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The fatigue and fracture behavior of a gamma-titanium aluminide intermetallic: Influence of ductile phase reinforcement
Engineering Fracture Mechanics Vol. 56, No. 3, pp. 379-395, 1997
Pergamon
PH: S0013-7944(96)00101-4
THE FATIGUE AND FRACTURE GAMMA-TITANIUM ALUMINIDE INFLUENCE OF DUCTILE PHASE
Copyright © 1996 Elsevier Science Ltd Printed in Great Britain. All fights reserved 0013-7944/97 $17.oo + o.oo
BEHAVIOR OF A INTERMETALLIC: REINFORCEMENT
W. O. SOBOYEJO and F. YE Department of Materials Science and Engineering, The Ohio State University, 2041, College Road, Columbus, Ohio 43210, U.S.A. T. S. SRIVATSAN Department of Mechanical Engineering, The University of Akron, Akron, Ohio 44325, U.S.A. Abstract--This paper presents the results of recent studies on the ambient temperature fracture toughness and cyclic crack growth characteristics of a ductile-phase toughened gamma-titanium aluminide intermetallic alloy reinforced with TiNb particles. Under monotonic loading, substantial toughening is achieved in the composite microstructure and is ascribed to bridging of the crack by uncracked TiNb particle reinforcements in the wake of the crack tip. Crack-particle interactions such as crack bridging, bifurcation, deflection and crack renucteation contribute to an enhancement in toughness. Under cyclic loading, ductile phase toughening was found to be less effective resulting in an inferior fatigue crack growth resistance of the composite when compared to the monolithic counterpart. The lower cyclic fatigue resistance is attributed to an increased susceptibility of the ductile phase to fatigue failure, coupled with the mutually interactive influences of an absence of bridging contributions to crack-tip shielding, higher crack-tip opening displacements and the intrinsic effects of inelastic strains that are developed in the ductile phase-reinforced composite matrix. Copyright © 1996 Elsevier Science Ltd
INTRODUCTION THE LIMITATIONSimposed by the melting temperature of the family of nickel-base, and to some extent the cobalt-base, superalloys have necessitated a growing need for more efficient, lightweight, high temperature materials to meet the demands of structural, engine and propulsion systems in the hypersonic transportation arena such as the supersonic national aerospace plane (NASP), jet engines and spacecraft requiring high performance in the temperature range 10881586 K. This led to a focus on the family of ordered intermetallic materials. As monolithic materials, they offered potentially attractive properties which include: good strength at elevated temperatures, high creep resistance, high stiffness, low density, and good resistance to oxidation and sulfidation compared to the conventional titanium alloys, and even the family of superalloys [1-8]. In fact, their density is about half the density of conventional nickel-base alloys. Their reduced density makes them an attractive candidate for future aero-engine applications. This engendered considerable scientific and technological interest in both monolithic intermetallics and composites with intermetallic matrices. Several of the ordered intermetallic compounds of TiAI, Ti3A1, FeA1, NiA1, Ni3AI and MoSi2 have emerged as attractive and viable high-temperature materials based on their melting point and the limited diffusion that is a consequence of their ordered structure. They have several advantages to offer over contending ceramics in that: (1) they are electrically conducting, and consequently can be machined into shapes by techniques such as electro-discharge machining; (2) are amenable to conventional joining techniques; (3) can be non-destructively evaluated by techniques used for metals. A corollary to the stronger bonding energy and strength of the family of monolithic intermetallics is that they suffer from several drawbacks. Firstly, their structural use is severely restricted by their inferior room-temperature tensile ductility, inadequate fracture toughness properties and inferior fracture resistance. The key reasons contributing to their inferior room temperature ductility, fracture toughness and damage tolerance capability are quite varied depending on the monolithic intermetallic, and include: (a) poor grain boundary cohesion, (b) an insufficient 379
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al.
number of slip systems to satisfy von-Mises criterion for deformation of polycrystalline materials, (c) limited cross-slip, and (d) impurity locking of dislocations[9-11]. Other factors which are of concern with these materials are a susceptibility to environmental embrittlement [7] and their inability to readily lend themselves to conventional thermomechanical processing (rolling, forging and extrusion) and machining[12, 13]. Attempts to improve the ambient temperature ductility, toughness and damage tolerance of the family of titanium aluminide intermetallic materials have centered around using alloy modification, innovative processing (wrought and secondary) and composite toughening (extrinsic) techniques[I-4, 14-16]. Alloy design through judicious microalloying with elements such as vanadium, niobium, molybdenum, manganese and chromium coupled with use of novel wrought and secondary processing techniques have resulted in engineering desirable microstructural features which offer attractive improvements in properties through an intrinsic change in the deformation mechanism [2-4, 17, 18]. Composite toughening is achieved through the incorporation of either a brittle or ductile second-phase in the form of particulates, whiskers and fibers [19-26]. Besides improved fatigue resistance [19, 20] and fracture toughness [19, 20], such reinforcements may also offer improvements in creep resistance, especially when they are brittle. An attractive and viable approach that has been recently used to achieve significant toughness improvements in intrinsically brittle intermetallic materials is the incorporation of small volume fractions of ductile reinforcements. Typically, these reinforcements are niobium, TiNb and Ti-6A1-4V particles [14, 16, 20, 27]. The ductile reinforcements undergo plastic deformation and stretch during matrix crack opening, thereby, producing closure tractions over the crack faces and thus partially shielding the crack-tip from the remotely applied stress. The enhancement in fracture toughness arises from the tractions produced by the unbroken ductile ligaments bridging the crack wake and the concomitant far-field stress required to overcome the tractions before the occurrence of unstable crack tip extension. By maximizing the extent of plastic stretch by the ductile reinforcement phase before crack tip extension, the toughness of the composite can be enhanced [20, 28]. Additional factors which contribute to crack-tip shielding are crack trapping, crack renucleation, crack deflection and decohesion along the matrix-reinforcement interfaces. The extent of toughening depends on the conjoint and mutually interactive influences of: (a) the length of the bridging zone in the wake of the crack-tip and (b) the deformation behavior of the ductile reinforcement phase. The strength of the interface (weak or strong) appears to exert an influence on the constraint, and hence, the flow behavior [29-31]. A constrained reinforcement, exhibiting high flow stress and low plastic stretch, has been suggested to have only a marginal influence on fracture toughness. Debonding at the matrix-reinforcement interfaces during crack opening would lower the constraints of the particulate reinforcement, facilitating the occurrence of extensive plastic stretch[29], while concurrently increasing the work of rupture[31]. The preponderance of research studies directed towards understanding the mechanical properties of ductile particle reinforced intrinsically brittle intermetallic and ceramic materials have focussed on modeling and analysis of strength and fracture toughness behavior[14, 31-37]. An adequate understanding of the contributions to toughening under monotonic loading conditions has been achieved. In these studies, the intrinsic effects of particle constitutive behavior on toughness were analyzed quantitatively based on assumptions of small scale bridging of an infinite crack located in an infinite specimen [38-40]. It was also observed that improved fracture toughness and steep fracture resistance curve (R-curve) can be obtained under conditions when the ductile reinforcement phase intersecting the crack path remains unbroken over significant dimensions behind the crack tip[19, 20]. However, there have been only a few experimental studies devoted exclusively to characterizing and understanding the response of these materials to Table 1. Chemical compositionof componentsof Ti-48AI/TiNb compositedeterminedby EDAX Composition (at. pet.) Ti A1 Nb Matrix Interface Reinforcement
48.5 48.1 57.5
51.5 51.5 15.3
-0.4 27.2
Gamma-titanium aluminide intermetallic
Fig. 1. Optical micrograph of the monolithic y-(Ti48AI) alloy.
Fig. 2. Optical micrograph of the TiAI/TiNb composite showing: (a) low magnification view of Ti48A1 + 20 vol.% TiNb structure, (b) high magnification view of the Ti-48A1 + 20 vol.% TiNb composite, (c) structure of the matrix and interface.
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Fig. 3. Scanning electron micrographs of the tensile fracture surface of the monolithic 7-Ti48AI alloy: (a) 25°C, (b) 700°C, (c) 982°C.
Gamma-titanium aluminide intermetallic
Fig. 4. Scanning electron micrographs of the tensile fracture surface of the Ti-48A1 + 20 vol.% TiNb composite showing: (a) 25°C, (b) 700°C, (c) 982°C.
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Fig. 5. Stable crack growth during R-curve testing showing: (a) crack initiation from notch, (b) crack extension through the matrix and TiNb strips, (c) crack tip blunting and bridging by the thicker (100/an) strips.
Gamma-titanium aluminide intermetallic
Fig. 7. Scanning electron micrographs of the fracture toughness specimens showing: (a) cleavage fracture in Ti-48A1, (b) Ti-48AI + 20 vol.% TiNb.
Fig. 8. Scanning electron micrographs of the fatigue crack propagation specimens showing overall fracture surface morphology: (a) Ti-48A1, (b) Ti~8AI + 20 vol.% TiNb.
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Fig. 12. Optical micrograph showing the interaction of the fatigue crack path with the microstructure (a) Ti-48 AI and (b) Ti-48A1 + 20 vol.% TiNb.
G a m m a - t i t a n i u m aluminide intermetallic
387
cyclic loading[17, 18, 26, 36]. Such information on cyclic response, at both ambient and elevated temperatures is essential and important because of the potential use of these materials in damage tolerant, temperature-critical and fatigue-sensitive applications. The broad objective of the present paper is to present the results of recent studies aimed at understanding the influence of ductile phase reinforcements on ambient temperature fracture toughness, cyclic crack growth and fracture behavior of a gamma-based titanium-aluminide alloy (Ti-48AI). The influence of ductile TiNb reinforcement phase on fatigue crack growth is compared with response of the composite microstructure under monotonic loading. The fracture characteristics of the composite material is discussed in the light of the competing influences of nature of loading, intrinsic microstructural effects and matrix deformation characteristics. MATERIAL Intermetallics rich in aluminum are attractive as candidate matrix materials because of their potentially low density, reasonably high melting point and good oxidation resistance [7, 8]. With this in mind the y-TiAl [Ti-48A1] alloy was chosen as the candidate intermetallic. The exact chemical composition (in weight percent) of the material is given in Table 1. The y-matrix material was procured from Nuclear Metals Inc. (Concord, MA) in the form o f - 3 5 mesh (average diameter 500 x 10-6 m) powder. The powder was produced by the plasma rotating electrode process. The monolithic Ti-48AI powder was canned in Ti-6A1-4V and extruded at a die temperature of 1343°C with a reduction ratio of 14:1. This resulted in the monolithic material having a duplex microstructure comprising ct2 + y lamellar packets and equiaxed y grains, as shown in Fig. 1. The Ti-48A1 + 20 vol.% TiNb intermetallic composite was produced by blending 35 mesh niobium and Ti-48AI powders. The mixture was processed under identical conditions that were used for the monolithic alloy. The resultant material had a duplex microstructure with non-uniform elongated strips of TiNb, as shown in Fig. 2. This microstructure has been shown in an earlier study[19] to consist predominantly of ~2. Energy dispersive X-ray analysis (EDAX) of the TiNb particles, interfacial regions and the Ti-48A1 matrix are summarized in Table 2. The actual composition of the interfacial region indicates the presence of both ~2 and v-phases at the interface. EXPERIMENTAL TECHNIQUES Uniaxial tensile tests were conducted on smooth cylindrical "button-headed" specimens with a gauge diameter of c a 3.18 mm. The tests were performed in laboratory air environment at temperatures between 25°C and 982°C. The elevated temperature tests were conducted after soaking the specimen at the temperature for one hr. The monolithic Ti-48AI and y-TiAl/TiNb composite specimens were loaded monotonically to failure at a strain rate of 5 x 104 s-~. The fatigue crack growth and fracture toughness resistance curve tests on the monolithic alloy and the composite counterpart were investigated using 25.4 mm long single edge notched (SEN) bend specimens having a square (6.35 m m x 6.35 mm) cross-section. The cyclic crack growth tests were conducted in laboratory air environment (temperature = 27°C; relative humidity of 55%), with specimens subjected to three-point constant amplitude loading after precracking under far-field cyclic compression [41,42]. A cyclic frequency of 10 Hz was employed. The growth of the crack was monitored using a high resolution (2.5 ~m) telescope connected to a video monitoring unit. An initial load-increasing scheme was used until crack growth was detected after about 106 cycles. The load range (AP = Pmax-Pmin) was then maintained Table 2. Comparison of tensile properties in as-extruded Ti-48AI and Ti-48A1 + 20 vol. pct, TiNb Temperature
Yield stress M P a (ksi)
Ultimate tensile (MPa (ksi)
Plastic elongation to failure(pct.)
(°c) 25 700 815 982
594 483 476 262
y
~ + TiNb
~
~ + TiNb
~
~ + TiNb
(86) (70) (69) (38)
--394 (57) 117 (17)
686 (100) 519 (84) 586 (85) 317 (46)
523 (76) 461 (67) 464(67) 132 (19)
1.7 2,8 6.0 26.0
--2.4 84.5
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w.o. SOBOYEJOet al.
constant until the end of the test. The crack propagation tests were stopped prior to specimen fracture. Metallographic examination of the interaction of the crack path with the microstructure was carried out, and the resultant micrographs used to determine crack deflection angles. Room temperature fracture toughness tests were conducted on SEN specimens that were precracked under far-field compression[41-43]. The specimens were loaded monotonically to failure at a ramp rate that corresponded to a stress intensity factor increase rate of 0.92 MPa(m)°5s -~. The conventional fracture toughness tests were conducted under monotonic loading to failure at a ramp rate of 0.92 MPa(m)°Ss-1. The specimens were pre-cracked under far-field compression [20] prior to monotonic loading under four-point (pure) bending. Resistance-curve (Rcurve) experiments were also performed to study the initiation and propagation of stable crack growth under monotonic loading. The R-curve experiments were performed under manual load control. They involved initial loading increments at stress intensity factors close to 6 MPa(m) °5. After unloading, both sides of the specimens were examined under an optical microscope to determine the precise conditions required for the onset of stable crack growth. The maximum stress intensity factors were increased by 1 MPa(m) °s during each loading increment and the tests were stopped only after the specimen had failed. Fracture surfaces of the deformed test specimens were examined in a scanning electron microscope (SEM) to: (a) determine the macroscopic fracture mode and (b) characterize the fine-scale topography and microscopic mechanisms governing fracture. The distinction between the macroscopic mode and microscopic fracture mechanism is based on the magnification level at which the observations are made. The macroscopic mode refers to the nature of failure, while the microscopic mechanism refers to the local failure process (cleavage, microvoid formation, growth, coalescence and nature of cracking).
RESULTS AND DISCUSSION Tensile behavior
Tensile properties of the as-extruded Ti-48AI and Ti-48AI + 20 Vol. % TiNb composite are summarized and compared in Table 2. Both the monolithic and composite material retain their relatively high strengths at temperatures up to 815°C. Increase in test temperature to 982°C was observed to cause a significant degradation in strength and a concomitant improvement in plastic elongation [Cp] to failure. The degradation in ultimate tensile strength is as high as 50% for the monolithic 7-TiA1 material and 80% for the Ti-48A1/TiNb composite. At the 982°C test temperature the plastic elongation to failure of the monolithic ~-TiAI alloy is 20% and 84.5% for the Ti-48AI/TiNb composite. The significant improvement in ductility with increase in test temperature is ascribed to the conjoint and mutually interactive influences of: (1) an activation of multiple slip systems[44,45] and (2) transitions in the micromechanisms of fracture that occur at temperatures around 650700°C[41,44,46] and 900-982°C[41]. The predicted ambient temperature strength of the Ti--48AI/TiNb composite based on the simple rule-of-mixtures (ROM) is 563 MPa (O'matrix = 594 MPa and trTiNb = 440 MPa). This value lies well within 10% of the experimentally determined value of 523 MPa, in spite of failure to account for: (i) the influence of reinforcement-matrix interfaces and (ii) the constraint effects associated with deformation of the TiNb reinforcements. The TiAI/TiNb composite exhibits no ductility below 815°C, while the monolithic ~-TiAI is ductile even at room temperature. The ductility of the composite improves significantly at 982°C where it exceeds that of the monolithic Ti-48AI. Examination of the tensile fracture surfaces provides useful information on the effect of microstructure on tensile ductility and fracture properties of the monolithic ~-TiAI alloy and the Ti--48AI/TiNb composite. At ambient temperature, fracture of the monolithic alloy occurred predominantly by cleavage [Fig. 3(a)]. However, at the elevated temperatures between 700°C and 815°C failure of the monolithic alloy was predominantly intergranular [Fig. 3(b)]. At the el-
Gamma-titanium aluminide intermetallic
389
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Crack Growth Aa (mm)
Fig. 6. R-curvebehaviorin the Ti48AI/TiNbcomposite. evated temperature of 982°C, failure of the monolithic alloy revealed features reminiscent of ductile failure, i.e. dimples and a transgranular fracture mode [Fig. 3(c)]. Monotonic fracture of the composite at ambient temperature (25°C) showed mixture of brittle failure through cleavage and intergranular cracking in the matrix, and ductile dimpled plus shear rupture mode in the TiNb reinforcements [Fig. 4(a)]. At temperatures between 700°C and 815°C tensile failure of the matrix was brittle intergranular, while the TiNb reinforcements failed by ductile transgranular fracture mode [Fig. 4(b)]. At the elevated test temperature of 982°C fracture was predominantly ductile and revealed transgranular failure in both the matrix and TiNb reinforcements [Fig. 4(c)].
Fracture toughness behavior Prior to discussing the data obtained from the fracture toughness test, it is important to present the visual evidence of stable crack growth that was obtained during R-curve testing. Stable crack growth was observed to initiate at a stress intensity factor of 15 MPa(m) °5 [Fig. 5(a)]. This is close to the matrix fracture toughness reported in earlier studies [20, 41]. Subsequent R-curve behavior was associated with stable crack growth through the gamma matrix and the thinner strips of TiNb [Fig. 5(b)]. From this observation it appears convincing that the thinner TiNb strips did not offer significant resistance to the growth of microscopic cracks. However, the thicker strips (of the order of 100/~m) of TiNb did offer some resistance to crack growth, with the cracks showing visible evidence of blunting caused by the strips [Fig. 5(c)]. The microscopic cracks deflected around the thicker (100/~m thick) strips and renucleated in the matrix on the opposite side. The resulting bridged cracks [Fig. 5(c)] extended easily through the thinner TiNb strips, while they were retarded and/or blunted by the thicker TiNb strips. Ductile phase toughening and resistance curve behavior (Fig. 6) in this TiNb-reinforced Ti-48AI intermetallic is attributed largely to the conjoint and interactive influences of crack bridging and crack-tip blunting by the thicker TiNb strips. The toughening arising from bridging of the crack by the thicker TiNb strips can be modeled using existing micromechanics models [40]. Using an analysis developed originally by Budiansky et al, [40], the toughness of the bridged composite is given by: Kc = KM + AKb = KM + 2(2/Tr)°'5otVf[cry/(X )0.5] dX,
(1)
where 0t is a constraint parameter between 1 and 6, Vf is the volume fraction of the secondphase (fibers) that participate in the bridging process (at 10%), KM is the matrix fracture toughness, X is the distance behind the crack tip and L is the bridge length, which is equal to the distance between the crack-tip and the last unfractured TiNb strip behind the crack tip. The
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SOBOYEJO
et al.
toughening arising solely due to crack bridging may be expressed as: ~.b = K c / K M = 1 + (2/KM)[(2/zr)°'5]otVf[cry/(X )°s] dX.
(2)
Using typical values of Vf = 0.1, KM = 15 MPa(m) °'5, ~ = 1.3 and L = 200 #m, the toughening ratio due to crack bridging is estimated from eq. (2) to be 1.16. The fracture toughness of the Ti--48AI/TiNb composite is thus estimated to be 17.4 MPa(m) °'5. This is less than the measured composite fracture toughness of 20 MPa(m) °5. The potential effects of crack-tip blunting by the ductile TiNb phase must therefore be considered in this analysis. This can be modeled using the following expression developed by Chan[47]: ~.bl =
K / K M = [1 +
Vf(E
-
l)]n-1/2n[1 -}-
Vf(l" -
I)n+I/2n[Ec/EM] n+l/2n
(3)
and
AKbl :
(~.bl - - 1 ) K M
,
(4)
where d M s = [,7;/% ]
(s)
r = [~rd/~],
(6)
and
where: (a) AKbl is the toughening increment due to crack-tip blunting, (b) K is the applied stress intensity factor, (c) try is the yield stress, (d) ef is the fracture strain, (e) n is the inverse of strain hardening exponent (N), (0 EM is the matrix modulus, and (g) the subscripts d and M denote the ductile and matrix phases, respectively. Using typical matrix and reinforcement values, E = 1, F = 4, E c / E M = 0.8, n = 20 and Vf = 0.1, the value of 2bl is estimated from eq. (3) to be 1.16. The toughening increment due to crack-tip blunting is therefore estimated to be 2.38 MPa(m) °5. The overall toughening can now be obtained by superposing the contributions to toughening from both crack-tip blunting and crack bridging. This predicts an overall composite fracture toughness of 19.8 MPa(m) °5, which is close to the measured fracture toughness of 20 MPa(m) °5. Toughening in this Ti-48A1/TiNb composite is ascribed to the conjoint and mutually competitive influences of crack bridging and crack-tip blunting. In summary, toughness of the ductile TiNb phase-reinforced composite is dictated by the mutually interactive influences of: (a) particle constitutive behavior, (b) fracture toughness of the matrix and (c) crack bridging by the TiNb strips. The composite would attain maximum toughness under the conjoint influence of crack bridging and crack-tip blunting before the occurrence of unstable crack extension. Thus, promoting or prolonging the fully-bridged configuration during stable crack-tip extension, by appropriate selection of reinforcements, can result in notable improvements in the fracture toughness of the composite. In fact, the important reinforcement parameters that directly contribute to improvements in fracture toughness of the composite are either a decrease in constraint and a lower flow stress, which would result in an increase in plastic rupture displacement of the particle, or a constrained particle exhibiting a higher maximum flow stress and a low plastic stretch. In fact, the occurrence of debonding or decohesion at the matrix-particle interfaces during crack opening facilitates a loss of through-thickness constraint, thereby, promoting extensive stretch while concurrently increasing the work of fracture. Also, the occurrence of such debonding at the interfaces minimizes the constraint (triaxiality) ahead of the crack-tip by relieving throughthickness stresses stemming from a transition from fully-constrained plane strain fracture process to an unconstrained "near-plane stress" ligament, which fractures at higher K c values. Fracture surface features of the fracture toughness specimens of the monolithic Ti--48AI alloy are shown in Fig. 7. Fracture in the matrix occurred primarily by cleavage across the equiaxed, ~ grains, with a "step-like" faceted fracture mode observed within the lamellar colonies [Fig. 7(a)]. Fracture of the Ti-48AI/TiNb composite occurred by: (a) a mixture of cleavage and intergranular fracture mode across the matrix and (b) a ductile dimpled fracture mode in the TiNb reinforcements [Fig. 7(b)] under monotonic loading conditions. This is in sharp con-
Gamma-titanium aluminide intermetallic
391
Stress Intmslty factor r a n ~ , A~ (]ksl On)u )
.......
I ,;
.......
I
. . . . . . .
~ 10"z
lO-I -O 'n-24AI-UNb 1093"C/@..~h/FAC A aI-6AI-4V mla mmmled [] aI-4VA r . ~ ,Ided St ~ R = 0.1
i
10-3
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A
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111-7 10' !
I I Itlil[
I
I
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I illll
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trast to the transgranular fracture mode observed in the TiNb reinforcements [Fig. 8(b)] and the predominantly cleavage fracture mode [Fig. 8(a)] observed in the Ti-48A1 matrix under cyclic loading conditions.
Fatigue crack growth The crack growth behavior in the monolithic Ti-48AI alloy under cyclic loading conditions is shown in Fig. 9 in terms of the growth rate per cycle (da/dN) as a function of the nominal stress intensity range AK (= Kmax - Kmin). Results of mill-annealed (MA) Ti-6A1-4V [48] and ~2 + fl processed Ti-24Al-11Nb[43] are also shown in the same figure for purposes of comparison. Under cyclic loads, the cracks tend to propagate subcritically at stress intensities well below the crack initiation toughness [15 MPa(m) °5] needed to initiate cracking under monotonic loads. Therefore, subcritical crack propagation under cyclic loads is an important mode contributing to degradation of the monolithic Ti-48AI alloy. The monolithic Ti48AI alloy is observed to have lower near-threshold fatigue crack growth rates than the mill-annealed Ti-6A1-4V and at2 + fl forged Ti-24Al-llNb. However, in the Paris regime the growth rates in the millannealed Ti-6A1-4V and Ti-24Al-llNb are similar to those in the monolithic Ti-48Ai. The fatigue crack growth resistance of the monolithic Ti-48A1 alloy was generally superior to that of the Ti-24AI-11Nb, although the toughness of the gamma-based alloy [15 MPa(m) °5] [41] is inferior to the toughness of the ~2 + fl alloy [23.4 MPa(m) °'5][43]. The fatigue crack growth rate data obtained for the as-extruded Ti-48AI/TiNb composite are compared with data obtained previously for fl-TiNb [19, 48, 49] and extruded Ti-48AI alloy in Fig. 10. The crack growth rates for the composite lie between those obtained for the Ti-48 AI matrix and the fl-TiNb reinforcements. Similar results have been reported in earlier studies on HIPped TiAI + 20 vol.% TiNb (Fig. 11)[19, 48, 49]. It is evident that crack propagation rates in the TiA1/TiNb composite are higher than those in the matrix alloy, especially at nearthreshold levels, in spite of the improved fracture toughness of the composite under monotonic loading. In fact, the ambient temperature fatigue threshold of the TiAI/TiNb composite is 20%
392
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I
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|
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6
,
i
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'lrUU+ 20'~ ( ~ ~m) 1 ~ C-R {EdllP)Ori,.,,e,,a,~: R = 0.I
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o 10-w
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Smm ~ z j t y
,
I S
,
,
,
,
10
~Etor rm~e, aK ( M ~ (m)")
Fig. 11. Comparison of the room temperature fatigue crack growth rates in HIEPped gamma alloys and fl-TiNb[17].
G a m m a - t i t a n i u m aluminide intermetallic
393
lower than in the unreinforced matrix. However, the crack growth rates in both the TiAI/TiNb composite and ~-TiAI are sensitive to the applied stress intensity range. The region of power law dependence is relatively small for both the TiA1/TiNb composite and the monolithic-~-TiAl alloy. With increasing stress intensity factor ranges (AK), a rapid increase in crack growth rate (da/dN) is observed for small changes in AK. The degradation in ambient temperature crack growth resistance is ascribed to the absence of crack bridging and interfacial debonding or crack deflection under cyclic loading conditions (Fig. 12). Since the degree of crack bridging is small, the near tip and applied stress intensities are expected to be equal, except for contributions to shielding from other competing mechanisms such as surface roughness-induced crack closure [50]. Additionally, the degradation in fatigue crack growth resistance of the TiAI/TiNb composite is most likely caused by the differences in microscopic fracture mechanisms under monotonic and cyclic loading (Figs 7 and 8) and the inelastic strains that are developed during cyclic deformation. The portioning and progressive accumulation of such strains results in a rapid degradation of the fatigue resistance of the TiA1/TiNb composite, and the resultant crack growth rates are between those of the matrix and the reinforcements. An alternative factor contributing to the degradation in cyclic crack growth resistance is the influence of ductile phase reinforcements on crack-tip opening displacement (CTOD), which is assumed to be directly proportional to crack-tip extension under cyclic loading. The plane strain crack opening displacement (COD) under cyclic loading is: At$ = ~max - - t~min =
(AK)2/2Etrys,
(7)
where 6 is the crack-tip opening displacement, and the subscripts "max" and "min" correspond to the maximum and minimum stress intensity factors. Assuming the extent of irreversibility to be the same for the monolithic Ti-48AI alloy and TiA1/TiNb composite, the increase in cyclic crack growth rates of the ductile phase-reinforced composite is ascribed to an absence of cracktip shielding from bridging. In fact, it is the fatigue-induced failure of the reinforcing TiNb particles that occurs before the advancing crack can establish a bridging zone. This is aided by the low stress intensities associated with cyclic loading and the concomitant small crack-tip opening displacements hinder the formation of a bridging zone. Fatigue crack propagation surfaces revealed transgranular fatigue crack growth in both the Ti-48AI monolithic alloy and the TiA1/TiNb composite. There was no evidence of interfacial decohesion and plastic stretching of the TiNb strips under cyclic loading [Fig. 12(b)] nor was there visible evidence of cracking of the TiNb particles ahead of the propagating crack tip [Fig. 12(a)]. Such an observation is consistent with similar observations on the fatigue and fracture behavior of ductile/brittle material systems, i.e. metal-ceramic interfaces. In these systems fracture predominates in the brittle ceramic phase under monotonic loading due to its much lower fracture toughness, whereas under cyclic loading the crack path is focussed in the metal [51]. The attraction and retention of the crack in the metal phase occurs because the metal matrix presents a region where strain can be either accumulated or localized thereby influencing progressive crack advance. This confirms the enhanced susceptibility of the metal matrix to failure by fatigue. The monolithic and cyclic fatigue tests demonstrate that the microstructure of the TiA1/ TiNb composite engineered for superior fracture toughness under monotonic loading fails to result in improved cyclic crack growth resistance. In spite of strong reinforcement (TiNb) matrix interfaces and a low strain hardening of the metallic phase, the occurrence of crack bridging and crack-tip blunting by the TiNb reinforcements significantly enhances monotonic fracture toughness. However, failure of the strong metal-reinforcement interfaces to provide impedance to fatigue crack growth under cyclic loading limits contributions from crack-tip shielding due to bridging. CONCLUSIONS The following are the observations of a study on the influence of ductile phase reinforcement on the fatigue and fracture behavior of a gamma-TiA1 intermetallic. (1) Increase in test temperature caused a loss of strength for both the monolithic alloy and EFM 56/3---D
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the composite counterpart. The extent of degradation was greater in the ductile phase reinforced composite. The loss of strength resulted in appreciable improvements in ductility. (2) Under monotonic loading significant improvements in room temperature fracture toughness of the brittle gamma-based Ti-48A1 intermetallic is achieved through the introduction of ductile TiNb particles. Toughening and resistance curve behavior of the TiAI/ TiNb intermetallic composite arises from the conjoint and mutually interactive influences of crack bridging and blunting by the thick (> 100/Am) TiNb strips. The thinner TiNb strips are fractured during stable crack growth under monotonic loading, and consequently they do not participate in either the bridging or the blunting processes. (3) Ductile phase toughening was found to be less effective under cyclic loading. This is attributed to mutually interactive influences of an absence of crack bridging, the effects of the ductile TiNb reinforcement on crack-tip opening displacement under cyclic loading, and the accumulation and portioning of the inelastic strains during cyclic deformation. The accumulation of inelastic strains during cyclic loading promotes subcritical fatigue damage in the composite microstructure.
Acknowledgements--WOS and FY would like to thank The National Science Foundation (Grant No. DMR 9458018) for financial support. Appreciation is also extended to Dr Bruce McDonald at The National Science Foundation for his encouragement and support of this research. One of the authors (TSS) gratefully acknowledges the State of Ohio: Board of Regents and The University of Akron for supporting this research study, through a faculty research grant.
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Mataga, P. A., Acta Metallurgica, 1989, 37, 3349. Ravichandran, K. S., Acta Metallurgica Materialia, 1992, 40(5), 1009. Przystupa, M. A. and Courtney, T. H., Metallurgical Transactions, 1982, 13A, 873. Cao, H. C., Dalgeleish, B, J., Deve, H. E., Elliott, C., Evans, A. G., Mehrabian, R. and Odette, G. R., Aeta Metallurgica, 1989, 37, 2969. Flinn, B. D., Ruhle, M. and Evans, A. G., Acta Metallurgica, 1989, 37, 3001. Nekkanti, R. M. and Dimiduk, D. M., Proceedings of Materials Research Society Symposium, Vol, 194, 1990, p. 95. Krstic, V. D., Nicholson, P. S. and Hoagland, R. G., Journal of American Ceramic Society, 1981, 64, 499. Ashby, M. F., Blunt, F. J. and Bannister, M., Aeta Metallurgica, 1989, 37, 1847. Venkatasubramanian, T. V. and Baker, T. J., Proceedings of Sixth International Conference on Fracture, ICF 6, Vol. 4, ed. S. R. Valluri. New Delhi, 1984, p. 2451. Erdogan, F. and Joseph, P. F., Journal of American Ceramic Society, 1989, 72, 262. Budiansky, B., Amazigo, J. C. and Evans, A. G., Journal of Mechanics and Physics of Solids, 1988, 36, 167. Soboyejo, W. O., Schwartz, D. S. and Sastry, S. M. L., Metallurgical Transactions, 1992, 23A, 2039. Soboyejo, W. O., Fatigue crack growth in gamma-based titanium aluminides. Fatigue 1993, ed. Barlan and Dickson. Montreal, Canada, EMAS, Warley, U.K., 1993, p. 1031. Soboyejo, W. O., Proceedinigs of M.E. Fine Symposium, ed. P. K. Liaw, J. R. Weertman, H. L. Marcus and J. S. Santner. TMS: The Minerals, Metals and Materials Society, Warrendale, PA, 1991, p. 399. Sastry, S. M. L. and Lipsitt, H. A., Metallurgical Transactions, 1977, 8A, 299. Huang, S. C. and Hall, E. L., Metallurgical Transactions, 1991, 22A, 427-439. Ritchie, R. O., International Metals Reviews, 1979, 20, 205. Chan, K. S., Metallurgical Transactions, 1992, 23A, 1737-1750. Venkateswara Rao, K. T., Odette, G. R. and Ritchie, R. O., Fatigue of Advanced Materials, ed. R. O. Ritchie, R. H. Dauskardt and B. N. Cox. Materials and Component Engineering Publications, Birmingham, U.K., 1991, p. 429. Brooks, C. L., McDonnell Douglas Corporation, St. Louis, MO, unpublished research, 1991. Suresh, S. and Ritchie, R. O., Metallurgical Transactions, 1982, 13A, 1627. Cannon, R. M., Dalgeish, B. J., Dauskardt, R. H., Oh, T. S. and Ritchie, R. O., Acta Metallurgica Materialia, 1991, 39, 2145.
(Received 10 May 1995)
Pergamon
PH: S0013-7944(96)00101-4
THE FATIGUE AND FRACTURE GAMMA-TITANIUM ALUMINIDE INFLUENCE OF DUCTILE PHASE
Copyright © 1996 Elsevier Science Ltd Printed in Great Britain. All fights reserved 0013-7944/97 $17.oo + o.oo
BEHAVIOR OF A INTERMETALLIC: REINFORCEMENT
W. O. SOBOYEJO and F. YE Department of Materials Science and Engineering, The Ohio State University, 2041, College Road, Columbus, Ohio 43210, U.S.A. T. S. SRIVATSAN Department of Mechanical Engineering, The University of Akron, Akron, Ohio 44325, U.S.A. Abstract--This paper presents the results of recent studies on the ambient temperature fracture toughness and cyclic crack growth characteristics of a ductile-phase toughened gamma-titanium aluminide intermetallic alloy reinforced with TiNb particles. Under monotonic loading, substantial toughening is achieved in the composite microstructure and is ascribed to bridging of the crack by uncracked TiNb particle reinforcements in the wake of the crack tip. Crack-particle interactions such as crack bridging, bifurcation, deflection and crack renucteation contribute to an enhancement in toughness. Under cyclic loading, ductile phase toughening was found to be less effective resulting in an inferior fatigue crack growth resistance of the composite when compared to the monolithic counterpart. The lower cyclic fatigue resistance is attributed to an increased susceptibility of the ductile phase to fatigue failure, coupled with the mutually interactive influences of an absence of bridging contributions to crack-tip shielding, higher crack-tip opening displacements and the intrinsic effects of inelastic strains that are developed in the ductile phase-reinforced composite matrix. Copyright © 1996 Elsevier Science Ltd
INTRODUCTION THE LIMITATIONSimposed by the melting temperature of the family of nickel-base, and to some extent the cobalt-base, superalloys have necessitated a growing need for more efficient, lightweight, high temperature materials to meet the demands of structural, engine and propulsion systems in the hypersonic transportation arena such as the supersonic national aerospace plane (NASP), jet engines and spacecraft requiring high performance in the temperature range 10881586 K. This led to a focus on the family of ordered intermetallic materials. As monolithic materials, they offered potentially attractive properties which include: good strength at elevated temperatures, high creep resistance, high stiffness, low density, and good resistance to oxidation and sulfidation compared to the conventional titanium alloys, and even the family of superalloys [1-8]. In fact, their density is about half the density of conventional nickel-base alloys. Their reduced density makes them an attractive candidate for future aero-engine applications. This engendered considerable scientific and technological interest in both monolithic intermetallics and composites with intermetallic matrices. Several of the ordered intermetallic compounds of TiAI, Ti3A1, FeA1, NiA1, Ni3AI and MoSi2 have emerged as attractive and viable high-temperature materials based on their melting point and the limited diffusion that is a consequence of their ordered structure. They have several advantages to offer over contending ceramics in that: (1) they are electrically conducting, and consequently can be machined into shapes by techniques such as electro-discharge machining; (2) are amenable to conventional joining techniques; (3) can be non-destructively evaluated by techniques used for metals. A corollary to the stronger bonding energy and strength of the family of monolithic intermetallics is that they suffer from several drawbacks. Firstly, their structural use is severely restricted by their inferior room-temperature tensile ductility, inadequate fracture toughness properties and inferior fracture resistance. The key reasons contributing to their inferior room temperature ductility, fracture toughness and damage tolerance capability are quite varied depending on the monolithic intermetallic, and include: (a) poor grain boundary cohesion, (b) an insufficient 379
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number of slip systems to satisfy von-Mises criterion for deformation of polycrystalline materials, (c) limited cross-slip, and (d) impurity locking of dislocations[9-11]. Other factors which are of concern with these materials are a susceptibility to environmental embrittlement [7] and their inability to readily lend themselves to conventional thermomechanical processing (rolling, forging and extrusion) and machining[12, 13]. Attempts to improve the ambient temperature ductility, toughness and damage tolerance of the family of titanium aluminide intermetallic materials have centered around using alloy modification, innovative processing (wrought and secondary) and composite toughening (extrinsic) techniques[I-4, 14-16]. Alloy design through judicious microalloying with elements such as vanadium, niobium, molybdenum, manganese and chromium coupled with use of novel wrought and secondary processing techniques have resulted in engineering desirable microstructural features which offer attractive improvements in properties through an intrinsic change in the deformation mechanism [2-4, 17, 18]. Composite toughening is achieved through the incorporation of either a brittle or ductile second-phase in the form of particulates, whiskers and fibers [19-26]. Besides improved fatigue resistance [19, 20] and fracture toughness [19, 20], such reinforcements may also offer improvements in creep resistance, especially when they are brittle. An attractive and viable approach that has been recently used to achieve significant toughness improvements in intrinsically brittle intermetallic materials is the incorporation of small volume fractions of ductile reinforcements. Typically, these reinforcements are niobium, TiNb and Ti-6A1-4V particles [14, 16, 20, 27]. The ductile reinforcements undergo plastic deformation and stretch during matrix crack opening, thereby, producing closure tractions over the crack faces and thus partially shielding the crack-tip from the remotely applied stress. The enhancement in fracture toughness arises from the tractions produced by the unbroken ductile ligaments bridging the crack wake and the concomitant far-field stress required to overcome the tractions before the occurrence of unstable crack tip extension. By maximizing the extent of plastic stretch by the ductile reinforcement phase before crack tip extension, the toughness of the composite can be enhanced [20, 28]. Additional factors which contribute to crack-tip shielding are crack trapping, crack renucleation, crack deflection and decohesion along the matrix-reinforcement interfaces. The extent of toughening depends on the conjoint and mutually interactive influences of: (a) the length of the bridging zone in the wake of the crack-tip and (b) the deformation behavior of the ductile reinforcement phase. The strength of the interface (weak or strong) appears to exert an influence on the constraint, and hence, the flow behavior [29-31]. A constrained reinforcement, exhibiting high flow stress and low plastic stretch, has been suggested to have only a marginal influence on fracture toughness. Debonding at the matrix-reinforcement interfaces during crack opening would lower the constraints of the particulate reinforcement, facilitating the occurrence of extensive plastic stretch[29], while concurrently increasing the work of rupture[31]. The preponderance of research studies directed towards understanding the mechanical properties of ductile particle reinforced intrinsically brittle intermetallic and ceramic materials have focussed on modeling and analysis of strength and fracture toughness behavior[14, 31-37]. An adequate understanding of the contributions to toughening under monotonic loading conditions has been achieved. In these studies, the intrinsic effects of particle constitutive behavior on toughness were analyzed quantitatively based on assumptions of small scale bridging of an infinite crack located in an infinite specimen [38-40]. It was also observed that improved fracture toughness and steep fracture resistance curve (R-curve) can be obtained under conditions when the ductile reinforcement phase intersecting the crack path remains unbroken over significant dimensions behind the crack tip[19, 20]. However, there have been only a few experimental studies devoted exclusively to characterizing and understanding the response of these materials to Table 1. Chemical compositionof componentsof Ti-48AI/TiNb compositedeterminedby EDAX Composition (at. pet.) Ti A1 Nb Matrix Interface Reinforcement
48.5 48.1 57.5
51.5 51.5 15.3
-0.4 27.2
Gamma-titanium aluminide intermetallic
Fig. 1. Optical micrograph of the monolithic y-(Ti48AI) alloy.
Fig. 2. Optical micrograph of the TiAI/TiNb composite showing: (a) low magnification view of Ti48A1 + 20 vol.% TiNb structure, (b) high magnification view of the Ti-48A1 + 20 vol.% TiNb composite, (c) structure of the matrix and interface.
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Fig. 3. Scanning electron micrographs of the tensile fracture surface of the monolithic 7-Ti48AI alloy: (a) 25°C, (b) 700°C, (c) 982°C.
Gamma-titanium aluminide intermetallic
Fig. 4. Scanning electron micrographs of the tensile fracture surface of the Ti-48A1 + 20 vol.% TiNb composite showing: (a) 25°C, (b) 700°C, (c) 982°C.
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Fig. 5. Stable crack growth during R-curve testing showing: (a) crack initiation from notch, (b) crack extension through the matrix and TiNb strips, (c) crack tip blunting and bridging by the thicker (100/an) strips.
Gamma-titanium aluminide intermetallic
Fig. 7. Scanning electron micrographs of the fracture toughness specimens showing: (a) cleavage fracture in Ti-48A1, (b) Ti-48AI + 20 vol.% TiNb.
Fig. 8. Scanning electron micrographs of the fatigue crack propagation specimens showing overall fracture surface morphology: (a) Ti-48A1, (b) Ti~8AI + 20 vol.% TiNb.
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Fig. 12. Optical micrograph showing the interaction of the fatigue crack path with the microstructure (a) Ti-48 AI and (b) Ti-48A1 + 20 vol.% TiNb.
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cyclic loading[17, 18, 26, 36]. Such information on cyclic response, at both ambient and elevated temperatures is essential and important because of the potential use of these materials in damage tolerant, temperature-critical and fatigue-sensitive applications. The broad objective of the present paper is to present the results of recent studies aimed at understanding the influence of ductile phase reinforcements on ambient temperature fracture toughness, cyclic crack growth and fracture behavior of a gamma-based titanium-aluminide alloy (Ti-48AI). The influence of ductile TiNb reinforcement phase on fatigue crack growth is compared with response of the composite microstructure under monotonic loading. The fracture characteristics of the composite material is discussed in the light of the competing influences of nature of loading, intrinsic microstructural effects and matrix deformation characteristics. MATERIAL Intermetallics rich in aluminum are attractive as candidate matrix materials because of their potentially low density, reasonably high melting point and good oxidation resistance [7, 8]. With this in mind the y-TiAl [Ti-48A1] alloy was chosen as the candidate intermetallic. The exact chemical composition (in weight percent) of the material is given in Table 1. The y-matrix material was procured from Nuclear Metals Inc. (Concord, MA) in the form o f - 3 5 mesh (average diameter 500 x 10-6 m) powder. The powder was produced by the plasma rotating electrode process. The monolithic Ti-48AI powder was canned in Ti-6A1-4V and extruded at a die temperature of 1343°C with a reduction ratio of 14:1. This resulted in the monolithic material having a duplex microstructure comprising ct2 + y lamellar packets and equiaxed y grains, as shown in Fig. 1. The Ti-48A1 + 20 vol.% TiNb intermetallic composite was produced by blending 35 mesh niobium and Ti-48AI powders. The mixture was processed under identical conditions that were used for the monolithic alloy. The resultant material had a duplex microstructure with non-uniform elongated strips of TiNb, as shown in Fig. 2. This microstructure has been shown in an earlier study[19] to consist predominantly of ~2. Energy dispersive X-ray analysis (EDAX) of the TiNb particles, interfacial regions and the Ti-48A1 matrix are summarized in Table 2. The actual composition of the interfacial region indicates the presence of both ~2 and v-phases at the interface. EXPERIMENTAL TECHNIQUES Uniaxial tensile tests were conducted on smooth cylindrical "button-headed" specimens with a gauge diameter of c a 3.18 mm. The tests were performed in laboratory air environment at temperatures between 25°C and 982°C. The elevated temperature tests were conducted after soaking the specimen at the temperature for one hr. The monolithic Ti-48AI and y-TiAl/TiNb composite specimens were loaded monotonically to failure at a strain rate of 5 x 104 s-~. The fatigue crack growth and fracture toughness resistance curve tests on the monolithic alloy and the composite counterpart were investigated using 25.4 mm long single edge notched (SEN) bend specimens having a square (6.35 m m x 6.35 mm) cross-section. The cyclic crack growth tests were conducted in laboratory air environment (temperature = 27°C; relative humidity of 55%), with specimens subjected to three-point constant amplitude loading after precracking under far-field cyclic compression [41,42]. A cyclic frequency of 10 Hz was employed. The growth of the crack was monitored using a high resolution (2.5 ~m) telescope connected to a video monitoring unit. An initial load-increasing scheme was used until crack growth was detected after about 106 cycles. The load range (AP = Pmax-Pmin) was then maintained Table 2. Comparison of tensile properties in as-extruded Ti-48AI and Ti-48A1 + 20 vol. pct, TiNb Temperature
Yield stress M P a (ksi)
Ultimate tensile (MPa (ksi)
Plastic elongation to failure(pct.)
(°c) 25 700 815 982
594 483 476 262
y
~ + TiNb
~
~ + TiNb
~
~ + TiNb
(86) (70) (69) (38)
--394 (57) 117 (17)
686 (100) 519 (84) 586 (85) 317 (46)
523 (76) 461 (67) 464(67) 132 (19)
1.7 2,8 6.0 26.0
--2.4 84.5
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constant until the end of the test. The crack propagation tests were stopped prior to specimen fracture. Metallographic examination of the interaction of the crack path with the microstructure was carried out, and the resultant micrographs used to determine crack deflection angles. Room temperature fracture toughness tests were conducted on SEN specimens that were precracked under far-field compression[41-43]. The specimens were loaded monotonically to failure at a ramp rate that corresponded to a stress intensity factor increase rate of 0.92 MPa(m)°5s -~. The conventional fracture toughness tests were conducted under monotonic loading to failure at a ramp rate of 0.92 MPa(m)°Ss-1. The specimens were pre-cracked under far-field compression [20] prior to monotonic loading under four-point (pure) bending. Resistance-curve (Rcurve) experiments were also performed to study the initiation and propagation of stable crack growth under monotonic loading. The R-curve experiments were performed under manual load control. They involved initial loading increments at stress intensity factors close to 6 MPa(m) °5. After unloading, both sides of the specimens were examined under an optical microscope to determine the precise conditions required for the onset of stable crack growth. The maximum stress intensity factors were increased by 1 MPa(m) °s during each loading increment and the tests were stopped only after the specimen had failed. Fracture surfaces of the deformed test specimens were examined in a scanning electron microscope (SEM) to: (a) determine the macroscopic fracture mode and (b) characterize the fine-scale topography and microscopic mechanisms governing fracture. The distinction between the macroscopic mode and microscopic fracture mechanism is based on the magnification level at which the observations are made. The macroscopic mode refers to the nature of failure, while the microscopic mechanism refers to the local failure process (cleavage, microvoid formation, growth, coalescence and nature of cracking).
RESULTS AND DISCUSSION Tensile behavior
Tensile properties of the as-extruded Ti-48AI and Ti-48AI + 20 Vol. % TiNb composite are summarized and compared in Table 2. Both the monolithic and composite material retain their relatively high strengths at temperatures up to 815°C. Increase in test temperature to 982°C was observed to cause a significant degradation in strength and a concomitant improvement in plastic elongation [Cp] to failure. The degradation in ultimate tensile strength is as high as 50% for the monolithic 7-TiA1 material and 80% for the Ti-48A1/TiNb composite. At the 982°C test temperature the plastic elongation to failure of the monolithic ~-TiAI alloy is 20% and 84.5% for the Ti-48AI/TiNb composite. The significant improvement in ductility with increase in test temperature is ascribed to the conjoint and mutually interactive influences of: (1) an activation of multiple slip systems[44,45] and (2) transitions in the micromechanisms of fracture that occur at temperatures around 650700°C[41,44,46] and 900-982°C[41]. The predicted ambient temperature strength of the Ti--48AI/TiNb composite based on the simple rule-of-mixtures (ROM) is 563 MPa (O'matrix = 594 MPa and trTiNb = 440 MPa). This value lies well within 10% of the experimentally determined value of 523 MPa, in spite of failure to account for: (i) the influence of reinforcement-matrix interfaces and (ii) the constraint effects associated with deformation of the TiNb reinforcements. The TiAI/TiNb composite exhibits no ductility below 815°C, while the monolithic ~-TiAI is ductile even at room temperature. The ductility of the composite improves significantly at 982°C where it exceeds that of the monolithic Ti-48AI. Examination of the tensile fracture surfaces provides useful information on the effect of microstructure on tensile ductility and fracture properties of the monolithic ~-TiAI alloy and the Ti--48AI/TiNb composite. At ambient temperature, fracture of the monolithic alloy occurred predominantly by cleavage [Fig. 3(a)]. However, at the elevated temperatures between 700°C and 815°C failure of the monolithic alloy was predominantly intergranular [Fig. 3(b)]. At the el-
Gamma-titanium aluminide intermetallic
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24
n ml
22
20
II
S
m m m
w e~
18 m
u mm 16
m 14
m
,
0.0
I
0.1
i
I
0.2
i
I
0.3
i
I
,
0.4
I
0.5
,
0.6
Crack Growth Aa (mm)
Fig. 6. R-curvebehaviorin the Ti48AI/TiNbcomposite. evated temperature of 982°C, failure of the monolithic alloy revealed features reminiscent of ductile failure, i.e. dimples and a transgranular fracture mode [Fig. 3(c)]. Monotonic fracture of the composite at ambient temperature (25°C) showed mixture of brittle failure through cleavage and intergranular cracking in the matrix, and ductile dimpled plus shear rupture mode in the TiNb reinforcements [Fig. 4(a)]. At temperatures between 700°C and 815°C tensile failure of the matrix was brittle intergranular, while the TiNb reinforcements failed by ductile transgranular fracture mode [Fig. 4(b)]. At the elevated test temperature of 982°C fracture was predominantly ductile and revealed transgranular failure in both the matrix and TiNb reinforcements [Fig. 4(c)].
Fracture toughness behavior Prior to discussing the data obtained from the fracture toughness test, it is important to present the visual evidence of stable crack growth that was obtained during R-curve testing. Stable crack growth was observed to initiate at a stress intensity factor of 15 MPa(m) °5 [Fig. 5(a)]. This is close to the matrix fracture toughness reported in earlier studies [20, 41]. Subsequent R-curve behavior was associated with stable crack growth through the gamma matrix and the thinner strips of TiNb [Fig. 5(b)]. From this observation it appears convincing that the thinner TiNb strips did not offer significant resistance to the growth of microscopic cracks. However, the thicker strips (of the order of 100/~m) of TiNb did offer some resistance to crack growth, with the cracks showing visible evidence of blunting caused by the strips [Fig. 5(c)]. The microscopic cracks deflected around the thicker (100/~m thick) strips and renucleated in the matrix on the opposite side. The resulting bridged cracks [Fig. 5(c)] extended easily through the thinner TiNb strips, while they were retarded and/or blunted by the thicker TiNb strips. Ductile phase toughening and resistance curve behavior (Fig. 6) in this TiNb-reinforced Ti-48AI intermetallic is attributed largely to the conjoint and interactive influences of crack bridging and crack-tip blunting by the thicker TiNb strips. The toughening arising from bridging of the crack by the thicker TiNb strips can be modeled using existing micromechanics models [40]. Using an analysis developed originally by Budiansky et al, [40], the toughness of the bridged composite is given by: Kc = KM + AKb = KM + 2(2/Tr)°'5otVf[cry/(X )0.5] dX,
(1)
where 0t is a constraint parameter between 1 and 6, Vf is the volume fraction of the secondphase (fibers) that participate in the bridging process (at 10%), KM is the matrix fracture toughness, X is the distance behind the crack tip and L is the bridge length, which is equal to the distance between the crack-tip and the last unfractured TiNb strip behind the crack tip. The
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SOBOYEJO
et al.
toughening arising solely due to crack bridging may be expressed as: ~.b = K c / K M = 1 + (2/KM)[(2/zr)°'5]otVf[cry/(X )°s] dX.
(2)
Using typical values of Vf = 0.1, KM = 15 MPa(m) °'5, ~ = 1.3 and L = 200 #m, the toughening ratio due to crack bridging is estimated from eq. (2) to be 1.16. The fracture toughness of the Ti--48AI/TiNb composite is thus estimated to be 17.4 MPa(m) °'5. This is less than the measured composite fracture toughness of 20 MPa(m) °5. The potential effects of crack-tip blunting by the ductile TiNb phase must therefore be considered in this analysis. This can be modeled using the following expression developed by Chan[47]: ~.bl =
K / K M = [1 +
Vf(E
-
l)]n-1/2n[1 -}-
Vf(l" -
I)n+I/2n[Ec/EM] n+l/2n
(3)
and
AKbl :
(~.bl - - 1 ) K M
,
(4)
where d M s = [,7;/% ]
(s)
r = [~rd/~],
(6)
and
where: (a) AKbl is the toughening increment due to crack-tip blunting, (b) K is the applied stress intensity factor, (c) try is the yield stress, (d) ef is the fracture strain, (e) n is the inverse of strain hardening exponent (N), (0 EM is the matrix modulus, and (g) the subscripts d and M denote the ductile and matrix phases, respectively. Using typical matrix and reinforcement values, E = 1, F = 4, E c / E M = 0.8, n = 20 and Vf = 0.1, the value of 2bl is estimated from eq. (3) to be 1.16. The toughening increment due to crack-tip blunting is therefore estimated to be 2.38 MPa(m) °5. The overall toughening can now be obtained by superposing the contributions to toughening from both crack-tip blunting and crack bridging. This predicts an overall composite fracture toughness of 19.8 MPa(m) °5, which is close to the measured fracture toughness of 20 MPa(m) °5. Toughening in this Ti-48A1/TiNb composite is ascribed to the conjoint and mutually competitive influences of crack bridging and crack-tip blunting. In summary, toughness of the ductile TiNb phase-reinforced composite is dictated by the mutually interactive influences of: (a) particle constitutive behavior, (b) fracture toughness of the matrix and (c) crack bridging by the TiNb strips. The composite would attain maximum toughness under the conjoint influence of crack bridging and crack-tip blunting before the occurrence of unstable crack extension. Thus, promoting or prolonging the fully-bridged configuration during stable crack-tip extension, by appropriate selection of reinforcements, can result in notable improvements in the fracture toughness of the composite. In fact, the important reinforcement parameters that directly contribute to improvements in fracture toughness of the composite are either a decrease in constraint and a lower flow stress, which would result in an increase in plastic rupture displacement of the particle, or a constrained particle exhibiting a higher maximum flow stress and a low plastic stretch. In fact, the occurrence of debonding or decohesion at the matrix-particle interfaces during crack opening facilitates a loss of through-thickness constraint, thereby, promoting extensive stretch while concurrently increasing the work of fracture. Also, the occurrence of such debonding at the interfaces minimizes the constraint (triaxiality) ahead of the crack-tip by relieving throughthickness stresses stemming from a transition from fully-constrained plane strain fracture process to an unconstrained "near-plane stress" ligament, which fractures at higher K c values. Fracture surface features of the fracture toughness specimens of the monolithic Ti--48AI alloy are shown in Fig. 7. Fracture in the matrix occurred primarily by cleavage across the equiaxed, ~ grains, with a "step-like" faceted fracture mode observed within the lamellar colonies [Fig. 7(a)]. Fracture of the Ti-48AI/TiNb composite occurred by: (a) a mixture of cleavage and intergranular fracture mode across the matrix and (b) a ductile dimpled fracture mode in the TiNb reinforcements [Fig. 7(b)] under monotonic loading conditions. This is in sharp con-
Gamma-titanium aluminide intermetallic
391
Stress Intmslty factor r a n ~ , A~ (]ksl On)u )
.......
I ,;
.......
I
. . . . . . .
~ 10"z
lO-I -O 'n-24AI-UNb 1093"C/@..~h/FAC A aI-6AI-4V mla mmmled [] aI-4VA r . ~ ,Ided St ~ R = 0.1
i
10-3
10~
A
0
10.5
[]
111-7 10' !
I I Itlil[
I
I
I
I illll
.....
104
Smm ~ . ~ j taam-m.ge,, ~ O~a (m~) Fig. 9. Comparison of the room temperature fatigue crack growth for the monolithic alloys[20].
trast to the transgranular fracture mode observed in the TiNb reinforcements [Fig. 8(b)] and the predominantly cleavage fracture mode [Fig. 8(a)] observed in the Ti-48A1 matrix under cyclic loading conditions.
Fatigue crack growth The crack growth behavior in the monolithic Ti-48AI alloy under cyclic loading conditions is shown in Fig. 9 in terms of the growth rate per cycle (da/dN) as a function of the nominal stress intensity range AK (= Kmax - Kmin). Results of mill-annealed (MA) Ti-6A1-4V [48] and ~2 + fl processed Ti-24Al-11Nb[43] are also shown in the same figure for purposes of comparison. Under cyclic loads, the cracks tend to propagate subcritically at stress intensities well below the crack initiation toughness [15 MPa(m) °5] needed to initiate cracking under monotonic loads. Therefore, subcritical crack propagation under cyclic loads is an important mode contributing to degradation of the monolithic Ti-48AI alloy. The monolithic Ti48AI alloy is observed to have lower near-threshold fatigue crack growth rates than the mill-annealed Ti-6A1-4V and at2 + fl forged Ti-24Al-llNb. However, in the Paris regime the growth rates in the millannealed Ti-6A1-4V and Ti-24Al-llNb are similar to those in the monolithic Ti-48Ai. The fatigue crack growth resistance of the monolithic Ti-48A1 alloy was generally superior to that of the Ti-24AI-11Nb, although the toughness of the gamma-based alloy [15 MPa(m) °5] [41] is inferior to the toughness of the ~2 + fl alloy [23.4 MPa(m) °'5][43]. The fatigue crack growth rate data obtained for the as-extruded Ti-48AI/TiNb composite are compared with data obtained previously for fl-TiNb [19, 48, 49] and extruded Ti-48AI alloy in Fig. 10. The crack growth rates for the composite lie between those obtained for the Ti-48 AI matrix and the fl-TiNb reinforcements. Similar results have been reported in earlier studies on HIPped TiAI + 20 vol.% TiNb (Fig. 11)[19, 48, 49]. It is evident that crack propagation rates in the TiA1/TiNb composite are higher than those in the matrix alloy, especially at nearthreshold levels, in spite of the improved fracture toughness of the composite under monotonic loading. In fact, the ambient temperature fatigue threshold of the TiAI/TiNb composite is 20%
392
W.O. SOBOYEJO et al.
s~m ~m~ty ~
rzmee,A~ (hi On)u)
1
1O ,
,
,
,
20 30
, ,,,i
,
,
,-,,I0"
lO ~ J'~'
@0 @@ •
Q
lO~
10"~ i t~ o "l'e.48Al + :~ voL% T i / ~ (As r.~a__-,d_ - No
zx 2X
o "I~.~AI + 20 ~oL% " 1 ' ~ (.~ e ~ d M - W'tl~ I I ~
:
~)
acres ~
~Ti-48A! (As-e~ruded) I
1
I
I Illl
I
10 Stress
katemity
range, AK
factor
I
L
20 3 0 4 0 5 0 (MPa
(m) ~)
Fig. 10. Comparison of the room temperature fatigue crack growth of the as-extruded gamma alloys and the fl-TiNb reinforcement[20]. 10~
I
I
|
,
|
6
,
i
,
'lrUU+ 20'~ ( ~ ~m) 1 ~ C-R {EdllP)Ori,.,,e,,a,~: R = 0.I
10~ •
104
"IIAI + T I N b
10 4
0 "r-'llAl
@@
.o .t"oo 10.7
104
o 10-w
10 •
I 2
,
Smm ~ z j t y
,
I S
,
,
,
,
10
~Etor rm~e, aK ( M ~ (m)")
Fig. 11. Comparison of the room temperature fatigue crack growth rates in HIEPped gamma alloys and fl-TiNb[17].
G a m m a - t i t a n i u m aluminide intermetallic
393
lower than in the unreinforced matrix. However, the crack growth rates in both the TiAI/TiNb composite and ~-TiAI are sensitive to the applied stress intensity range. The region of power law dependence is relatively small for both the TiA1/TiNb composite and the monolithic-~-TiAl alloy. With increasing stress intensity factor ranges (AK), a rapid increase in crack growth rate (da/dN) is observed for small changes in AK. The degradation in ambient temperature crack growth resistance is ascribed to the absence of crack bridging and interfacial debonding or crack deflection under cyclic loading conditions (Fig. 12). Since the degree of crack bridging is small, the near tip and applied stress intensities are expected to be equal, except for contributions to shielding from other competing mechanisms such as surface roughness-induced crack closure [50]. Additionally, the degradation in fatigue crack growth resistance of the TiAI/TiNb composite is most likely caused by the differences in microscopic fracture mechanisms under monotonic and cyclic loading (Figs 7 and 8) and the inelastic strains that are developed during cyclic deformation. The portioning and progressive accumulation of such strains results in a rapid degradation of the fatigue resistance of the TiA1/TiNb composite, and the resultant crack growth rates are between those of the matrix and the reinforcements. An alternative factor contributing to the degradation in cyclic crack growth resistance is the influence of ductile phase reinforcements on crack-tip opening displacement (CTOD), which is assumed to be directly proportional to crack-tip extension under cyclic loading. The plane strain crack opening displacement (COD) under cyclic loading is: At$ = ~max - - t~min =
(AK)2/2Etrys,
(7)
where 6 is the crack-tip opening displacement, and the subscripts "max" and "min" correspond to the maximum and minimum stress intensity factors. Assuming the extent of irreversibility to be the same for the monolithic Ti-48AI alloy and TiA1/TiNb composite, the increase in cyclic crack growth rates of the ductile phase-reinforced composite is ascribed to an absence of cracktip shielding from bridging. In fact, it is the fatigue-induced failure of the reinforcing TiNb particles that occurs before the advancing crack can establish a bridging zone. This is aided by the low stress intensities associated with cyclic loading and the concomitant small crack-tip opening displacements hinder the formation of a bridging zone. Fatigue crack propagation surfaces revealed transgranular fatigue crack growth in both the Ti-48AI monolithic alloy and the TiA1/TiNb composite. There was no evidence of interfacial decohesion and plastic stretching of the TiNb strips under cyclic loading [Fig. 12(b)] nor was there visible evidence of cracking of the TiNb particles ahead of the propagating crack tip [Fig. 12(a)]. Such an observation is consistent with similar observations on the fatigue and fracture behavior of ductile/brittle material systems, i.e. metal-ceramic interfaces. In these systems fracture predominates in the brittle ceramic phase under monotonic loading due to its much lower fracture toughness, whereas under cyclic loading the crack path is focussed in the metal [51]. The attraction and retention of the crack in the metal phase occurs because the metal matrix presents a region where strain can be either accumulated or localized thereby influencing progressive crack advance. This confirms the enhanced susceptibility of the metal matrix to failure by fatigue. The monolithic and cyclic fatigue tests demonstrate that the microstructure of the TiA1/ TiNb composite engineered for superior fracture toughness under monotonic loading fails to result in improved cyclic crack growth resistance. In spite of strong reinforcement (TiNb) matrix interfaces and a low strain hardening of the metallic phase, the occurrence of crack bridging and crack-tip blunting by the TiNb reinforcements significantly enhances monotonic fracture toughness. However, failure of the strong metal-reinforcement interfaces to provide impedance to fatigue crack growth under cyclic loading limits contributions from crack-tip shielding due to bridging. CONCLUSIONS The following are the observations of a study on the influence of ductile phase reinforcement on the fatigue and fracture behavior of a gamma-TiA1 intermetallic. (1) Increase in test temperature caused a loss of strength for both the monolithic alloy and EFM 56/3---D
394
W.O. SOBOYEJO et al.
the composite counterpart. The extent of degradation was greater in the ductile phase reinforced composite. The loss of strength resulted in appreciable improvements in ductility. (2) Under monotonic loading significant improvements in room temperature fracture toughness of the brittle gamma-based Ti-48A1 intermetallic is achieved through the introduction of ductile TiNb particles. Toughening and resistance curve behavior of the TiAI/ TiNb intermetallic composite arises from the conjoint and mutually interactive influences of crack bridging and blunting by the thick (> 100/Am) TiNb strips. The thinner TiNb strips are fractured during stable crack growth under monotonic loading, and consequently they do not participate in either the bridging or the blunting processes. (3) Ductile phase toughening was found to be less effective under cyclic loading. This is attributed to mutually interactive influences of an absence of crack bridging, the effects of the ductile TiNb reinforcement on crack-tip opening displacement under cyclic loading, and the accumulation and portioning of the inelastic strains during cyclic deformation. The accumulation of inelastic strains during cyclic loading promotes subcritical fatigue damage in the composite microstructure.
Acknowledgements--WOS and FY would like to thank The National Science Foundation (Grant No. DMR 9458018) for financial support. Appreciation is also extended to Dr Bruce McDonald at The National Science Foundation for his encouragement and support of this research. One of the authors (TSS) gratefully acknowledges the State of Ohio: Board of Regents and The University of Akron for supporting this research study, through a faculty research grant.
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(Received 10 May 1995)