- Email: [email protected]
Mechanical behavior and fatigue damage of a titanium matrix composite reinforced with continuous SiC fibers
Materials Science and Engineering A 452–453 (2007) 536–544
Mechanical behavior and fatigue damage of a titanium matrix composite reinforced with continuous SiC fibers D. Bettge a , B. G¨unther a , W. Wedell a , P.D. Portella a , J. Hemptenmacher b , P.W.M. Peters b , B. Skrotzki a,∗ a
Federal Institute for Materials Research and Testing (BAM), 12200 Berlin, Germany b German Aerospace Center (DLR), 51147 Cologne, Germany
Received 31 July 2006; received in revised form 18 October 2006; accepted 20 October 2006
Abstract A titanium (Ti-6242) matrix composite reinforced with continuous SiC fibers was studied. The mechanical behavior of the matrix and the composite was characterized by tensile, creep and isothermal fatigue tests at room temperature and up to 550 ◦ C. The thermo-mechanical fatigue behavior under in-phase and out-of-phase conditions was investigated for the composite between 100 and 550 ◦ C. Fracture surfaces were characterized by confocal light microscopy and by scanning electron microscopy to identify the damage mechanisms. © 2006 Elsevier B.V. All rights reserved. Keywords: Metal matrix composites; Tensile strength; Creep strength; Isothermal fatigue; Thermo-mechanical fatigue; Damage analysis
1. Introduction Near-␣ titanium alloys are classical high temperature materials for applications in the temperature range between 500 and 550 ◦ C. They were developed to meet demands for higher operating temperatures in the compressor section of aircraft engines [1]. These alloys combine the good creep resistance of ␣-alloys with the high strength of (␣ + )-alloys. During the last few years, new engine concepts were developed with higher power densities, which increased the demands with respect to mechanical and thermal loads even further [2]. At the same time, there is the need for consistent light weight design which requires essentially new material concepts. Intermetallic titanium aluminides and reinforced titanium matrix composites (TMCs) are examples for such developments. Reinforcing titanium alloys with continuous SiC fibers for applications in aero engines represents a class of materials, which combines the high strength, stiffness and creep resistance of the ceramic SiC fiber with the damage tolerance of titanium alloys. The use of the light ceramic fiber (ρ = 3.0 g/cm3 , E = 380 GPa) in the titanium alloy matrix reduces the density of the latter even further and provides the potential for consider-
∗
Corresponding author. Tel.: +49 30 8104 1520; fax: +49 30 8104 1527. E-mail address: [email protected] (B. Skrotzki).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.10.107
able weight savings in aero engines when replacing bulk titanium alloy parts. Potential applications are redesigned discs and shafts of the compressor section of the engines. Disks are supposed to become rings and eventually monolithic bladed rings (“blings”). Presently, a major obstacle on the way to applications for safety relevant components made of TMCs is the complex fatigue damage mechanism. In the last decade, the isothermal fatigue behavior of TMCs has been investigated at room temperature [3] and at elevated temperature [4,5]. It is evident that the room temperature fatigue behavior of long fiber reinforced TMCs can be interpreted using general concepts for metal matrix composites [6]. The damage mechanisms commonly are fiber failure (at high stresses and low cycles), matrix cracking (at intermediate stresses and intermediate cycles), and no damage or damage not giving rise to final failure (at low stresses and high cycles). At high temperatures, the fatigue behavior generally can be described similarly. Realistic testing of TMCs for aircraft engine applications, however, incorporates the variation of load as well as of temperature, i.e. thermo-mechanical loading. Thermo-mechanical fatigue (TMF) tests are commonly used to characterize the simultaneous influence of load and temperature on fatigue life of TMCs. Some extensive studies on the behavior of titanium matrix composites containing SiC fibers under TMF loading were published (e.g. [7–12]). The fatigue life of TMCs strongly depends on the experimental details of the TMC tests. Hold
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
times have a major influence on the damage mechanisms because the matrix alloy becomes very soft at elevated temperatures as opposed to the SiC fibers. It was found that out-of-phase (maximum load applied at minimum temperature) thermo-mechanical loading generally leads to a stronger reduction in fatigue life than in-phase (maximum load applied at maximum temperature) loading [9,10]. This seems to be controlled by matrix cracking whereas failure during in-phase cycling appears to be dominated by fiber failure [11]. In the present work, results of uniaxial tensile and creep tests as well as of isothermal and thermo-mechanical fatigue tests are presented for the material system Ti-6242/SiC. The mechanical data were determined to identify the material parameters of a viscoplastic material model, which will be used to simulate the material behavior of the composite. These results will be published elsewhere, first data are given in Ref. [13]. 2. Experimental procedure The alloy Ti-6242 was used as matrix material. This alloy is mainly employed in gas turbines for rotating components. Its nominal composition is Ti–6%Al–2%Sn–4%Zr–2%Mo– 0.1%Si (wt.%). Silicon is added for improved creep resistance [1]. The alloy has a near-␣ microstructure and finds its applications in rotating components at temperatures up to 540 ◦ C
537
[14]. Ti-6242 is widely used in the aerospace industry and its behavior is well characterized. Therefore, it was chosen as a suitable matrix material for this study. The unreinforced matrix was available as thermo-mechanically processed rod material with an ␣-grain size of about 4 m, Fig. 1a. The SiC fiber SCS-6 (Textron Specialty Materials, Lowell, MA, USA) was used as long fiber reinforcement. The fiber has an outer diameter of 142 m and a carbon core of 30 m diameter. The fiber is covered with a protection layer, which consists mainly of carbon. The composite material was processed by deposition of the matrix alloy on the SCS-6 fibers using the magnetron sputtering process [15]. Subsequently, the coated fibers are cut to the required length and stacked into a tube of Ti-6242. The preform is hipped at 950 ◦ C and 190 MPa [16]. The fiber volume fraction in the gauge length was measured to be about 35%. The grain size of the sputtered matrix material was about 1 m, Fig. 1b–d. Cylindrical specimens were machined with an overall length of 81 mm, a gauge length of 21 mm (unreinforced matrix) and 10 mm (TMC), and a gauge diameter of 3.5 mm. The transition to the threaded heads was machined to a parabolic shape. The surface of the gauge length was mechanically polished (diamond suspension, min. size 3 m). The spatial distribution of fibers in the gauge length visualized by X-ray computer tomography (CT) is shown in Fig. 2 in a cross-section. From the CT-images,
Fig. 1. (a) Ti-6242 matrix consisting of dark ␣-grains and bright -grains (SEM, BSE-mode). (b) SEM-image of the TMC; overview at low magnification. (c) TMC at higher magnification. (d) TMC at higher magnification. Grain size is smaller than in the unreinforced matrix due to sputtering.
538
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 2. Low resolution CT-image of the TMC-test piece cross-section in the gauge length.
the actual fiber volume fraction was determined for each test piece. The yield strength and ultimate tensile strength were determined in tensile tests at room temperature and at 550 ◦ C. Constant strain rates of 10−3 and 10−4 s−1 were applied. Yielding was determined at 0.2% offset. Most of the tests on the matrix material were interrupted after reaching 10% strain. All fatigue tests were accomplished on a 100 kN MTS servohydraulic testing system equipped with induction heating. For LCF and TMF testing, the specimen temperature was measured using two ribbon thermo-couples (Pt-RhPt type S) at the upper and lower end of the gauge section. The temperature gradient in the gauge length was less than ±15 ◦ C. All fatigue tests were carried out in air under load control with a stress ratio of R = 0.1. The reason to run tests under load control is (i) that the ceramic fibers fail already at small strains and (ii) that a compressive stress state had to be avoided because the specimen tend to buckle due to the small diameter in the gauge length. The frequency of the isothermal tests was 0.25 Hz for an estimated lifetime of up to 105 cycles, 1 Hz for up to 106 cycles and 5 Hz exceeding 106 cycles. For TMF-tests, the temperature was cycled with the same frequency as the stress, Fig. 3, with a temperature rate of 5 K/s and an upper and lower temperature of 550 and 100 ◦ C, respectively, resulting in a heating and cooling time intervals of 1.5 min each. After each heating and cooling ramp, the temperature was held for 2 min. Consequently, a full TMF cycle lasted 7 min. The applied stress was varied in-phase (IP) or out-of-phase (OP) to temperature with R = 0.1. A high temperature extensometer with a gauge length of 10 mm and a precision of 1% was used. Uniaxial creep tests were carried out under constant tensile load on the matrix material at initial stresses of 80–200 MPa and at a temperature of 550 ◦ C. One test each was run at 450 and 500 ◦ C at σ = 150 MPa. The TMC was tested at 550 ◦ C and at stresses varying between 1100 and 1400 MPa. The geometry of the creep specimen was the same as for the fatigue tests. The strain measurement was performed using a high temperature extensometer with 21.5 mm gauge length. The creep strain was measured, i.e. the time dependent strain that occurs after the
Fig. 3. Schematic representation of the temperature-time- and stress-time-path of the TMF-tests.
complete application of the load. Three Pt-RhPt (type S) thermocouples were connected to the upper and lower end and in the center of the gauge length. A standard three zone cylindrical radiation furnace was used for heating. The temperature gradient within the gauge length was less than ±3 ◦ C. A preload of 10% of the total load was applied. The composite material and the fracture surfaces were characterized using conventional light microscopy (LM), a confocal light scanning microscope (CLSM) Leica DSCM1000, scanning electron microscopes (SEM) CamScan 2 and LEO Gemini 1530 VP. Low resolution computer tomography has been applied to characterize each individual test piece with respect to the actual number of fibers. 3. Results 3.1. Tensile tests Tensile properties were determined for the matrix and for the composite material at room temperature and at 550 ◦ C, as given in Fig. 4. Ti-6242 has a yield strength Rp0.2 of 1163 MPa at room temperature (strain rate 10−3 s−1 ) which is considerably reduced to 523 MPa at 550 ◦ C. Applying a lower strain rate of 10−4 s−1 reduces the strength at RT somewhat, while the strength at 550 ◦ C is decreased by 50% due to creep effects. The fiber reinforcement leads to a clear improvement of the mechanical behavior at both test temperatures (Rp0.2 = 1907 MPa at room temperature, and Rmax = 1518 MPa at 550 ◦ C). The temperature dependence of the tensile behavior of the metallic matrix is shown in Fig. 4b. All strength data are summarized in Table 1. The mean strength of a single SiC fiber at room temperature is about 4600 MPa which drops only a little if temperature is increased up to 550 ◦ C [17]. 3.2. Creep tests The matrix material and the composite were tested in shorttime creep tests. Creep strain ε versus time t was measured. The creep data were then further analyzed and represented as creep rate ε˙ versus creep strain ε in Fig. 5a and b. The Ti-
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
phenomenological equation Qc σn, ε˙ = A exp − RT
539
(1)
where A represents a constant, Qc is the apparent activation energy for creep, R the universal gas constant, T the temperature, σ the creep stress and n is the stress exponent. Analyzing the plot of log creep rate versus log stress (not shown) provides the stress exponent n which was determined to be 3.2 for the matrix at 550 ◦ C. It was not possible to determine the stress exponent for the TMC because the tests were interrupted before the minimum creep rate was reached. Fig. 5c shows results of the short-time creep tests on TMC material in terms of applied stress versus time to fracture. Results obtained in this work are compared to those published in the literature. The figure shows that due to the high homogeneity with respect to fiber distribution and matrix microstructure, the material produced by DLR shows very high creep strength. 3.3. Isothermal fatigue behavior
Fig. 4. (a) Stress–strain diagram for Ti-6242 and the TMC; strain rate was 10−4 s−1 . (TMC: tested to fracture, Ti-6242: test interrupted at 6% and 10% strain, respectively.) (b) Temperature dependence of matrix strength; strain rate was 10−3 s−1 .
6242 matrix creeps considerably at 550 ◦ C. Strain increases with increasing applied stress (80–200 MPa) at a constant temperature of T = 550 ◦ C, and with increasing temperature (from 450 to 550 ◦ C) at a constant stress of 150 MPa. Fig. 5a shows that the creep rate increases accordingly. Raising the stress from 100 to 200 MPa increases the strain rate by one order of magnitude. A fiber reinforcement improves the creep resistance considerably. The creep strain is much smaller, although the applied stress is higher by a factor of 10 (Fig. 5b). The stress dependence of the minimum creep rate was calculated for the matrix material at 550 ◦ C using the well known
The results of the isothermal fatigue tests at room temperature and at 550 ◦ C are summarized in Fig. 6. The stress values have been corrected according to the actual fiber volume fraction. In the LCF-regime, i.e. at Nf < 104 (Nf represents the number of cycles to fracture), the fatigue strength at room temperature is higher than that at 550 ◦ C. There is, however, a cross-over between 104 and 105 cycles, and when exceeding this point, the fatigue strength is about 350 MPa higher at 550 ◦ C than at room temperature. The fatigue limit at 550 ◦ C is about 950 MPa, while at room temperature a fatigue limit was determined to be about 600 MPa. Similar results have been reported for other TMCs [18,19]. At both temperatures the fatigue failure mode of the TMC changes from fiber dominated failure at high loads to propagation of one or more matrix cracks at lower loads, while fibers are bridging the matrix crack until the specimen fails. The matrix crack can cover a substantial area of the cross-section, as e.g. demonstrated in Fig. 7. The transition from fatigue to catastrophic crack propagation, Fig. 8a, is associated with a change in micro-roughness of the matrix surface, which drops to about half, Fig. 8b. 3.4. Thermo-mechanical fatigue behavior
Table 1 Tensile test results of the unreinforced matrix and the composite Material
Strain rate (s−1 )
T (◦ C)
Rp0.2 (MPa)
Rm (MPa)
Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 TMC TMC
10−3 10−3 10−3 10−3 10−3 10−4 10−4 10−4 10−4
RT 300 400 500 550 RT 550 RT 550
1163 818 767 665 523 1138 255 1907 –
1166 871 816 727 654 1153 372
Rmax (MPa)
2141 1518
The results of the TMF tests on the composite material are shown in Fig. 6 as well. Under TMF in-phase conditions, small changes in the applied stress cause large variation in fatigue life, e.g., an increase of stress from 1100 to 1300 MPa results in a reduction of fatigue life by a factor of about 500. At a stress level of about 1100 MPa the fatigue life in IP and OP tests is almost the same. On examining the broken specimens, no cracks are observed on the outer surface of the specimens under IP TMF loading conditions. Thus, there is no evidence that cracks initiated at the surface before final failure of the specimens occurred. The fracture
540
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 5. (a) Creep rate vs. creep strain for matrix material. (b) Creep rate vs. creep strain for TMC. (c) Results of creep tests as applied creep stress vs. time to fracture for different TMCs at 550 and 600 ◦ C, respectively.
surfaces are similar to those seen after quasi-static tensile testing. Under out-of-phase TMF conditions, there is a change in the fracture mode. At stresses higher than 1800 MPa, the fatigue life curve is almost parallel to that of in-phase tests. It is, however, shifted to much higher stresses, i.e. lifetime under high stresses is longer than under in-phase TMF conditions. At stresses lower than 1800 MPa, the fatigue life curve deviates and shows a sharp decline. Small changes in applied stress now have a reduced effect on fatigue life than under in-phase conditions. This behavior is associated with the observation of multiple matrix cracks
Fig. 6. Maximum applied stress vs. number of cycles to fracture for TMC loaded under LCF conditions at RT and 550 ◦ C, and under TMF conditions between 550 and 100 ◦ C.
Fig. 7. Fracture surface of TMC specimen after isothermal fatigue (LCF) at room temperature, σ max = 700 MPa, and Nf = 4.6 × 106 (SEM, SE). White line marks the transition from fatigue to catastrophic failure (bar = 2 mm).
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
541
Fig. 9. (a) Development of crack density with increasing distance from fracture surface. Corresponding surfaces of test pieces after testing at 1700 MPa (b), 1600 MPa (c), and 1400 MPa (d).
Fig. 8. (a) Transition from stable to unstable crack growth at higher magnification. (b) Confocal light scanning microscopy image of a fracture surface after LCF testing.
initiated at the outer surface of the specimens, see e.g. Fig. 9b–d. Fig. 9a shows the crack density versus distance from the final fracture location after OP-TMF-testing in a stress range between 1400 and 1700 MPa. The crack density is high if high stresses are applied. The crack density decreases with increasing distance from the fracture surface. Only a small number of cracks is observed for = 1400 MPa. The specimens fail when one or more cracks propagating into the specimen reach the first or second row of fibers. The fibers degrade at 550 ◦ C due to oxidation [17], overloading the remaining fibers and leading to specimen failure. Matrix cracks can be easily detected by oxidation colors, Fig. 10. At the tip of those cracks that have reached the fibers, striations can be found followed by smooth regions similar to those known from stretched zones in steels. At a maximum stress of 1400 MPa and lower, fiber bridging is observed, Fig. 10b. To study the effect of oxidation on TMF-life, two preoxidized (300 h at 550 ◦ C in air) TMC specimens were tested under out-of-phase conditions and at a maximum stress of 1600 MPa. This results in a shortening of lifetime by a factor of two to three (see Fig. 6).
4. Discussion Due to the processing technique, the composite has a much finer grain size than the titanium alloy rod material. Fig. 1 shows that the grain size is smaller by a factor of 4. This has to be taken into account when describing the deformation behavior involving creep and relaxation. If a composite is subjected to a constant load in a creep test, then stresses within the matrix relax and, due to the balanced total stress, the fiber stresses increase, which may cause overloading of the fibers and subsequently statistically distributed fiber failure and finally failure of the composite. This has been shown in Ref. [20]. Sinclair et al. [21] proved experimentally that intact fibers have high peak strains in the vicinity of broken fibers as compared to those which have no broken fibers in their neighborhood. Other investigations support the view that fiber failure is the dominating mechanism controlling creep failure [19,22]. Fibers fail under a lower stress under constant load conditions at high temperature than in a quasi-static tensile test [19], which has been attributed to the growth of fiber defects. The isothermal fatigue limit of the Ti-6242/SCS-6 TMC at high temperature determined in this work is much higher than at room temperature. A similar behavior has been reported for other TMCs [23]. This is attributed to the relaxation of internal stresses in the matrix near to the interface to the fibers at elevated temperatures. This has to be taken into account when describing the damage mechanisms because it results in a cyclic
542
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 10. Fracture surface of a TMC specimen after OP-TMF testing (SEM, SE).
tension–compression state of the matrix, although the composite cycles in tension–tension [5]. This state is reached already after a small number of cycles if the frequency is low. It has been shown that this causes a suppression of matrix crack formation, leading to a high endurance limit [24]. Tanaka et al. [25] have demonstrated using in situ fatigue tests in SEM that the fatigue crack growth rate at 550 ◦ C was lower than that at room temperature when cycling at R = 0.1. They explained their observations with matrix residual compressive stresses at the crack front due to asymmetry in matrix tensile and compressive creep at elevated temperature during each cycle. This effect combined with crack closure at crack wake leads to a decrease in crack tip stress intensity factor and consequently in a lower fatigue crack growth rate at high temperature. Stress relaxation tests on titanium alloys [20] showed that Ti-6242 is sensitive to creep at high temperatures so that internal stresses in the TMC specimen are reduced quickly while being subjected to an external tensile stress at high temperature. However, during room temperature isothermal fatigue tests and during out-of-phase TMF tests, the matrix is subjected to the applied stresses and also to the internal stresses in an accumulative way, which results in smaller life times. The internal tensile stress in the matrix near to the interface to the fibers is introduced during the cooling phase after the HIP process of the TMC specimens due to the mismatch of the
coefficients of thermal expansion (CTE) of matrix and fiber (␣SiC-fiber = 4.21 × 10−6 K−1 , ␣Ti-matrix = 10.47 × 10−6 K−1 ). Bobet et al. [26] calculated the thermal residual stresses in a SiC/Ti-15-3 TMC and showed that after cooling down from processing temperature, the matrix is under tensile stress, while the fiber is under compression. The calculated results are strongly affected by the uncertainty of the CTE of the matrix. They also calculated the evolution of the stresses during fatigue testing and showed that the thermal residual stresses relax and that the relaxation increases with increasing number of cycles. The axial stresses relaxed most quickly and after about 100 cycles, the thermal stresses were reduced to half of the initial value. After about 1000 cycles, the thermal residual stress reached a saturation at about one third of the value in the initial state. Stress relaxation also plays a role during in-phase TMF loading. Calculations have shown that matrix stresses relax already during the first cycles while, simultaneously, stresses in the fibers and calculated strain increase [27]. The damage during in-phase TMF test is mainly creep controlled and failure is dominated by fiber failure. Matrix damage plays a minor role because the strain range in the matrix is small [28]. Since there are no surface cracks, oxidation does not play a decisive role. The oxidation speed into the crack free specimen has been proven to be low at 550 ◦ C [29]. Therefore, the limiting factor for IP-TMF strength is the fiber strength.
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Out-of-phase TMF, however, is controlled by matrix crack formation and growth. These processes are enhanced by oxidation and oxygen uptake in the surface region, which mainly occurs during the dwell time at 550 ◦ C (and low stress). Under these conditions, oxidation at high temperature and high tensile stress at low temperature alternate. The applied tensile stress adds to the residual stress in the matrix. Jin and Johnson have shown that dwell times at 500 and 650 ◦ C reduce the composite life due to enhanced oxidation processes, i.e. thicker oxide scale [30]. Dwell times are well known to have a harmful effect on fatigue life in ␣ and ␣/ titanium alloys. McBagonluri et al. [31] studied dwell fatigue crack growth in Ti-6242 at room temperature and found that in the short crack growth regime, crack growth rates are much greater under dwell fatigue than under pure fatigue conditions. The authors attribute this effect to possible creep effects, which give rise to a mean stress effect during dwell fatigue. Stress relaxation and creep processes play a minor role under out-of-phase TMF conditions because high stresses are applied at low temperature. Further, as a consequence of the surface oxidation, the specimen’s surface embrittles and numerous cracks are initiated on the surface. It is reasonable to assume that the interaction of stressing the cracked regions and oxidizing the crack tips is responsible for the crack propagation mechanism. The cracks grow into the unreinforced outer rim of the test piece and eventually reach the first row of fibers. Then oxidation becomes even more critical because fiber oxidation severely affects the fiber strength. A previous work has shown that oxidation of SCS6 fibers at 600 ◦ C in air results in weight loss and reduced carbon coating thickness [17]. After 4 h of oxidation the strength is reduced from 4610 MPa to 3001 MPa which has been attributed to an embrittlement of the carbon coating on the fiber giving rise to early coating failure and formation of stress intensity. Therefore, the drop in TMF strength under OP conditions and Nf > 500 is ascribed to the oxidation attack on the fibers. The high stress range results in high strains, which contribute to crack formation in the matrix. Hence, matrix cracks are initiated already after a small number of cycles. 5. Conclusions A Ti-6242/SiC composite was characterized with respect to mechanical properties under creep, isothermal and thermomechanical fatigue conditions. The main results were: (1) SiC fiber reinforcement improves the creep resistance considerably at 550 ◦ C. Failure is dominated by fiber failure. (2) The isothermal fatigue limit of the composite at 550 ◦ C is considerably higher than at room temperature. Residual stresses due to the mismatch of the coefficient of thermal expansion of matrix and reinforcement may relax at high temperatures resulting in a cyclic tension–compression stress state in the matrix and, consequently, a suppression of matrix crack formation. (3) In contrast, applied stresses and internal stresses accumulate in the matrix during isothermal room temperature fatigue tests and OP-TMF tests.
543
(4) Under TMF conditions, life time of the composite is higher under OP- than IP-conditions if Nf < 1000. Life time is controlled by fiber strength under IP-conditions. Residual stresses in the matrix relax during the first cycles while fiber stresses increase. (5) Under OP-TMF conditions we find a change in behavior at about 1800 MPa. Above 1800 MPa, the curve is shifted to higher stresses parallel to IP-TMF. Below 1800 MPa, multiple cracks are found, which proceed to the first row of fibers. This results in oxidation of the SiC fibers, which causes a reduction in fiber strength and consequently in fiber failure. Acknowledgements Financial Support by the “Deutsche Forschungsgemeinschaft” (DFG Po 405/5-1 and 405/5-3) is gratefully acknowledged. The technical support of BAM Division V.1 (Microstructure of Engineering Materials) is gratefully appreciated. We also wish to thank Dr. Goebbels of BAM Division VIII.3 (Radiological Methods) for CT-measurements.
References [1] I.J. Polmear, Light Alloys—Metallurgy of the Light Metals, 3rd ed., Hodder Headline, PLC, London, 1996, p. 261. [2] C. Leyens, J. Hausmann, J. Kumpfert, in: M. Peters, C. Leyens (Eds.), Titan und Titanlegierungen, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim, 2002, pp. 321–350. [3] G.T. Ward, J. Herrmann, M. Hillberry, J. Compos. Tech. Res. 17 (1995) 205–211. [4] T.P. Gabb, J. Gayda, A. MacKay, J. Compos. Mater. 24 (1990) 667–686. [5] B.P. Sander, S. Mall, J. Compos. Tech. Res. 16 (1994) 304–313. [6] R. Talreja, Mater. Sci. Engg. A200 (1995) 21–28. [7] T.P. Gabb, J. Gayda, P.A. Bartolotta, M.G. Castelli, Int. J. Fatigue 15 (1993) 413–422. [8] M.G. Castelli, in: AGARD Conf. Proc. 569, 2.-4.10.1995, Banff, Canada, PP. 12–1. [9] M.G. Castelli, J. Gayda, in: Proc. Conf. on Reliability, stress analysis and failure prevention, 1993, Albuquerque, New Mexico, USA, 19.-22. Sept. 1993, DE-Vol. 55, ASME 1993, 213. [10] M.G. Castelli, M.G. Life Prediction Methodology for Titanium Matrix Composites, ASTM STP 1253, American Soc. for Testing and Materials, 1996, 412–431. [11] R. Neu, T. Nicholas, J. Compos. Tech. Res. 16 (1994) 214–224. [12] T. Nicholas, Mater. Sci. Eng. A200 (1995) 29–37. [13] L. Figel, B. G¨unther, Modeling the high temperature longitudinal fatigue behavior of metal matrix composites (SiC/Ti-6242): non-linear time dependent matrix behavior, Int. J. Fatigue, submitted for publication. [14] R.R. Boyer, Adv. Perform. Mater. 2 (1995) 349. [15] R. Leucht, H.J. Dudek, K. Weber, A. Werner, Z. Metallkd. 87 (1996) 424. [16] E. Olevsky, H.J. Dudek, W.A. Kaysser, Acta Mater. 44 (1996) 715. [17] P.W.M. Peters, J. Hemptenmacher, Composites, Part A 33 (2002) 1373–1379. [18] M.P. Thomas, G. Howard, R. Pather, M.R. Brown, P.H. Tranter, Int. J. Fatigue 23 (2001) 851–856. [19] P.W.M. Peters, J. Hempenmacher, K. Weber, H. Assler, J. Comp. Techn. Res. 24 (2002) 246–253. [20] J. Hemptenmacher, P.W.M. Peters, K. Weber, Adv. Eng. Mater. 6 (2004) 129–133. [21] R. Sinclair, M. Preuss, E. Maire, J.-Y. Buffiere, P. Brown, P.J. Withers, Acta Mater. 52 (2004) 1423–1438. [22] S.W. Schwenker, D. Eylon, Metall. Mater. Trans. A 27 (1996) 4193–4204.
544
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
[23] P.W.M. Peters, H. Assler, J. Hemptenmacher, Z. Xia, Proc. ECCM-9, Brighton, June 4–7, 2000. [24] P.W.M. Peters, W.D. Zeng, J. Hemptenmacher, Z. Xia, Comp. Mater., Test. Design, ASTM STP 1436 (2003) 273–288. [25] Y. Tanaka, Z.-Y. Deng, Y.-F. Liu, C. Masuda, Acta Mater. 51 (2003) 6329–6340. [26] J.-L. Bobet, C. Masuda, Y. Kagawa, J. Mater. Sci. 32 (1997) 6357–6369.
[27] P.W.M. Peters, J. Hemptenmacher, B. G¨unther, D. Bettge, P. Portella, in: Proc. ECCM-11, Rhodos, May 31–June 3, 2004. [28] R. Neu, Fatigue Fract. Eng. M 16 (1993) 811–828. [29] C. Todd, DLR Internal Report IB 334-02/06, 2002. [30] O. Jin, W.S. Johnson, Composites A35 (2004) 113–120. [31] F. McBagonluri, E. Akpan, C. Mercer, W. Shen, W.O. Soboyejo, Mater. Sci. Engg. A405 (2005) 111–134.
Mechanical behavior and fatigue damage of a titanium matrix composite reinforced with continuous SiC fibers D. Bettge a , B. G¨unther a , W. Wedell a , P.D. Portella a , J. Hemptenmacher b , P.W.M. Peters b , B. Skrotzki a,∗ a
Federal Institute for Materials Research and Testing (BAM), 12200 Berlin, Germany b German Aerospace Center (DLR), 51147 Cologne, Germany
Received 31 July 2006; received in revised form 18 October 2006; accepted 20 October 2006
Abstract A titanium (Ti-6242) matrix composite reinforced with continuous SiC fibers was studied. The mechanical behavior of the matrix and the composite was characterized by tensile, creep and isothermal fatigue tests at room temperature and up to 550 ◦ C. The thermo-mechanical fatigue behavior under in-phase and out-of-phase conditions was investigated for the composite between 100 and 550 ◦ C. Fracture surfaces were characterized by confocal light microscopy and by scanning electron microscopy to identify the damage mechanisms. © 2006 Elsevier B.V. All rights reserved. Keywords: Metal matrix composites; Tensile strength; Creep strength; Isothermal fatigue; Thermo-mechanical fatigue; Damage analysis
1. Introduction Near-␣ titanium alloys are classical high temperature materials for applications in the temperature range between 500 and 550 ◦ C. They were developed to meet demands for higher operating temperatures in the compressor section of aircraft engines [1]. These alloys combine the good creep resistance of ␣-alloys with the high strength of (␣ + )-alloys. During the last few years, new engine concepts were developed with higher power densities, which increased the demands with respect to mechanical and thermal loads even further [2]. At the same time, there is the need for consistent light weight design which requires essentially new material concepts. Intermetallic titanium aluminides and reinforced titanium matrix composites (TMCs) are examples for such developments. Reinforcing titanium alloys with continuous SiC fibers for applications in aero engines represents a class of materials, which combines the high strength, stiffness and creep resistance of the ceramic SiC fiber with the damage tolerance of titanium alloys. The use of the light ceramic fiber (ρ = 3.0 g/cm3 , E = 380 GPa) in the titanium alloy matrix reduces the density of the latter even further and provides the potential for consider-
∗
Corresponding author. Tel.: +49 30 8104 1520; fax: +49 30 8104 1527. E-mail address: [email protected] (B. Skrotzki).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.10.107
able weight savings in aero engines when replacing bulk titanium alloy parts. Potential applications are redesigned discs and shafts of the compressor section of the engines. Disks are supposed to become rings and eventually monolithic bladed rings (“blings”). Presently, a major obstacle on the way to applications for safety relevant components made of TMCs is the complex fatigue damage mechanism. In the last decade, the isothermal fatigue behavior of TMCs has been investigated at room temperature [3] and at elevated temperature [4,5]. It is evident that the room temperature fatigue behavior of long fiber reinforced TMCs can be interpreted using general concepts for metal matrix composites [6]. The damage mechanisms commonly are fiber failure (at high stresses and low cycles), matrix cracking (at intermediate stresses and intermediate cycles), and no damage or damage not giving rise to final failure (at low stresses and high cycles). At high temperatures, the fatigue behavior generally can be described similarly. Realistic testing of TMCs for aircraft engine applications, however, incorporates the variation of load as well as of temperature, i.e. thermo-mechanical loading. Thermo-mechanical fatigue (TMF) tests are commonly used to characterize the simultaneous influence of load and temperature on fatigue life of TMCs. Some extensive studies on the behavior of titanium matrix composites containing SiC fibers under TMF loading were published (e.g. [7–12]). The fatigue life of TMCs strongly depends on the experimental details of the TMC tests. Hold
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
times have a major influence on the damage mechanisms because the matrix alloy becomes very soft at elevated temperatures as opposed to the SiC fibers. It was found that out-of-phase (maximum load applied at minimum temperature) thermo-mechanical loading generally leads to a stronger reduction in fatigue life than in-phase (maximum load applied at maximum temperature) loading [9,10]. This seems to be controlled by matrix cracking whereas failure during in-phase cycling appears to be dominated by fiber failure [11]. In the present work, results of uniaxial tensile and creep tests as well as of isothermal and thermo-mechanical fatigue tests are presented for the material system Ti-6242/SiC. The mechanical data were determined to identify the material parameters of a viscoplastic material model, which will be used to simulate the material behavior of the composite. These results will be published elsewhere, first data are given in Ref. [13]. 2. Experimental procedure The alloy Ti-6242 was used as matrix material. This alloy is mainly employed in gas turbines for rotating components. Its nominal composition is Ti–6%Al–2%Sn–4%Zr–2%Mo– 0.1%Si (wt.%). Silicon is added for improved creep resistance [1]. The alloy has a near-␣ microstructure and finds its applications in rotating components at temperatures up to 540 ◦ C
537
[14]. Ti-6242 is widely used in the aerospace industry and its behavior is well characterized. Therefore, it was chosen as a suitable matrix material for this study. The unreinforced matrix was available as thermo-mechanically processed rod material with an ␣-grain size of about 4 m, Fig. 1a. The SiC fiber SCS-6 (Textron Specialty Materials, Lowell, MA, USA) was used as long fiber reinforcement. The fiber has an outer diameter of 142 m and a carbon core of 30 m diameter. The fiber is covered with a protection layer, which consists mainly of carbon. The composite material was processed by deposition of the matrix alloy on the SCS-6 fibers using the magnetron sputtering process [15]. Subsequently, the coated fibers are cut to the required length and stacked into a tube of Ti-6242. The preform is hipped at 950 ◦ C and 190 MPa [16]. The fiber volume fraction in the gauge length was measured to be about 35%. The grain size of the sputtered matrix material was about 1 m, Fig. 1b–d. Cylindrical specimens were machined with an overall length of 81 mm, a gauge length of 21 mm (unreinforced matrix) and 10 mm (TMC), and a gauge diameter of 3.5 mm. The transition to the threaded heads was machined to a parabolic shape. The surface of the gauge length was mechanically polished (diamond suspension, min. size 3 m). The spatial distribution of fibers in the gauge length visualized by X-ray computer tomography (CT) is shown in Fig. 2 in a cross-section. From the CT-images,
Fig. 1. (a) Ti-6242 matrix consisting of dark ␣-grains and bright -grains (SEM, BSE-mode). (b) SEM-image of the TMC; overview at low magnification. (c) TMC at higher magnification. (d) TMC at higher magnification. Grain size is smaller than in the unreinforced matrix due to sputtering.
538
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 2. Low resolution CT-image of the TMC-test piece cross-section in the gauge length.
the actual fiber volume fraction was determined for each test piece. The yield strength and ultimate tensile strength were determined in tensile tests at room temperature and at 550 ◦ C. Constant strain rates of 10−3 and 10−4 s−1 were applied. Yielding was determined at 0.2% offset. Most of the tests on the matrix material were interrupted after reaching 10% strain. All fatigue tests were accomplished on a 100 kN MTS servohydraulic testing system equipped with induction heating. For LCF and TMF testing, the specimen temperature was measured using two ribbon thermo-couples (Pt-RhPt type S) at the upper and lower end of the gauge section. The temperature gradient in the gauge length was less than ±15 ◦ C. All fatigue tests were carried out in air under load control with a stress ratio of R = 0.1. The reason to run tests under load control is (i) that the ceramic fibers fail already at small strains and (ii) that a compressive stress state had to be avoided because the specimen tend to buckle due to the small diameter in the gauge length. The frequency of the isothermal tests was 0.25 Hz for an estimated lifetime of up to 105 cycles, 1 Hz for up to 106 cycles and 5 Hz exceeding 106 cycles. For TMF-tests, the temperature was cycled with the same frequency as the stress, Fig. 3, with a temperature rate of 5 K/s and an upper and lower temperature of 550 and 100 ◦ C, respectively, resulting in a heating and cooling time intervals of 1.5 min each. After each heating and cooling ramp, the temperature was held for 2 min. Consequently, a full TMF cycle lasted 7 min. The applied stress was varied in-phase (IP) or out-of-phase (OP) to temperature with R = 0.1. A high temperature extensometer with a gauge length of 10 mm and a precision of 1% was used. Uniaxial creep tests were carried out under constant tensile load on the matrix material at initial stresses of 80–200 MPa and at a temperature of 550 ◦ C. One test each was run at 450 and 500 ◦ C at σ = 150 MPa. The TMC was tested at 550 ◦ C and at stresses varying between 1100 and 1400 MPa. The geometry of the creep specimen was the same as for the fatigue tests. The strain measurement was performed using a high temperature extensometer with 21.5 mm gauge length. The creep strain was measured, i.e. the time dependent strain that occurs after the
Fig. 3. Schematic representation of the temperature-time- and stress-time-path of the TMF-tests.
complete application of the load. Three Pt-RhPt (type S) thermocouples were connected to the upper and lower end and in the center of the gauge length. A standard three zone cylindrical radiation furnace was used for heating. The temperature gradient within the gauge length was less than ±3 ◦ C. A preload of 10% of the total load was applied. The composite material and the fracture surfaces were characterized using conventional light microscopy (LM), a confocal light scanning microscope (CLSM) Leica DSCM1000, scanning electron microscopes (SEM) CamScan 2 and LEO Gemini 1530 VP. Low resolution computer tomography has been applied to characterize each individual test piece with respect to the actual number of fibers. 3. Results 3.1. Tensile tests Tensile properties were determined for the matrix and for the composite material at room temperature and at 550 ◦ C, as given in Fig. 4. Ti-6242 has a yield strength Rp0.2 of 1163 MPa at room temperature (strain rate 10−3 s−1 ) which is considerably reduced to 523 MPa at 550 ◦ C. Applying a lower strain rate of 10−4 s−1 reduces the strength at RT somewhat, while the strength at 550 ◦ C is decreased by 50% due to creep effects. The fiber reinforcement leads to a clear improvement of the mechanical behavior at both test temperatures (Rp0.2 = 1907 MPa at room temperature, and Rmax = 1518 MPa at 550 ◦ C). The temperature dependence of the tensile behavior of the metallic matrix is shown in Fig. 4b. All strength data are summarized in Table 1. The mean strength of a single SiC fiber at room temperature is about 4600 MPa which drops only a little if temperature is increased up to 550 ◦ C [17]. 3.2. Creep tests The matrix material and the composite were tested in shorttime creep tests. Creep strain ε versus time t was measured. The creep data were then further analyzed and represented as creep rate ε˙ versus creep strain ε in Fig. 5a and b. The Ti-
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
phenomenological equation Qc σn, ε˙ = A exp − RT
539
(1)
where A represents a constant, Qc is the apparent activation energy for creep, R the universal gas constant, T the temperature, σ the creep stress and n is the stress exponent. Analyzing the plot of log creep rate versus log stress (not shown) provides the stress exponent n which was determined to be 3.2 for the matrix at 550 ◦ C. It was not possible to determine the stress exponent for the TMC because the tests were interrupted before the minimum creep rate was reached. Fig. 5c shows results of the short-time creep tests on TMC material in terms of applied stress versus time to fracture. Results obtained in this work are compared to those published in the literature. The figure shows that due to the high homogeneity with respect to fiber distribution and matrix microstructure, the material produced by DLR shows very high creep strength. 3.3. Isothermal fatigue behavior
Fig. 4. (a) Stress–strain diagram for Ti-6242 and the TMC; strain rate was 10−4 s−1 . (TMC: tested to fracture, Ti-6242: test interrupted at 6% and 10% strain, respectively.) (b) Temperature dependence of matrix strength; strain rate was 10−3 s−1 .
6242 matrix creeps considerably at 550 ◦ C. Strain increases with increasing applied stress (80–200 MPa) at a constant temperature of T = 550 ◦ C, and with increasing temperature (from 450 to 550 ◦ C) at a constant stress of 150 MPa. Fig. 5a shows that the creep rate increases accordingly. Raising the stress from 100 to 200 MPa increases the strain rate by one order of magnitude. A fiber reinforcement improves the creep resistance considerably. The creep strain is much smaller, although the applied stress is higher by a factor of 10 (Fig. 5b). The stress dependence of the minimum creep rate was calculated for the matrix material at 550 ◦ C using the well known
The results of the isothermal fatigue tests at room temperature and at 550 ◦ C are summarized in Fig. 6. The stress values have been corrected according to the actual fiber volume fraction. In the LCF-regime, i.e. at Nf < 104 (Nf represents the number of cycles to fracture), the fatigue strength at room temperature is higher than that at 550 ◦ C. There is, however, a cross-over between 104 and 105 cycles, and when exceeding this point, the fatigue strength is about 350 MPa higher at 550 ◦ C than at room temperature. The fatigue limit at 550 ◦ C is about 950 MPa, while at room temperature a fatigue limit was determined to be about 600 MPa. Similar results have been reported for other TMCs [18,19]. At both temperatures the fatigue failure mode of the TMC changes from fiber dominated failure at high loads to propagation of one or more matrix cracks at lower loads, while fibers are bridging the matrix crack until the specimen fails. The matrix crack can cover a substantial area of the cross-section, as e.g. demonstrated in Fig. 7. The transition from fatigue to catastrophic crack propagation, Fig. 8a, is associated with a change in micro-roughness of the matrix surface, which drops to about half, Fig. 8b. 3.4. Thermo-mechanical fatigue behavior
Table 1 Tensile test results of the unreinforced matrix and the composite Material
Strain rate (s−1 )
T (◦ C)
Rp0.2 (MPa)
Rm (MPa)
Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 Ti-6242 TMC TMC
10−3 10−3 10−3 10−3 10−3 10−4 10−4 10−4 10−4
RT 300 400 500 550 RT 550 RT 550
1163 818 767 665 523 1138 255 1907 –
1166 871 816 727 654 1153 372
Rmax (MPa)
2141 1518
The results of the TMF tests on the composite material are shown in Fig. 6 as well. Under TMF in-phase conditions, small changes in the applied stress cause large variation in fatigue life, e.g., an increase of stress from 1100 to 1300 MPa results in a reduction of fatigue life by a factor of about 500. At a stress level of about 1100 MPa the fatigue life in IP and OP tests is almost the same. On examining the broken specimens, no cracks are observed on the outer surface of the specimens under IP TMF loading conditions. Thus, there is no evidence that cracks initiated at the surface before final failure of the specimens occurred. The fracture
540
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 5. (a) Creep rate vs. creep strain for matrix material. (b) Creep rate vs. creep strain for TMC. (c) Results of creep tests as applied creep stress vs. time to fracture for different TMCs at 550 and 600 ◦ C, respectively.
surfaces are similar to those seen after quasi-static tensile testing. Under out-of-phase TMF conditions, there is a change in the fracture mode. At stresses higher than 1800 MPa, the fatigue life curve is almost parallel to that of in-phase tests. It is, however, shifted to much higher stresses, i.e. lifetime under high stresses is longer than under in-phase TMF conditions. At stresses lower than 1800 MPa, the fatigue life curve deviates and shows a sharp decline. Small changes in applied stress now have a reduced effect on fatigue life than under in-phase conditions. This behavior is associated with the observation of multiple matrix cracks
Fig. 6. Maximum applied stress vs. number of cycles to fracture for TMC loaded under LCF conditions at RT and 550 ◦ C, and under TMF conditions between 550 and 100 ◦ C.
Fig. 7. Fracture surface of TMC specimen after isothermal fatigue (LCF) at room temperature, σ max = 700 MPa, and Nf = 4.6 × 106 (SEM, SE). White line marks the transition from fatigue to catastrophic failure (bar = 2 mm).
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
541
Fig. 9. (a) Development of crack density with increasing distance from fracture surface. Corresponding surfaces of test pieces after testing at 1700 MPa (b), 1600 MPa (c), and 1400 MPa (d).
Fig. 8. (a) Transition from stable to unstable crack growth at higher magnification. (b) Confocal light scanning microscopy image of a fracture surface after LCF testing.
initiated at the outer surface of the specimens, see e.g. Fig. 9b–d. Fig. 9a shows the crack density versus distance from the final fracture location after OP-TMF-testing in a stress range between 1400 and 1700 MPa. The crack density is high if high stresses are applied. The crack density decreases with increasing distance from the fracture surface. Only a small number of cracks is observed for = 1400 MPa. The specimens fail when one or more cracks propagating into the specimen reach the first or second row of fibers. The fibers degrade at 550 ◦ C due to oxidation [17], overloading the remaining fibers and leading to specimen failure. Matrix cracks can be easily detected by oxidation colors, Fig. 10. At the tip of those cracks that have reached the fibers, striations can be found followed by smooth regions similar to those known from stretched zones in steels. At a maximum stress of 1400 MPa and lower, fiber bridging is observed, Fig. 10b. To study the effect of oxidation on TMF-life, two preoxidized (300 h at 550 ◦ C in air) TMC specimens were tested under out-of-phase conditions and at a maximum stress of 1600 MPa. This results in a shortening of lifetime by a factor of two to three (see Fig. 6).
4. Discussion Due to the processing technique, the composite has a much finer grain size than the titanium alloy rod material. Fig. 1 shows that the grain size is smaller by a factor of 4. This has to be taken into account when describing the deformation behavior involving creep and relaxation. If a composite is subjected to a constant load in a creep test, then stresses within the matrix relax and, due to the balanced total stress, the fiber stresses increase, which may cause overloading of the fibers and subsequently statistically distributed fiber failure and finally failure of the composite. This has been shown in Ref. [20]. Sinclair et al. [21] proved experimentally that intact fibers have high peak strains in the vicinity of broken fibers as compared to those which have no broken fibers in their neighborhood. Other investigations support the view that fiber failure is the dominating mechanism controlling creep failure [19,22]. Fibers fail under a lower stress under constant load conditions at high temperature than in a quasi-static tensile test [19], which has been attributed to the growth of fiber defects. The isothermal fatigue limit of the Ti-6242/SCS-6 TMC at high temperature determined in this work is much higher than at room temperature. A similar behavior has been reported for other TMCs [23]. This is attributed to the relaxation of internal stresses in the matrix near to the interface to the fibers at elevated temperatures. This has to be taken into account when describing the damage mechanisms because it results in a cyclic
542
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Fig. 10. Fracture surface of a TMC specimen after OP-TMF testing (SEM, SE).
tension–compression state of the matrix, although the composite cycles in tension–tension [5]. This state is reached already after a small number of cycles if the frequency is low. It has been shown that this causes a suppression of matrix crack formation, leading to a high endurance limit [24]. Tanaka et al. [25] have demonstrated using in situ fatigue tests in SEM that the fatigue crack growth rate at 550 ◦ C was lower than that at room temperature when cycling at R = 0.1. They explained their observations with matrix residual compressive stresses at the crack front due to asymmetry in matrix tensile and compressive creep at elevated temperature during each cycle. This effect combined with crack closure at crack wake leads to a decrease in crack tip stress intensity factor and consequently in a lower fatigue crack growth rate at high temperature. Stress relaxation tests on titanium alloys [20] showed that Ti-6242 is sensitive to creep at high temperatures so that internal stresses in the TMC specimen are reduced quickly while being subjected to an external tensile stress at high temperature. However, during room temperature isothermal fatigue tests and during out-of-phase TMF tests, the matrix is subjected to the applied stresses and also to the internal stresses in an accumulative way, which results in smaller life times. The internal tensile stress in the matrix near to the interface to the fibers is introduced during the cooling phase after the HIP process of the TMC specimens due to the mismatch of the
coefficients of thermal expansion (CTE) of matrix and fiber (␣SiC-fiber = 4.21 × 10−6 K−1 , ␣Ti-matrix = 10.47 × 10−6 K−1 ). Bobet et al. [26] calculated the thermal residual stresses in a SiC/Ti-15-3 TMC and showed that after cooling down from processing temperature, the matrix is under tensile stress, while the fiber is under compression. The calculated results are strongly affected by the uncertainty of the CTE of the matrix. They also calculated the evolution of the stresses during fatigue testing and showed that the thermal residual stresses relax and that the relaxation increases with increasing number of cycles. The axial stresses relaxed most quickly and after about 100 cycles, the thermal stresses were reduced to half of the initial value. After about 1000 cycles, the thermal residual stress reached a saturation at about one third of the value in the initial state. Stress relaxation also plays a role during in-phase TMF loading. Calculations have shown that matrix stresses relax already during the first cycles while, simultaneously, stresses in the fibers and calculated strain increase [27]. The damage during in-phase TMF test is mainly creep controlled and failure is dominated by fiber failure. Matrix damage plays a minor role because the strain range in the matrix is small [28]. Since there are no surface cracks, oxidation does not play a decisive role. The oxidation speed into the crack free specimen has been proven to be low at 550 ◦ C [29]. Therefore, the limiting factor for IP-TMF strength is the fiber strength.
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
Out-of-phase TMF, however, is controlled by matrix crack formation and growth. These processes are enhanced by oxidation and oxygen uptake in the surface region, which mainly occurs during the dwell time at 550 ◦ C (and low stress). Under these conditions, oxidation at high temperature and high tensile stress at low temperature alternate. The applied tensile stress adds to the residual stress in the matrix. Jin and Johnson have shown that dwell times at 500 and 650 ◦ C reduce the composite life due to enhanced oxidation processes, i.e. thicker oxide scale [30]. Dwell times are well known to have a harmful effect on fatigue life in ␣ and ␣/ titanium alloys. McBagonluri et al. [31] studied dwell fatigue crack growth in Ti-6242 at room temperature and found that in the short crack growth regime, crack growth rates are much greater under dwell fatigue than under pure fatigue conditions. The authors attribute this effect to possible creep effects, which give rise to a mean stress effect during dwell fatigue. Stress relaxation and creep processes play a minor role under out-of-phase TMF conditions because high stresses are applied at low temperature. Further, as a consequence of the surface oxidation, the specimen’s surface embrittles and numerous cracks are initiated on the surface. It is reasonable to assume that the interaction of stressing the cracked regions and oxidizing the crack tips is responsible for the crack propagation mechanism. The cracks grow into the unreinforced outer rim of the test piece and eventually reach the first row of fibers. Then oxidation becomes even more critical because fiber oxidation severely affects the fiber strength. A previous work has shown that oxidation of SCS6 fibers at 600 ◦ C in air results in weight loss and reduced carbon coating thickness [17]. After 4 h of oxidation the strength is reduced from 4610 MPa to 3001 MPa which has been attributed to an embrittlement of the carbon coating on the fiber giving rise to early coating failure and formation of stress intensity. Therefore, the drop in TMF strength under OP conditions and Nf > 500 is ascribed to the oxidation attack on the fibers. The high stress range results in high strains, which contribute to crack formation in the matrix. Hence, matrix cracks are initiated already after a small number of cycles. 5. Conclusions A Ti-6242/SiC composite was characterized with respect to mechanical properties under creep, isothermal and thermomechanical fatigue conditions. The main results were: (1) SiC fiber reinforcement improves the creep resistance considerably at 550 ◦ C. Failure is dominated by fiber failure. (2) The isothermal fatigue limit of the composite at 550 ◦ C is considerably higher than at room temperature. Residual stresses due to the mismatch of the coefficient of thermal expansion of matrix and reinforcement may relax at high temperatures resulting in a cyclic tension–compression stress state in the matrix and, consequently, a suppression of matrix crack formation. (3) In contrast, applied stresses and internal stresses accumulate in the matrix during isothermal room temperature fatigue tests and OP-TMF tests.
543
(4) Under TMF conditions, life time of the composite is higher under OP- than IP-conditions if Nf < 1000. Life time is controlled by fiber strength under IP-conditions. Residual stresses in the matrix relax during the first cycles while fiber stresses increase. (5) Under OP-TMF conditions we find a change in behavior at about 1800 MPa. Above 1800 MPa, the curve is shifted to higher stresses parallel to IP-TMF. Below 1800 MPa, multiple cracks are found, which proceed to the first row of fibers. This results in oxidation of the SiC fibers, which causes a reduction in fiber strength and consequently in fiber failure. Acknowledgements Financial Support by the “Deutsche Forschungsgemeinschaft” (DFG Po 405/5-1 and 405/5-3) is gratefully acknowledged. The technical support of BAM Division V.1 (Microstructure of Engineering Materials) is gratefully appreciated. We also wish to thank Dr. Goebbels of BAM Division VIII.3 (Radiological Methods) for CT-measurements.
References [1] I.J. Polmear, Light Alloys—Metallurgy of the Light Metals, 3rd ed., Hodder Headline, PLC, London, 1996, p. 261. [2] C. Leyens, J. Hausmann, J. Kumpfert, in: M. Peters, C. Leyens (Eds.), Titan und Titanlegierungen, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim, 2002, pp. 321–350. [3] G.T. Ward, J. Herrmann, M. Hillberry, J. Compos. Tech. Res. 17 (1995) 205–211. [4] T.P. Gabb, J. Gayda, A. MacKay, J. Compos. Mater. 24 (1990) 667–686. [5] B.P. Sander, S. Mall, J. Compos. Tech. Res. 16 (1994) 304–313. [6] R. Talreja, Mater. Sci. Engg. A200 (1995) 21–28. [7] T.P. Gabb, J. Gayda, P.A. Bartolotta, M.G. Castelli, Int. J. Fatigue 15 (1993) 413–422. [8] M.G. Castelli, in: AGARD Conf. Proc. 569, 2.-4.10.1995, Banff, Canada, PP. 12–1. [9] M.G. Castelli, J. Gayda, in: Proc. Conf. on Reliability, stress analysis and failure prevention, 1993, Albuquerque, New Mexico, USA, 19.-22. Sept. 1993, DE-Vol. 55, ASME 1993, 213. [10] M.G. Castelli, M.G. Life Prediction Methodology for Titanium Matrix Composites, ASTM STP 1253, American Soc. for Testing and Materials, 1996, 412–431. [11] R. Neu, T. Nicholas, J. Compos. Tech. Res. 16 (1994) 214–224. [12] T. Nicholas, Mater. Sci. Eng. A200 (1995) 29–37. [13] L. Figel, B. G¨unther, Modeling the high temperature longitudinal fatigue behavior of metal matrix composites (SiC/Ti-6242): non-linear time dependent matrix behavior, Int. J. Fatigue, submitted for publication. [14] R.R. Boyer, Adv. Perform. Mater. 2 (1995) 349. [15] R. Leucht, H.J. Dudek, K. Weber, A. Werner, Z. Metallkd. 87 (1996) 424. [16] E. Olevsky, H.J. Dudek, W.A. Kaysser, Acta Mater. 44 (1996) 715. [17] P.W.M. Peters, J. Hemptenmacher, Composites, Part A 33 (2002) 1373–1379. [18] M.P. Thomas, G. Howard, R. Pather, M.R. Brown, P.H. Tranter, Int. J. Fatigue 23 (2001) 851–856. [19] P.W.M. Peters, J. Hempenmacher, K. Weber, H. Assler, J. Comp. Techn. Res. 24 (2002) 246–253. [20] J. Hemptenmacher, P.W.M. Peters, K. Weber, Adv. Eng. Mater. 6 (2004) 129–133. [21] R. Sinclair, M. Preuss, E. Maire, J.-Y. Buffiere, P. Brown, P.J. Withers, Acta Mater. 52 (2004) 1423–1438. [22] S.W. Schwenker, D. Eylon, Metall. Mater. Trans. A 27 (1996) 4193–4204.
544
D. Bettge et al. / Materials Science and Engineering A 452–453 (2007) 536–544
[23] P.W.M. Peters, H. Assler, J. Hemptenmacher, Z. Xia, Proc. ECCM-9, Brighton, June 4–7, 2000. [24] P.W.M. Peters, W.D. Zeng, J. Hemptenmacher, Z. Xia, Comp. Mater., Test. Design, ASTM STP 1436 (2003) 273–288. [25] Y. Tanaka, Z.-Y. Deng, Y.-F. Liu, C. Masuda, Acta Mater. 51 (2003) 6329–6340. [26] J.-L. Bobet, C. Masuda, Y. Kagawa, J. Mater. Sci. 32 (1997) 6357–6369.
[27] P.W.M. Peters, J. Hemptenmacher, B. G¨unther, D. Bettge, P. Portella, in: Proc. ECCM-11, Rhodos, May 31–June 3, 2004. [28] R. Neu, Fatigue Fract. Eng. M 16 (1993) 811–828. [29] C. Todd, DLR Internal Report IB 334-02/06, 2002. [30] O. Jin, W.S. Johnson, Composites A35 (2004) 113–120. [31] F. McBagonluri, E. Akpan, C. Mercer, W. Shen, W.O. Soboyejo, Mater. Sci. Engg. A405 (2005) 111–134.