THE EFFECT OF STRAIN RATE ON THE TENSILE DEFORMATION OF Ti-6Al-4V/SiC COMPOSITES

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1 Scripta mater. 44 (2001) THE EFFECT OF STRAIN RATE ON THE TENSILE DEFORMATION OF Ti-6Al-4V/SiC COMPOSITES F. Gálvez, C. González, P. Poza and J. LLorca Department of Materials Science. Polytechnic University of Madrid E. T. S. de Ingenieros de Caminos Madrid, Spain (Received November 2, 2000) (Accepted in revised form January 15, 2001) Keywords: Composites; Mechanical properties; Strain rate; Fracture Introduction The search for lighter, stiffer and stronger materials has led to the reinforcement of light alloys with ceramic fibers. In this respect, Ti alloys reinforced with SiC monofilaments stand out as the materials with the best performance in term of specific stiffness and strength at intermediate temperatures (300 C to 700 C). They were mainly aimed at gas turbine engines for aerospace vehicles, and significant progress has been achieved over the last ten years, although their future is still dubious due to cost and other considerations (1 2). The effect of temperature and loading conditions (static or cyclic) on the mechanical behavior of these composites has been extensively analyzed, but there is no information available on their performance at high strain rates, so this investigation is focussed on this question. Apart from of its scientific interest, the response of fiber-reinforced Ti-alloy composites to those conditions is important from the practical viewpoint, as gas turbine components are often subjected to high strain rate and impact loading. Material and Experimental Techniques The composite material analyzed in this investigation was a Ti-6Al-4V alloy uniaxially reinforced with 35 vol. % Sigma 1140 SiC fibers. Composite panels of 1.35 mm in thickness were manufactured by DERA (UK) by the foil-fiber-foil method. The panels were consolidated in vacuum at 940 C during 1 hour under 30 MPa of pressure. The matrix filled completely the spaces between the fibers, whose centers formed an imperfect hexagonal array. The Sigma 1140 fibers were processed by chemical vapor deposition of -SiC on a 15 m diameter tungsten wire, and a carbon coating of 4 m was finally deposited by the same method. Dog-bone specimens of 15 mm gage length and 3 mm width were machined with the fibers oriented in the loading direction using a water-jet. Tensile tests were carried out in a servo-hydraulic testing machine with nominal cross-head speeds of 0.08 mm/min and 9 mm/min, which correspond to average strain rates of approximately s 1 and210 3 s 1. Loads and strains were measured with a standard load cell and a resistive extensometer of 12.5 mm gage length, respectively. The high strain rate tests were performed in a Hopkinson-split pressure bar. A scheme of the experimental set-up is shown in Fig. 1. A hollow cylindrical projectile impacts the free end of the incident bar, inducing a tensile elastic wave which travels along the incident bar in the opposite /01/$ see front matter Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S (01)

2 2668 Ti-6Al-4V/SiC COMPOSITES Vol. 44, No. 11 Figure 1. Scheme of the experimental set-up for the tensile tests at high strain rate. direction and is transferred to the output bar through the specimen. The force F transmitted through the composite specimen as a function of time can be computed as F(t) E (t), where, E, and stand, respectively, for the area, the elastic modulus, and the elastic strain in the output steel bar, the latter measured with a strain gage bonded to the bar. Special care was taken to design the gripping system for the composite specimens, so that its influence on the wave propagation through the specimen was minimized (3). The specimen strain as a function of time (and, thus, the strain rate) was determined with another strain gage glued to composite specimen. A detailed description of the hypotheses on which the experimental set-up is based can be found elsewhere (4). They show that a few microseconds are spent to reach quasi-static equilibrium; it was not possible to determine accurately the initial part (the elastic slope) of the tensile stress-strain curve from the load and strain data. Metallographic sections were prepared in the longitudinal direction from the specimens broken at different strain rates. The composite samples were embedded in a conductive resin, polished on SiC paper to 500 grit finish followed by a diamond slurry (up to 1 m) and finally on magnesia. They were observed in the scanning electron microscope together with the fracture surfaces to analyze the failure micromechanisms. Results and Discussion Representative stress-strain curves obtained at different strain rates are shown in Fig. 2a. Those measured at low strain rate exhibit a bilinear shape, which is typical of this kind of composites when loaded in the fiber direction (5). The initial part corresponds to the elastic deformation of the matrix and the fibers, and its slope was equal to 186 GPa. This value is very close to the theoretical result computed from the elastic modulus of the Ti alloy (E m 110 GPa (2)), the fibers (E f 330 GPa (6)), and the volume fraction of the latter, f, using the rule of mixtures, E E ƒ ƒ E m 1 ƒ 187 GPa (1) which indicates that an isostrain model is adequate to represent the composite deformation in the fiber direction. The elastic region ends when the matrix begins to deform plastically, and it is followed by another linear region where the stress increase is mainly borne by the elastic deformation of the fibers as the Ti-matrix shows very little strain hardening in the annealed condition. Consequently, the slope of the stress-strain curve in this region is close to fe f 115 GPa. The experimental results showed that the composite yield stress (represented by the beginning of the matrix plastic deformation) increased with the strain rate, while the overall ductility of the composite was significantly reduced (Fig. 2b). As a result, the second linear region in the stress-strain curve

3 Vol. 44, No. 11 Ti-6Al-4V/SiC COMPOSITES 2669 Figure 2. (a) Representative stress-strain curves at different strain rates. (b) Influence of strain rate on the ductility ( u ) and the tensile strenth ( u ). practically disappeared in the specimens tested at 500 s 1 (Fig. 2a). The composite tensile strength was not, however, modified with the strain rate, very likely because the negative effect of the ductility was offset by the increase in yield strength. The increase in the composite yield stress with strain rate can be easily rationalized from the matrix behavior. Follansbee and Gray (7) and Ruiz et al. (8) reviewed the effect of the strain rate on the mechanical properties of Ti-6Al-4V alloys, and their results included data for materials in the annealed condition with an - microstructure, which is similar to that of the matrix in the composite. They reported that the flow stress increased on average approximately from MPaat10 3 s 1 to about MPa at 100 s 1, i. e. a rate sensitivity of 34 MPa/decade, which is in perfect agreement with the composite behavior. In addition, the strain hardening of the matrix at high rates of strain was negligible. The analysis of the longitudinal sections near the fracture surface showed that damage in the form of multiple fiber fracture was concentrated a few hundred microns above and below the fracture surfaces (Fig. 3a). Two or three fiber cracks were seen in each fiber, and the distance between the cracks was in the range 100 m to300 m. The extent of damaged localization did not change significantly with the strain rate, although the fiber crack spacing was qualitatively shorter in the samples tested at 500 s 1. This is in agreement with the shorter fiber pull-out length observed in the fracture surfaces (Fig. 3b), but this may be a consequence of the reduction in ductility with strain rate. The longitudinal sections showed that damage prior to failure was localized in a given section of the composite. The failure sequence started with the random fracture of the SiC fibers along the gage

4 2670 Ti-6Al-4V/SiC COMPOSITES Vol. 44, No. 11 Figure 3. (a) Longitudinal section of a specimen broken at 500 s 1 near the fracture surface. (b) Fracture surface of a specimen broak at 500 s 1. (c) Detail of a fiber fractured at s 1. (d) Detail of a fiber fractured at 500 s 1. length. The load carried by the broken fiber has to be redistributed and is mainly transferred to the adjacent fibers through the matrix, increasing the failure probability of these fibers upon further loading. The failure probability of the surrounding fibers will be further enhanced when one of the neighboring fibers breaks, leading sooner or later to a chain-reaction which will induce specimen fracture by the successive fracture of the fibers followed by the ductile fracture of the matrix. As a result, the composite ductility would depend on whether the failure of one fiber leads to a significant increase in the load carried by the intact neighbor fibers. This behavior (denominated local load sharing) precipitates the composite failure and leads to a reduction in strength, as opposed to global load sharing, where the load borne by the broken fiber is distributed equally among the rest of the fibers, eliminating the localization of damage in the early stages of deformation (9). Obviously, the magnitude of the stress concentration induced by one fiber fracture in the neighboring fibers depends on a number of factors. If the matrix has deformed plastically and cannot harden any more, the stress borne by the broken fiber is transferred by shear deformation of the matrix or by

5 Vol. 44, No. 11 Ti-6Al-4V/SiC COMPOSITES 2671 frictional sliding at the fiber/matrix interface and, all other factors being equal, the maximum stress in the neighbor fibers increases with the shear stress (10). The analyses at higher magnification of the longitudinal surfaces showed that the fiber cracks were stopped at the C coating in the specimens tested at low strain rates (Fig. 3c) and further composite deformation led to the relative sliding between the fiber and the C coating. The fiber cracks propagated, however, into the C coating in the specimens tested at 500 s 1 and were stopped at the C/matrix interface (Fig. 3d), the load being transferred in this case by the shear deformation of the matrix and by the relative sliding at the C/matrix interface. The differences between the fiber crack opening displacements in Figs. 3c and 3d demonstrated that sliding resistance at the C/matrix interface was significantly higher than that at the C/fiber interface, and the results of push-out tests reported in a separate investigation are in agreement with this conclusion (11). As a result, the stress concentration around a broken fiber was higher in the specimens tested at high strain rates, leading to the reduction in ductility experimentally observed. Conclusions Specimens of Ti-6Al-4V matrix uniaxially reinforced with Sigma 1140 SiC fibers were tested in tension in the fiber direction at strain rates in the range s 1 to 500 s 1. The composite yield strength increased with the strain rate as a result of the matrix strain rate sensitivity. Analysis of longitudinal sections of broken specimens showed that damage in the form of fiber fracture was localized very near the fracture surfaces, indicating the predominance of local load sharing conditions in all cases. However, the fiber cracks were deflected at the C coating at low strain rates, while they penetrated the C coating and were stopped by the Ti-matrix at high rates of strain. As a result of this difference in the failure micromechanisms, the stress concentration in the neighbor fibers caused by one fiber failure was higher in the specimens tested at high strain rates and this led to a marked reduction in the overall composite ductility. The ultimate strength reflected the trade off between the increase in yield strength and the reduction in ductility and was insensitive to the rate of the deformation. Acknowledgments The financial support from CYCIT and the European Union through grant 2FD97-89-C02-02 is gratefully acknowledged. References 1. J. M. Larsen, S. M. Russ, and J. W. Jones, Metall. Mater. Trans. 26A, 3211 (1995). 2. S. Mall and T. Nicholas, ed., Titanium Matrix Composites. Mechanical Behavior, Technomic, Lancaster, PA(1998). 3. I. S. Chocron, J. Rodríguez, M. A. Martínez, and V. Sánchez-Gálvez, Int. J. Impact Eng. 19, 135 (1997). 4. J. A. Zukas, High Impact Dynamics, John Wiley & Sons, New York (1990). 5. C. H. Weber, X. Chen, S. J. Connell, and F. W. Zok, Acta Metall. Mater. 42, 3443 (1994). 6. C. González and J. LLorca, Adv. Comp. Lett. 9, 295 (2000). 7. P. S. Follansbee and G. T. Gray III, Metall. Trans. 20A, 863 (1989). 8. D. Ruiz, P. Hardy, and J. Harding, Report JH/000/1989, Solid Mechanics Group, Department of Engineering Science, University of Oxford (1989). 9. M. Y. He, A. G. Evans, and W. A. Curtin, Acta Metall. Mater. 41, 871 (1993). 10. C. M. Landis and R. M. McMeeking, Int. J. Solids Struct. 44, 4333 (1999). 11. C. González, Ph.D. thesis. Polytechnic University of Madrid (2000).