ambient air at 

1100 °C for the SiC/BSAS DNS specimens using an electromechanical test frame (Model 8562, Instron, 

Canton, MA). A simpler test-fixture configuration consisting of SiC upper and lower fixtures, as shown 

in Figure 3, was used due to the merits of test specimen’s configuration as well as its tight machining 

tolerances, with which test specimen could stand alone - a free-standing feature. A total of five test rates 

ranging from 3.3×10–5 to 3.3×10–1 mm/s were used in displacement control in both interlaminar and inplane shear testing. Typically, three specimens were used at each test rate for a given material direction. 

Each test specimen was held for about 20 min at the test temperature for a thermal equilibration prior to 
NASA/TM—2003-212182 3 

1

3

2

 (a) Plane (‘1-2’) normal to fiber direction 

(b) Plane (‘3-2’) parallel to fiber direction 

(c) Plane (‘3-1’) parallel to fiber direction 

Figure 1. Microstructure of unidirectional (1-D) SiC/BSAS ceramic matrix composite: (a) normal to fiber 

direction, (b) and (c) parallel to fiber direction. The insert defines material’s axes and planes. 
NASA/TM—2003-212182 4 

 (a) Interlaminar (‘2-3’ top plane) (b) In-plane (‘1-3’ top plane) 

 (c) 

Figure 2. Double notch shear (DNS) test specimens [5] used this work showing overall views of 

interlaminar (a), in-plane (b) test specimens, and notch detail (c). The material’s axes and planes are 

defined in the insert of Figure 1. 

 

Figure 3. A schematic showing test fixture and test specimen used in this work. 

Upper fixture

Test specimen

(self standing)

Lower fixture

Thermocouple

2

3

1

3
NASA/TM—2003-212182 5 

testing. The shear fracture stress, average shear stress at failure, was calculated using the following 

relation 

n

f

f WL

P

τ = (1) 

where τf is the shear strength, Pf is the fracture load, W and Ln are the specimen width and the distance 

between the two notches, respectively. A limited fractographic analysis was performed in an attempt to 

help understand mechanisms associated with shear failure. 

EXPERIMENTAL RESULTS

 The results of monotonic shear strength testing for the SiC/BSAS composite are presented in Figure 4, 

where shear strength was plotted as a function of applied test rate for both interlaminar and in-plane 

directions using a log-log scale. Each solid line in the figure represents a best-fit regression line based on 

the log (shear strength) versus log (applied test rate) relation (The reason for using the log-log relation 

will be described in the ‘Discussion’ section.). The decrease in shear strength with decreasing test rate, 

termed a susceptibility to slow crack growth or damage accumulation (or delayed failure), was significant 

for either interlaminar or in-plane direction. The shear strength degradation was about 48 and 59 percent, 

respectively, for interlaminar and in-plane directions when test rate decreased from the highest 

(3.3×10–1 mm/s) to the lowest (3.3×10–5 mm/s). For a given test rate, in-plane shear strength was about 

70 percent greater than interlaminar shear strength, an important factor when a conservative design is 

considered and sought. The interlaminar plane was relatively easier in cleavage in either shear or tension 

than the in-plane counterpart at elevated temperature, attributed to the nature of the composite’s tape layup (laminated) architecture and of the weakened ‘glassy’ (aluminosilicate) matrix due to high 

temperature. Note that except for the magnitude of shear strength, the trend in strength degradation with 

respect to test rate was basically the same for both interlaminar and in-plane directions. The shearstrength degradation was analogous to the trend of tensile strength with respect to test rate that has been 

observed in various CFCCs including SiC/MAS, SiC/CAS, SiC/BSAS, C/SiC, and SiC/SiC composites 

[6,7]. These CFCCs have exhibited significant degradation of tensile strength with decreasing test rates, 

with their degree of degradation being dependent on material and test temperature. 

 Figure 5 shows typical load-versus-time curves determined at two different test rates for both 

interlaminar and in-plane test specimens. At a fast test rate of 3.3×10–2 mm/s, an initial settling stage was 

followed by a linear region until the maximum load reached, beyond which a typical composite failure 

mode was followed. At a much slower test rate of 3.3×10–4 mm/s, the applied load increased linearly up 

to the peak load and dwelled at the peak a little while followed by a sudden drop. The dwelling of the 

peak load was clearly an indication of a phenomenon associated with slow crack growth. At faster test 

rates, a test specimen was subjected to shorter test time so that a crack had no enough time to grow, 

resulting in insignificant slow crack growth or little strength degradation. By contrast, at a slower test 

rate, a crack was subjected to longer time so that a crack had enough time for slow crack growth, thereby 

yielding significant strength degradation. Creep was insignificant even at lower test rates, as can be seen 

from the linear behavior of load versus time curves. The insignificant creep in shear for the 1-D 

SiC/BSAS composite was in contrast to the somewhat significant creep in tension at lower test rates by a 

similar 2-D SiC/BSAS composite tested at 1100 °C in air [7]. 

 Figure 6 shows fracture surfaces of both interlaminar and in-plane specimens tested at high 

(3.3×10–1 mm/s) and low (3.3×10–5 mm/s) test rates. The composite, regardless of material direction, did 

not exhibit any overall, noticeable signs of slow crack growth from their fracture surfaces. Also, the 

fracture surfaces showed little difference in the mode and feature of fracture between high and low test 

rates. However, fracture surfaces were generally smoother in interlamilar than in in-plane specimens. 
NASA/TM—2003-212182 6 

TEST RATE, [mm/s]

10-6 10-5 10-4 10-3 10-2 10-1 100 101

SHEAR STRENGTH, 

τf [MPa]

7

8

9

20

30

40

50

60

70

80

90

10

100

Hi-Nic/BSAS

x&

Interlaminar

In-plane

Unal [ref.5]

(DNS/1100oC)

Figure 4. Shear strength as a function of test rate in both interlaminar and in-plane directions for 

SiC/BSAS composite tested at 1100 °C in air. The arrow in the ‘test-rate’ axis indicates the test rate used 

by Ünal and Bansal [5]. 

0

200

400

600

800

1000

01234567

Time (s)

Load (N)

Interlaminar

3.3x10 In-Plane -2 mm/s

0

200

400

600

800

1000

0 50 100 150 200 250 300

Time (s)

Load (N)

Interlaminar

In-Plane

3.3x10-4 mm/s

 (a) (b)

Figure 5. Typical examples of load versus time curves of SiC/BSAS composite tested in interlaminar and 

in-plane shear at 1100 °C in air at: (a) fast test rate of 3.3×10–2 mm/s, and (b) slow test rate of 3.3×10–4 

mm/s. 
NASA/TM—2003-212182 7 

 (i) 3.3×10–1 mm/s (ii) 3.3×10–5 mm/s 

(a) Interlaminar 

 (i) 3.3×10–1 mm/s (ii) 3.3×10–5 mm/s 

(b) In-plane 

Figure 6. Typical examples of fracture surfaces of SiC/BSAS composite at high (3.3×10–1 mm/s) and low 

(3.3×10–5 mm/s) test rates, tested in shear at 1100 °C in air: (a) interlaminar and (b) in-plane shear 

specimens. 
NASA/TM—2003-212182 8 

(a) 

(b) 

Figure 7. Cross sectional views showing shear planes, normal to fiber direction, of SiC/BSAS composite 

tested in shear at 1100 °C in air: (a) interlaminar and (b) in-plane specimens. 

2

1

1

2

Shear plane 

Shear plane 
NASA/TM—2003-212182 9 

This also can be seen from planes perpendicular to fiber direction, in which shear planes were reasonable 

straight in interlaminar specimens but were more tortuous in in-plane specimens, as shown in Figure 7. 

A more detailed fractography is underway to better understand the mechanisms associated with shear 

failure in conjunction with SCG and material direction. 

DISCUSSION

 The strength degradation in shear with decreasing test rate exhibited by the SiC/BSAS composite in 

this work was very similar to that in tension shown by not only CFCCs such as SiC/CAS, SiC/MAS, 

SiC/SIC, C/SiC, SiC/BSAS (2-D) [6,7,9] but advanced monolithic ceramics such as silicon nitrides, 

silicon carbides and aluminas [10]. The strength degradation in tension with decreasing test rate for 

advanced monolithic ceramics is known as a slow-crack growth (“dynamic fatigue”) phenomenon, 

commonly expressed by the following empirical power-law relation [11] 

n KI KIc v = α ( / ) (2) 

where v, KI and KIc are crack velocity, stress intensity factor and fracture toughness under mode I loading, 

respectively. α and n are called slow crack growth (SCG) parameters in mode I. Based on this powerlaw relation, the tensile strength (σf) can be derived as a function of applied test rate or stress rate (σ& ) 

with some mathematical manipulations [12–14]. 

 1

1

[ ] + = n σ f D σ& (3) 

where D is another SCG parameter associated with inert strength, n and crack geometry. Equation (3) can 

be expressed in a more convenient form by taking logarithms of both sides 

D

n f log log 1

1 log +

+

σ = σ& (4) 

 A test methodology based on Equation (3) or (4) is called constant stress-rate (“dynamic fatigue”) 

testing and has been established as ASTM Test Methods (C1368 [13] and C1465 [14]) to determine SCG 

parameters of advanced monolithic ceramics at ambient and elevated temperatures. As mentioned before, 

the data fit to Equation (4) has shown very reasonable even for CFCCs including SiC/CAS (1-D), 

SiC/MAS (1-, 2-D), SiC/SiC (2-D), C/SiC (2-D), and SiC/BSAS (2-D) in tension at 1100 to 1200 °C in 

air. This indicates that SCG or damage evolution/accumulation or delayed failure of those composites 

would be adequately described by the power-law type relation, Equation (1). The SCG parameter n is the 

most important life prediction parameter since it controls a measure of susceptibility to slow crack growth 

or damage evolution/accumulation and in turn to strength degradation, as implied from Equation 2. For 

monolithic ceramics, glasses or glass-ceramics, the susceptibility to SCG is typically categorized such 

that significant SCG lies in the range of in n<30, intermediate SCG in n =30-40, and insignificant SCG is 

in the range of n≥50. For the aforementioned composites, the values of n were all within n=10-20, hence, 

the composites can be termed ‘highly’ susceptible to SCG or damage evolution/accumulation. Typical 

toughened modern monolithic silicon nitrides at 1200 °C exhibit n ≥ 50. Therefore, the CFCCs exhibited 

a significantly greater susceptibility to SCG or damage evolution/accumulation than toughened 

monolithic silicon nitrides. Examples showing invariably significant strength degradation and 

consequently significant SCG susceptibility for some CFCCs in constant stress-rate testing in tension 

[6,7] are presented in Figure 8. 

 
NASA/TM—2003-212182 10 

APPLIED STRESS RATE, [MPa/s]

10-4 10-3 10-2 10-1 100 101 102 10

ULTIMATE STRENGTH, 

3

σf [MPa]

50

60

70

80

90

200

300

400

500

100

σ&

SiC/CAS

SiC/MAS

SiC/BSAS

SiC/SiC

(n=9)

(n=6)

(n=7)

(n=18)

TENSION

Figure 8. Examples of ultimate tensile strength as a function of test rate at elevated temperatures in air for 

some CFCCs including SiC/MAS (2-D; 1100 °C), SiC/CAS (1-D; 1100 °C), SiC/SiC (2-D woven; 

1200 °C), and SiC/BSAS (2-D; 1100 °C) [6,7]. 

APPLIED SHEAR STRESS RATE, [MPa/s]

10-3 10-2 10-1 100 101 102 103 104

SHEAR STRENGTH, 

τf [MPa]

7

8

9

20

30

40

50

60

70

80

90

10

100

ns=11.2

Hi-Nic/BSAS

τ&

ns=11.4

Interlaminar

In-plane

Unal [ref.5]

(DNS/1100oC)

Figure 9. Shear strength as a function of applied shear stress rate for SiC/BSAS composite at 1100 °C in 

air, reconstructed from the data in Figure 4 based on Eq. (10). The arrow in the ‘stress-rate’ axis indicates 

the stress rate used by Ünal and Bansal [5]. 
NASA/TM—2003-212182 11 

Based on the experimental results of the SiC/BSAS composite showing the dependency of shear strength 

on test rate, the reasonable data fit to log-log relation, and on the indication of SCG from the load-vs-time 

curves, the governing failure mechanism in shear can be assumed to be the one associated with SCG, 

similar in expression to the power-law relation of Equation (2). Hence, the following empirical slow 

crack velocity formulation is proposed here in shear, modified from the one (Equation 2) in tension, 

s n

s KII KIIc dt

da

v = = α ( / ) (5) 

where v, a, t, KII and KIIc are crack velocity, crack size, time, mode II stress intensity factor, and fracture 

toughness in mode II loading, respectively. αs and ns are slow crack growth (SCG) parameters in shear. 

In monotonic shear testing, constant displacement rate or constant load rate can be employed to test 

specimens so that shear stress applied to test specimens is a linear function of test time 

t dt t

t

τ = τ& =τ& ∫

0

( ) (6) 

where τ is shear stress, and τ& is applied shear stress rate determined from [τ& =P/(WLnt)] from load versus 

time curves in displacement control or directly from [τ& = P& /WLn] in load control with P& being load rate 

applied through a test frame. The generalized expression of stress intensity factor in shear for the case of 

an infinite body with a semicircular crack takes the form [15] 

1/ 2 K Y a II s = τ (7) 

where Ys is a crack geometry factor, typically Ys=1/(2√π). Using Equations (5) to (7) and following a 

similar procedure as used to drive Equation (3) in mode I, one can obtain shear strength (τf) as a function 

of applied shear stress rate as follows: 

1

1

[ ] + = s n

f Ds τ τ& (8) 

where 

1

1

2 [ ( 1) ] − + = + s s n n

s s s i D B n τ (9) 

where Bs=2KIIc/[αsYs

2

(ns-2)] and, τi is the inert strength that is determined in an appropriate inert 

environment or at a fast test rate whereby no or little SCG occurs. Equation (8) can be expressed in a 

more convenient form by taking logarithms of both sides 

s

s

f D

n

log log 1

1 log +

+

τ = τ& (10) 

which is identical in form to Equation (3) of mode I loading. SCG parameters ns and Ds in shear can be 

determined from slope and intercept of a linear regression analysis of the log (individual shear strength) 

versus log (individual shear stress rate) data based on Equation (10). The SCG parameter αs can be 

estimated from Equation (9) with appropriate constants and parameters associated. 

 Figure 9 shows the results of shear strength as a function of applied shear stress rate plotted based on 

Equation (10), with displacement rate being converted to shear stress rate. The slow crack growth 

parameters were determined as ns=11.2 and Ds=19.24 for interlaminar direction and ns=11.4 and Ds=33.27 

for in-plane direction. The corresponding coefficients of correlation (rcoef) of curve fit were 0.8442 and 
NASA/TM—2003-212182 12 

0.8850, respectively, for interlaminar and in-plane directions, showing reasonable data fit to the equation.* 

It is noteworthy that the value of ns, a measure of susceptibility to SCG, was identical to each other, 

regardless of material direction. The value of SCG parameter ns (=11) in shear also compares reasonably 

with that of n (=7) in tension for the similar 2-D SiC/BSAS composite tested at 1100 °C in air [7], 

indicating that the SiC/BSAS composite exhibited a significant susceptibility to SCG in shear as well. In 

general, the patterns of SCG for many monolithic advanced ceramics tested at elevated temperatures are 

well defined in fracture surfaces in terms of its configuration and size. By contrast, the patterns of SCG 

for CFCCs subjected to either in tension or in shear are not obvious and rather obscured by the nature of 

composite architecture so that it would be very difficult and/or a great challenge to derive a definite 

evidence of SCG from overall fracture surface examinations. Hence, the method proposed here –shear 

strength degradation with respect to test rate through constant stress-rate testing- would be a simple, quick 

and convenient way to check and quantify the degree of SCG in shear for CFCCs at elevated 

temperatures. The residual glassy phase may have been a major cause of SCG in the SiC/BSAS 

composite under shear. 

 Constant stress rate testing in tension has shown as a possible alternative to life prediction testing, as 

verified with stress rupture (constant stress) testing for various CFCCs at elevated temperatures at 1100 to 

1200 °C [6,7]. This indicated that the same failure mechanism might have been operative, independent of 

loading configuration that was either in monotonic loading (constant stress rate) or in static loading 

(constant stress). Therefore, it is expected that a single failure mechanism, SCG, would be dominant in 

shear for the SiC/BSAS composite regardless of loading configuration and that life prediction in shear 

from one loading configuration to another could be made analytically or numerically depending on the 

complexity of loading configurations. A phenomenological, simplified life prediction is proposed based 

on the following relation that was modified to account for shear loading from the relation primarily used 

for brittle monolithic ceramics in tension [16] 

 s

s

n

s

n

s

f n

D

t −

+

+ = ]τ

1 [

1

 (11) 

where tf is the time to failure, τ is applied shear stress, and ns and Ds are SCG parameters determined for 

each material direction in constant stress rate testing. The resulting life prediction curves in shear are 

shown in Figure 10 for both interlaminar and in-plane directions. Since the SiC/BSAS composite used in 

this work was no longer available, verification of the prediction by conducting constant stress (“stress 

rupture”) testing was not feasible so that only a method of prediction was proposed. An alternative 

composite is being sought for a future work of verification. It has been shown previously that under 

tension loading the overall life prediction of CFCCs was in reasonable agreement at least for a short 

period of lives with experimental data despite that the SCG formulation (Equation (1)) was 

phenomenological in nature without taking into account fiber/matrix architecture and interfacial 

microscopic mechanics of the composites. 

 The results of shear strength behavior of the 1-D SiC/BSAS showed that constant stress-rate testing, 

commonly utilized in determining life prediction parameters of monolithic ceramics in tension, could be 

applicable to determine phenomenological life prediction parameters of the composite material even in 

shear. This is also indicative that the overall failure mechanism in shear would be the one governed by 

the power-law type of slow crack growth. The merit of constant stress rate testing, incorporated with 

the power-law formulation, is enormous in terms of simplicity, test economy (short test times), and less 

 

*

 If average shear strength and average shear stress rate were used, the coefficients of correlation were rcoef=0.9708 and 0.9723, 

respectively, for interlaminar and in-plane directions. The corresponding SCG parameters were found to be ns=12 and Ds=19.4 

and ns=11 and Ds= 33.5, respectively, for interlaminar and in-plane directions. Hence, the difference in SCG parameters between 

individual data and average data approaches was negligible; however, the difference in coefficient of correlation was somewhat 

amplified. It is generally recommended that individual data approach be used in regression analysis for improved statistical 

accuracy [13,14]. 
NASA/TM—2003-212182 13 

TIME TO FAILURE, tf [s]

10-1 100 101 102 103 104 105 106 107

APPLIED SHEAR STRESS, τ [MPa]

5

6

7

8

9

20

30

40

50

60

10

ns=11

Hi-Nic/BSAS

ns=11

Interlaminar

In-plane

(DNS/1100oC)

Figure 10. A life prediction diagram of SiC/BSAS composite at 1100 °C under constant stress condition, 

constructed from the data of Figure 9 based on Eq. (11). Note that the diagram represents at a failure 

probability of approximately 50 percent. 

(strength) data scatter over other stress rupture or cyclic fatigue testing. Although the experimental 

results and phenomenological SCG model were presented in this work, a more detailed study regarding 

microscopic shear failure mechanisms associated with matrix/fiber interaction, matrix cracking and its 

effect on slow crack growth, and environmental effects, etc, [17–20] is still needed. Finally, the results of 

this work also suggest that care must be exercised when characterizing elevated-temperature shear 

strength of composite materials. This is due to that fact that elevated-temperature shear strength has a 

relative meaning if a material exhibits rate dependency: the shear strength simply depends on which test 

rate one chooses. Therefore, at least two test rates (high and low) are recommended to better characterize 

high-temperature shear strength behavior of a composite material, as suggested previously for the 

determination of ultimate tensile strength of CFCCs at elevated temperatures [6,7]. 

CONCLUSIONS

 

 The interlaminar and in-plane shear strengths of a unidirectional Hi-Nicalon fiber-reinforced barium 

strontium aluminosilicate (SiC/BSAS) composite were determined at 1100 °C in air as a function of test 

rate. For a given test rate, the interlaminar shear strength was approximately 70 percent greater than the 

in-plane shear strength. The composite’s shear strength exhibited a significant effect on test rate, 

regardless of orientation which was either in interlaminar or in in-plane direction, resulting in significant 

shear-strength degradation of about 50 to 60 percent as test rate decreased from 3.3×10–1 mm/s to 

3.3×10–5 mm/s. The rate dependency of shear strength was analogous to that of ultimate tensile strength 

at 1100 °C observed in a similar composite material (2-D SiC/BSAS) in which tensile strength decreased 

by about 60 percent when test rate varied from the highest to the lowest. A phenomenological power-law 

slow crack growth formulation was proposed to take account for shear-strength degradation behavior of 

the composite. Furthermore, constant shear-stress rate testing was proposed as a possible means of life 
NASA/TM—2003-212182 14 

prediction testing methodology to determine life prediction parameters of the composite in shear at 

elevated temperatures. 

REFERENCES 

1. P. Brondsted, F.E. Heredia, and A.G. Evans, “In-Plane Shear Properties of 2-D Ceramic 

Composites,” J. Am. Ceram. Soc., 77, 2569 (1994). 

2. E. Lara-Curzio and M.K. Ferber, “Shear Strength of Continuous Fiber Ceramic Composites,” ASTM 

STP 1309, p. 31, American Society for Testing & Material, West Conshohocken, PA (1997). 

3. N.J.J. Fang and T.W. Chou, “Characterization of Interlaminar Shear Strength of Ceramic Matrix 

Composites,” J. Am. Ceram. Soc., 76 2539–48 (1993). 

4. ASTM C 1425, “Test Method for Interlaminar Shear Strength of 1-D and 2-D Continuous Fiberreinforced Advanced Ceramics at Elevated Temperatures,” Annual Book of ASTM Standards, 

Vol.15.01, American Society for Testing & Materials, West Conshohocken, PA (2002). 

5. Ö. Ünal and N.P. Bansal, “In-Plane and Interlaminar Shear Strength of a Unidirectional Hi-Nicalon 

Fiber-Reinforced Celsian Matrix Composite,” NASA/TM—2000-210608, National Aeronautics and 

Space Administration, Glenn Research Center, Cleveland, OH (2000). 

6. S.R. Choi and J.P. Gyekenyesi, “Effect of Load Rate on Ultimate Tensile Strength of Ceramic Matrix 

Composites at Elevated Temperatures,” J. Comp. Appl. Mech., vol. 3 [1], 15–26 (2002); also in 

NASA/TM—2001-211125, National Aeronautics & Space Administration, Glenn Research Center, 

Cleveland, OH (2001). 

7. S.R. Choi, N.P. Bansal, and J.P. Gyekenyesi, “Ultimate Tensile Strength as a Function of Test Rate 

for Various Ceramic Matrix Composites at Elevated Temperatures,” NASA/TM—2002-211579, 

National Aeronautics & Space Administration, Glenn Research Center, Cleveland, OH (2002); also 

presented at CIMTEC 2002 Conference, paper # SV–4:L05 (be published in the proceedings), June 

14–18, 2002, Florence, Italy. 

8. N.P. Bansal and J.A. Setlock, “Fabrication of Fiber-Reinforced Celsian Matrix Composites,” 

NASA/TM—2000-210216, National Aeronautics & Space Administration, Glenn Research Center, 

Cleveland, OH (2000). 

9. B. F. Sorenson and J. W. Holmes, “Effect of Loading Rate on the Monotonic Tensile Behavior of a 

Continuous-Fiber-Reinforced Glass-Ceramic Matrix Composite,” J. Am. Ceram. Soc., 79[2] 313-320 

(1996). 

10. S.R. Choi and J.P. Gyekenyesi, “’Ultra’-Fast Fracture Strength of Advanced Structural Ceramics at 

Elevated Temperatures: An Approach to High-Temperature ‘Inert’ Strength,” pp. 27–46 in Fracture 

Mechanics of Ceramics, Vol. 13, Edited by R.C. Bradt, D. Munz, M. Sakai, V.Ya. Shevchenko, and 

K.W. White, Kluwer Academic/Plenum Publishers, New York (2002). 

11. S.M. Wiederhorn, “Subcritical Crack Growth in Ceramics,” pp. 613–646 in Fracture Mechanics of 

Ceramics, Vol. 2, Edited by R.C. Bradt, D.P.H. Hasselman, and F.F. Lange, Plenum Press, New 

York (1974). 

12. A.G. Evans, “Slow Crack Growth in Brittle Materials under Dynamic Loading Condition,” Int. J. 

Fracture, 10 251–259 (1974). 

13. ASTM C 1368, “Standard Test Method for Determination of Slow Crack Growth Parameters of 

Advanced Ceramics by Constant Stress-Rate Flexural Testing at Ambient Temperature,” Annual 

Book of ASTM Standards, Vol. 15.01, American Society for Testing and Materials, West 

Conshohocken, PA (2001). 

14. ASTM C 1465, “Standard Test Method for Determination of Slow Crack Growth Parameters of 

Advanced Ceramics by Constant Stress-Rate Flexural Testing at Elevated Temperatures,” Annual 

Book of ASTM Standards, Vol. 15.01, American Society for Testing and Materials, West 

Conshohocken, PA (2001). 

15. G.C. Sih, Handbook of Stress Intensity Factors, p. 3.2.1–2, Lehigh University, Bethlehem, PA (1973). 

16. J.E. Ritter, “Engineering Design and Fatigue of Brittle Materials,” pp. 667–686 in Fracture 

Mechanics of Ceramics, Vol. 4, Edited by R.C. Bradt, D.P.H. Hasselman, and F.F. Lange, Plenum 

Publishing Co., New York (1978). 
NASA/TM—2003-212182 15 

17. W.C. Curtin and H.G. Halverson, “High Temperature Deformation and Failure in Oxide/Oxide 

Composites,” HITEMP Review: Advanced High Temperature Engine Materials Technology Project,

NASAS/CP—1999-208915/VOL2, Paper 48, NASA Glenn Research Center, Cleveland, OH (1999). 

18. C.A. Lewinsohn, C.H. Henager and R.H. Johns, “Environmentally Induced Time-Dependent Failure 

Mechanism in CFCCS at Elevated Temperatures,” Ceram. Eng. Sci. Proc., 19[4] 11–18 (1998). 

19. C.H. Henager and R.H. Jones, “Subcritical Crack Growth in CVI Silicon Carbide Reinforced with 

Nicalon Fibers: Experiment and Model,” J. Am. Ceram. Soc., 77[9] 2381–94 (1994). 

20. S.M. Spearing, F.W. Zok and A.G. Evans, “Stress Corrosion Cracking in a Unidirectional CeramicMatrix Composite, J. Am. Ceram. Soc., 77[2] 562–70 (1994). 
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21

Dependency of Shear Strength on Test Rate in SiC/BSAS Ceramic Matrix

Composite at Elevated Temperature

Sung R. Choi, Narottam P. Bansal, and John P. Gyekenyesi

Ceramic composites; SiC/BSAS; Elevated-temperature shear strength; Slow crack growth;

Mechanical testing; Lifing

Unclassified -Unlimited

Subject Category: 07 Distribution: Nonstandard

Sung R. Choi, Ohio Aerospace Institute, Brook Park, Ohio 44142; Narottam P. Bansal and John P. Gyekenyesi,

NASA Glenn Research Center. Responsible person, Sung R. Choi, organization code 5920, 216–433–8366.

Both interlaminar and in-plane shear strengths of a unidirectional Hi-NicalonTM fiber-reinforced barium strontium

aluminosilicate (SiC/BSAS) composite were determined at 1100 °C in air as a function of test rate using double notch

shear test specimens. The composite exhibited a significant effect of test rate on shear strength, regardless of orientation which was either in interlaminar or in in-plane direction, resulting in an appreciable shear-strength degradation of

about 50 percent as test rate decreased from 3.3×10–1 mm/s to 3.3×10–5 mm/s. The rate dependency of composite’s

shear strength was very similar to that of ultimate tensile strength at 1100 °C observed in a similar composite (2-D

SiC/BSAS) in which tensile strength decreased by about 60 percent when test rate varied from the highest (5 MPa/s)

to the lowest (0.005 MPa/s). A phenomenological, power-law slow crack growth formulation was proposed and

formulated to account for the rate dependency of shear strength of the composite.