The degradation of silicon carbide fibers during exposure to conditions typical of composite fabrication was investigated. The tensile strength of pristine fibers and fibers sputtered with thin metal coatings were determined before and after treatment at 870 C for one hour in vacuum. Each fiber strength distribution was related by an analytical procedure to a projected composite ultimate tensile strength (PC UTS). The results indicate that a thin aluminum diffusion barrier can yield a 150 percent increase in PC UTS over the baseline SiC/Ti system

Brewer, W. D.

Landis, H. S.

Naidu, S. V. N.

Unman, J.

English

NASA Technical Reports Server

1111111111111111111111111111111111111111111111111111111111111111 
3 1176 00168 7202 
NASA Technical Memorandum 83122 
NASA-TM-8312219810016617 
MODIFICATION OF SILICON CARBIDE FIBERS 
FOR USE IN SIC/TI COMPdsITES 
H. S. LANDIS) ,J. UNNAM) W. D. BREWER) 
AND S. V. N. NAIDU 
NI\SJ\ 
National Aeronautics and 
Space Administration 
Langley Research Center 
Hampton, Virginia 23665 
1111111111111 IIII 11111 11111 IIIII IIIII IIII 1111 
NF00394 
tlBRf\HY COpy 
MAY 1 91981 
LAr~GLEY K=.SUII'C;-j C:=':NTER 
LIBRARY. NASA 
HAMPTON, VIRGINI" 
https://ntrs.nasa.gov/search.jsp?R=19810016617 2020-03-21T12:20:37+00:00Z
ERRATA 
NASA Technical Memorandum 83122 
MODIFICATION OF SILICON CARBIDE FIBERS 
FOR USE IN SiC/Ti COMPOSITES 
H. S. Landis, J. Unnam, W. D. Brewer, 
and S. V. N. Naidu 
Apri 1 1981 
Error in the title on the cover of the report. "Si/Ti" should read 
ISiC/Ti." 
Repl ace the cover of the report \'Ji th the attached corrected cover. 
Issued 5-12-81 
SUMMARY 
SiC/Ti composites have been marked by tensile strengths well below those 
predicted by rule-of-mixtures (ref. 1). A commercially available high-strength 
SiC fiber was found to degrade significantly after exposure to chemical and 
thermal conditions typical of composite consolidation. Several modifications 
to the fiber surface have been utilized to investigate fiber strenghthening 
mechanisms, and to develop techniques to minimize the loss of fiber strength 
during composite fabrication. SiC fibers were coated with various thin metal 
films and subjected to elevated temperatures typical of composite fabrication. 
Changes in fiber tensile properties were related to fiber-matrix interaction. 
The effects of fiber strength and strength distribution on expected composite 
properties were investigated analytically. 
INTRODUCTION 
In recent years, filament reinforced composites have gained increased 
acceptance in materials applications calling for high strength-to-weight and 
stiffness-to-weight ratios. Since polymer matrix composites degrade at 
moderately elevated temperatures, much work has been done on fiber reinforced 
metal-matrix composites. 
Titanium matrix systems, although theoretically very promising, have been 
hindered by developmental problems. Specifically, titanium is very reactive 
with most ceramic fibers at the high temperatures necessary for composite 
consolidation. Although silicon carbide is thermally and chemically one of the 
most stable fiber types known, significant amounts of brittle titanium carbides 
and titanium silicides have been shown to form at the fiber-matrix interface 
during consolidation (ref. 1). These compounds degrade the composite mechani-
cal properties as stress-intensifying microcracks form in the reaction layers 
at low strains (refs. 2, 3 and 4). 
This paper presents some preliminary results of research in which various 
surface modifications to the silicon carbide fiber were investigated in an 
attempt to reduce the amount of fiber-matrix interaction. Because of the high 
cost and long turn around time associated with composite fabrication, the 
titanium/silicon carbide system was studied in its component parts. Prelimi-
nary work in this study showed that no significant composite degradation can be 
ascribed to the titanium matrix (commercially pure TiA55) itself. Moreover, 
with the silicon carbide fiber used in this study, no degradation was associ-
ated with the heating of the fiber itself. It was therefore concluded that the 
investigation should focus on the interaction of the fiber and the titanium 
matrix, specifically on the fiber-matrix interface. 
The general approach has been to study the effects of thin-film diffusion 
barriers on this interaction, and the associated changes in tensile strength. 
Because the fibers used in this study have- intrinsically large variations in 
tensile strength, an analytical model was developed to predict composite 
strengths based on experimental fiber strength distributions. 
SYMBOLS 
Ec = composite modulus 
Ef = fiber modulus, from ref. 5 
Em = matrix modulus, experimentally determined 
Q = unbroken fibers/total fibers 
V = volume fraction of fibers 
~pp1 = stress applied to the composite 
2 
~ = fiber stress 
~r = fiber residual stress, from differences in the coefficients 
of thermal expansion for the fiber and the matrix 
~ = matrix stress 
~r = matrix residual stress 
~y = matrix yield stress, experimentally determined 
EXPERIMENTAL METHODS 
In this investigation, a composite was simulated by sputtering thin metal 
films (1-2 ~m) directly onto individual silicon carbide fibers. All specimens 
had an outer layer of titanium so that the diffusion boundary conditions 
resembled those of a composite. 
The coated fibers were subjected to a heat treatment typical of that 
necessary for composite consolidation, 87GoC for one hour under vacuum. Some 
aluminum coated fibers were given an initial heat treatment at 635°C for 10 
minutes, followed by the standard one hour anneal at 870°C. The reasons for 
this two-stage treatment will be discussed later in the paper. Although actual 
composite consolidation requires pressures in the 70-100 MPa range, no account 
was made for these mechanical pressures in the present study. 
Fibers were mounted for individual tensile testing as follows (see figure 
1): each fiber was bonded to two aluminum tabs with cyanoacrylate adhesive. 
The fixture was mounted by pins (to allow for self-alignment) in a displacement 
controlled Instron tensile testing machine, and the fiber loaded to failure. 
Figure 2 provides a comparison between an actual composite and a schematic 
diagram of an experimental specimen used to model the composite. Whether the 
specimen had an aluminum or tungsten diffusion barrier, or no barrier at all, 
3 
an outer layer of titanium was deposited so that the diffusion and reaction 
kinetics at the fiber-matrix interface resembled those in the actual 
composite. Although a thicker titanium coating would be desirable, sputtered 
coatings greater than 2 ~m were found to debond during handling, presumably 
because of large internal stresses generated during the sputtering process. 
Since the ultimate objective of this research is to develop a composite 
material, some method of predicting composite properties from the experimental 
fiber strength distributions was needed. Because of the wide scatter in the 
SiC fiber systems of this investigation, simple rule-of-mixtures predictions 
based on an average fiber strength were considered inadequate (refs. 6 and 7). 
The computer program developed in this project takes explicit account of the 
variations within a fiber strength distribution. 
The computer model is based on the following assumptions: strain is 
assumed to be uniform. That is, composite strain equals fiber strain equals 
matrix strain. The matrix is assumed to be elastic-ideally plastic, and the 
fibers are assumed to be elastic to failure. A fractured fiber is assumed to 
have no load carrying capacity, and no stress intensification is assumed in the 
neighborhood of such a fractured fiber. 
The resulting governing equations are given below. 
<rm < ~my : <rm ~ ([my : 
1) E = (I-V) E + QVEf c m 
a;ppl - (l-V)~my 
1) (f f = <rf r + ......;;.;,j....a.-_Q-V---=--
2) ~ = ~r + ~PPI Ef E 2) cf - rr m 'i my 
c 
3) a = (fmr + ~PPI Ef m 
Ec 
4 
RESULTS AND DISCUSSION 
In order to establish a baseline and account for any variability in the 
as-received and as-sputtered fibers, the fracture strength of each type of 
fiber before annealing was investigated. The as-fabricated fiber strength 
distributions are shown in figure 3. The wide scatter in the as-received SiC 
fiber strengths is typical of present fiber technology. A significant 
improvement in the filament strength distribution due to the deposition of 1 ~m 
of titanium was observed. One ~m of tungsten (followed by 1 ~m of titanium) 
was found to degrade the silicon carbide fibers, particularly the high strength 
fibers (ref. 6). One ~m of aluminum (followed by 1 ~m of titanium) was the 
most effective in increasing the as-sputtered fiber strengths, eliminating 
most of the low-strength fiber failures. 
The degradation asso~iated with annealing can be seen in figure 4. The 
most dramatic decrease in strength occurred for the fibers with no diffusion 
barrier (SiC/Ti). Since the same degradation is to be expected during actual 
composite consolidation, these fiber-matrix interactions may be a major reason 
why experimental composite strengths have been somewhat disappointing. 
Although tungsten degrades the fibers in the as-sputtered state, it does 
provide some protection after annealing, relative to SiC/Ti. The most 
encouraging results occur in the SiC/A1/Ti system. Although some loss of 
strength does occur during annealing, after heat treatment the SiC/A1/Ti fibers 
have a better strength distribution than even the as-received SiC fibers. 
(This assessment is made on the basis of the predictions of composite ultimate 
tensile strength by the computer model, using the actual fiber strength 
distributions.) 
The data presented in figure 4 as annealed SiC/A1/Ti fibers resulted from 
the two-stage heat treatment: a 10 minute pre-anneal at 635°C and the one hour 
5 
at 870°C common to all of the annealed fibers. The benefits of the 
pre-treatment can be seen in the comparison of the histograms in figure 5. 
The proposed mechanisms are as follows: if no pre-anneal is given, the 
diffusion of titanium toward the fiber is greatly facilitated as soon as the 
temperature exceeds the melting point of aluminum (660°C). The molten aluminum 
can also diffuse rapidly into the titanium matrix to form a solid solution. In 
this case, the protection afforded the fiber is minimal. 
On the other hand, at 635°C the formation of titanium aluminides is 
thermodynamically favored. These compounds have high melting temperatures and 
are relatively stable. After 10 minutes at 635°C, it is assumed that most of 
the aluminum layer has been converted to aluminides of various compositions 
(TiA13 to Ti3Al). These compounds provide an activation barrier to reaction 
compound formation. Locally high concentrations of aluminum at the 
fiber-matrix interface (due to the decomposition of aluminides and the 
formation of titanium carbides and silicides, refs. 8, 9 and 10), also tend to 
inhibit further reaction. Moreover, bulk diffusion of titanium through the 
aluminide layers is much more difficult than through molten aluminum and, 
apparently, through tungsten. In reality, the diffusion barrier is not 
aluminum, but the titanium aluminides. 
Some typical results from the computer program are shown in figure 6. The 
fiber strength distribution is for the SiC/Al/Ti fiber system after annealing 
and a 0.40 volume fraction of fibers is assumed. 
In the computer generated curve, the discontinuity at 0.002 strain 
represents the fracture of the low strength (500 MPa) fiber. The failure of 
this fiber causes an instantaneous decrease in the composite modulus, and a 
consequent increase in strain at the same load. At 0.003 strain the matrix 
stress becomes equal to the matrix yield stress and the composite modulus is 
6 
again lowered. At 0.0044 strain, the next fiber fails, beginning a chain 
reaction of increased fiber stress and more fiber failures, which results in 
all of the remaining fibers breaking at this load. It is the applied stress at 
this point which is presented as the projected composite ultimate tensile 
strength (UTS). 
The simple rule-of-mixtures plot shows a different type of behavior. Here 
the average fiber strength is the only fracture parameter. No account is taken 
of the low strength fiber, except as it affects the average. In this model, 
matrix yielding is postponed and composite failure is not predicted until the 
applied stress reaches nearly 1100 MPa, as opposed to 875 MPa for the computer 
model. 
By utilizing this computer program, each experimental fiber strength 
distribution was related to a single number of some significance: the 
projected composite UTS. 
The projected composite ultimate tensile strengths for each fiber system, 
assuming an 0.40 volume fraction of fibers, are presented in figure 7. A 
SiC/Ti composite with no diffusion barriers can be expected to have a UTS of 
only 350 MPa. A tungsten barrier offers a mild improvement, as 435 MPa is the 
projected composite UTS for the SiC/W/Ti system. However, the SiC/Al/Ti system 
promises an increase over baseline of 150% in composite UTS, to 876 MPa. 
On the basis of these preliminary results, further optimization of the 
aluminum duffusion barrier approach seems warranted. It is also expected that 
the knowledge gained from the study of this particular type of silicon carbide 
fiber will be applicable to newer higher-strength silicon carbide fibers 
currently being developed. 
7 
CONCLUSIONS 
The important results of this research program can be summarized in three 
main points. First, an aluminum coating was found to increase the average 
as-fabricated SiC fiber strength 35% over the baseline SiC fibers. Second, the 
aluminum coating was found to increase the average annealed fiber strength 250% 
over baseline SiC/Ti fibers, provided optimized heat treatments are used. 
Finally, the aluminum diffusion barrier approach was found to increase the 
projected composite UTS 150% over baseline, based on the computer model 
developed for this research. 
REFERENCES 
1. L. House, "A Study of Interfacial Reactions in Silicon Carbide Reinforced 
Titanium," M.S. Thesis, The George Washington University, May 1979. 
2. A. Metcalfe and M. Klein, "Effect of the Interface on Longitudinal Tensile 
Properties," Composite Materials Vol. 1, Interfaces in Metal Matrix 
Composites. Academic Press, New York 1974, pp. 127-168. 
3. M. Shorshorov et. al., "Brittle Interface Layers and the Tensile Strength 
of Metal Matrix-Fibre Composites,·· Journal of Materials Science 14 (1979) 
pp. 1850-1861. 
4. S. Oehiai and Y. Murakami, "Tensile Strength of Composites with Brittle 
Reaction Zones at Interfaces." Journal of Materials Science 14 (1979) pp. 
8 
5. H. Debolt and V. Krukonis, "Improvement of Manufacturing Methods for the 
Production of Low Cost Silicon Carbide Filament," Technical Report 
AFML-TR-73-140, December 1972. 
6. C. Zweben, "New Understanding of Fiber Composite Materials," NASA Tech 
Brief 71-10161, NASA Pasadena Office, June 1971. 
7. D. Karpinos et. al., "Mechanical Properties of Plastics Reinforced with 
Silicon Carbide," Problemy Prochnosti 7 (7) July 1975, pp. 87-89. 
(Translated from Russian) 
8. W. Brewer, J. Unnam and D. Tenney, "Mechanical Property Degradation and 
Chemical Interactions in a Borsic/Titanium Composite," 24th National SAMPE 
Symposium and Exhibition, San Francisco, CA, May 1979. 
9. A. Metcalfe, "Physical Chemical Aspects of the Interface," Composite 
Materials Vol. 1, Interfaces in Metal Matrix Composites. Academic Press, 
New York 1974, pp. 67-123. 
10. R. Pailler et al., "Chemical Interaction Phenomena at High Temperature 
Betweeen Boron Fibers and Titanium Metal (or TA6V alloy)," TMS-AIME, Fall 
Meeting, Chicago, Oct. 1977. 
9 
SiC fiber 
<= 10 I / 
, 0] ==> 
Al tab 
Figure 1. - Fiber tensile test fixture. 
SiC fiber 
142 /-L m 
'i \ Diffusion barrier 
..... 1 /-Lm 
layer 
.... 1 /-L m 
Figure 2.- Composite cross-section and schematic diagram of an experimental fiber specimen. 
.4 
.2 SiC 
0 500 1000 1500 2000 2500 3000 3500 
.4 
.2 ~ SiC/Ti 
tll 
..... 
tll 
Q) 
..... 
'+-4 0 500 1000 1500 2000 2500 3000 3500 0 
s:: 
0 
..... 
..... 
.4 C) 
ro 
M 
~ 
.2 I- J ~ SiC/W/Ti I I I 
.. --- -.--.- . 
0 500 1000 1500 2000 2500 3000 3500 
.4 
.J SiC/Al/Ti 
o 500 1000 1500 2000 2500 3000 3500 
Fiber fracture strength, MPa 
Figure 3.- As-fabricated fibers strength distributions. 
.4 
.2 
0 
.4 
.J 
Ul 
..... 
Ul 
Q) 
0 ..... 
'+-4 
0 
s:: 
0 
.4 ..... 
..... 
u 
cod 
M 
.2 . ~ 
0 
.4 ) 
o 
500 1000 
L 
500 1000 
- .. I I I I .... I I ........ _ 
500 1000 
500 1000 
As-fabricated 
lW/aWA Annealed 
1500 2000 
1500 2000 
- -
~ ~ I 
1500 2000 
1500 2000 
Fiber fracture strength, MPa 
SiC 
2500 3000 
SiC/Ti 
2500 3000 
SiC/W/Al 
I 
2500 3000 
SiC/Al/Ti 
2500 3000 
Figure 4.- As-fabricated and annealed fiber strength distributions. 
3500 
3500 
3500 
3500 
Annealed 
.2J- m ~ ~ 870oC, 1 hr 
to 
....., 
to 
Q) 
....., 
Cl 0 500 1000 1500 2000 2500 
s:: 
0 
..... 
°l ....., C,) Annealed ~ f..4 6350 C, 10 min ~ 8700 C, 1 hr 
.2 
o 500 1000 1500 2000 2500 
Fiber fracture strength, MPa 
Figure 5.- Effects of heat treatments on SiC/ Al/Ti fibers. 
~ 
Pot 
~ 
... 
rt.l 
rt.l 
Cl) 
M 
~ 
u.l 
rt.l 
t.i .4 
Cl) 
~ 
'+-4 
o 
s:: 
o 
.... 
~ 
C) 
~ 
M 
~ 
1200 
1000 
800 
600 
400 
200 
o 
.2 
o 
Annealed SiC/ Al/Ti 
500 1000 1500 2000 2500 
Fiber fracture strength, MPa 
Computer model Simple rule-of-mixtures 
1200 
1000 
~ 800 
Pot 
~ 
... 600 rt.l 
rt.l 
Cl) 
M 
-I-> 
u.l 400 
200 
.002 .004 .006 o .002 .004 .006 
Strain, mm/mm Strain, mm/mm 
Figure 6.- A comparison of the computer program and simple rule-of-mixtures predictions 
of composite behavior. 
1000 [ 
('(! 
~ 
~ 800 I-
~ 
til l-E-! 
::> 
Q) 600 I-
..... 
..... 
Ul 
0 ..... 
0. 
S 
0 400 l-e,) 
"d 
Q) I-
..... 
e,) 
Q) 
.,......, 200 I-0 
M 
~ 
I 
0 
SiC/Ti SiC/W/Ti SiC/ Al/Ti 
Figure 7. - Proj ected compos ite strengths. 
I 2. Government Accession No. 1. Report No. 
NASA TM-83122 
4. Title and Subtitle 
Modification of Silicon Carbide Fibers for 
Use in SiCjTi Composites 
* ** *** i't.Au~~r(slandis, J. Unnam, W. D. Brewer, 
and S. V. N • N aid u ** 
3. Recipient's Catalog No. 
5. Report Date 
Apri 1 1981 
6. Performing Organization Code 
506-53-23-01 
8. Performing Organization Rp.p.>rt No. 
f-----------------------------i 10. Work Unit No. 
9. Performing Organization Name and Address 
NASA Langley Research Center 
Hampton, VA 23665 11. Contract or Grant No. 
12. Sponsoring Agency Name and Address 
National Aeronautics and Space Administration 
Washington, DC 20546 14. Sponsoring Agency Code 
15. Supplementary Notes 
*George Washington University, Joint Institute for the Advancement of Flight Sciences 
NASA Langley Research Center, Hampton, Virginia 23665 
**Virginia Polytechnic Institute and State University, Blacksburg, Virginia 24061 
***NASA Lanql ev Resea rch Cen te r~ Hampton,. Vi rqi n i a 23665 
16. Abstract 
The degradation of silicon carbide fibers during exposure to condi-
tions typical of composite fabrication was investigated. The tensile 
strength of pristine fibers and fibers sputtered with thin metal coat-
ings (Ti, Al, W) were determined before and after treatment at 870°C for 
one hour in vacuum. Each fiber strength distribution was related by an 
analytical procedure to a projected composite ultimate tensile strength 
(PC UTS). The results indicate that a thin aluminum diffusion barrier 
can yield a 150% increase in PC UTS over the baseline SiCjTi system. 
17. Key Words (Suggested by Author(s)) 
Composite 
Filament 
Silicon Carbide 
Coating 
19. Security Classi!. (of this report) 
Unclassified 
18. Distribution Statement 
20. Security Classif. (pf this P.<Igtl). UnCIaSS1Tled 
Unclassified - Unlimited 
Subject Category 24 
21. No. of Pages 
16 
22. Price· 
A02 
N-305 For sale by the National Technical Information Service, Springfield. Virginia 22161 
End of Document 


Modification of silicon carbide fibers for use in SiC/Ti composites

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