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An investigation of the effects of ductile phase reinforcement on the mechanical behavior of advanced high temperature intermetallics
Materials" Science and Engineering, A 171 (1993) 95-104
95
An investigation of the effects of ductile phase reinforcement on the mechanical behavior of advanced high temperature intermetallics* W. O. Soboyejo Department of Materials Science and Engineering, The Ohio State University, 116 West 19th A venue, Columbus, OH 43210-1179 (USA)
S. M. L. Sastry Department of Mechanical Engineering, Washington University, 1 Brookings Drive, St. LouL~, MO 63130 (USA) (Received January 27, 1993; in revised form April ! 5, 1993)
Abstract The effects of ductile phase reinforcement on the mechanical behavior of Ti-48AI reinforced with 20 vol.% TiNb and MoSi 2 reinforced with 20 vol.% Nb are reported. Ductile phase reinforcement is shown to promote improved fracture toughness and a reduction in fatigue crack growth resistance. The role of crack/tip shielding by bridging and deflection mechanisms is also modeled, and the effects of temperature on the bend and tensile strengths are assessed.
1. Introduction
Recent efforts to develop hypersonic vehicles have stimulated research on advanced high temperature structural intermetallics for potential applications in the temperature range between 700 and 1500°C. In particular, y-based titanium aluminides have been considered for intermediate temperature (700-1000°C) applications [1, 2], while MoSi2-based systems have been proposed for higher temperature uses in the temperature range between 1000 and 1350°C [3-6]. The selection of these systems has been due to their relatively low density (compared with Ni-based superalloys), and strength retention in the stated temperature regimes [1-6]. Unfortunately, however, both the Y titanium aluminide and MoSi 2 based systems are limited by their room-temperature damage tolerance, and there is a clear need for improved fracture toughness and fatigue resistance in both systems. There are essentially two approaches for engineering damage tolerance improvements in y-based titanium aluminides and MoSi2. The first approach relies on intrinsic modification via alloying and variations in thermomechanical processing schedules, and this has been used successfully to engineer toughness *Paper presented at the ASM Symposium on High-Temperature Aluminides and Intermetallics, San Diego, CA, September 16-19, 1991. 0921-5093/93/$6.00
improvements and in y-based titanium aluminides [7-10]. The second approach utilizes brittle or ductile reinforcements to increase fracture toughness via crack-tip shielding mechanisms [11-24]. The latter "composite" approach has been applied, with varying degrees of success, to the extrinsic toughening of ybased titanium aluminides [ 11 - 15] and MoSi 2 [ 15-25]. In particular, ductile phase reinforcement of y alloys with Nb or TiNb has been shown to promote significant (200%-300%) improvements in the work of rupture WR (a measure of toughness), and modest improvements (24%) in the fracture toughness K o due to crack-tip shielding by crack bridging mechanisms [11-13]. Similar improvements in WR and K o have been reported for MoSi 2 reinforced with Nb [24, 25]. Brittle reinforcement of MoSi2 with 20 vol.% SiC whiskers [19] has also been shown to improve the fracture toughness from about 5.3 to 8.2 MPa m ~/2, while particulate reinforcement with brittle TiC has been shown to have little effect on fracture toughness
[2o[. In the current paper, we report the results of recent efforts to engineer improvements in the damage tolerance of a y-based titanium aluminide alloy (Ti-48AI) and MoSi 2 (all compositions are quoted in atomic per cent unless stated otherwise). The effects of ductile phase reinforcement on strength and fracture toughness are rationalized using micromechanics models. Ductile phase reinforcement is also shown to © 1993 - Elsevier Sequoia. All fights reserved
96
W. O. Soboyejo, S. M. I,. Sastry
/
degrade fatigue crack growth resistance, in spite of the toughening that is achieved largely via plastic deformation of the ductile phases and the deflection of cracks around the reinforcements under monotonic loading. The use of linear superposition in the prediction of the combined effects of multiple toughening mechanisms is also elucidated.
2.
Material
The MoSi 2 powder (99.5°/,,, - 3 2 5 mesh) was obtained from Cerac, Milwaukee, WI, and the Nb powder ( - 35 to 80 mesh) was procured from Nuclear Metals, Inc., Concord, MA. The composition of the powders is given in Table 1. The Cerac powder was found to be molybdenum rich relative to stoichiometric M o S i 2 which contains 36.92 wt.% Si. M o S i 2 powders were poured into an Nb can, degassed and sealed by electron beam welding. The can was then hot isostatically pressed at 1700 °C for 4 h under 207 MPa argon pressure. The resulting compact was then cooled slowly (10 h) to room temperature to minimize thermal gradients and cracking. The equiaxed microstructure of the resulting MoSi, compact is shown in Fig. 1. The mean linear intercept grain size was found to be 14/~m, which is consistent with the size of the original powder. This indicates that very limited grain growth occurred during hot isostatic pressing. The density of the material was approximately 97% of the theoretical density, and this is manifested by the incidence of "dark" pores on the photomicrograph shown in Fig. 1. Silicon and SiO, particles (Fig. 1) were also revealed in the monolithic MoSi 2 using energy dispersive X-ray (EDAX) spectroscopy. EDAX analysis revealed that many of the apparent w)ids contained remnants of these particles which were pulled out during metallographic specimen preparation. The Nb-reinforced MoSi2 was produced by phase blending and hot isostatic pressing of nominal 20 voi.% spherical niobium particles (approximately 250/~m in diameter) and MoSi: powder. No SiO2 contamination
Effects o f ductile phase reinforcement
was detected in the Nb by EDAX analysis, and the same processing schedules that were used for the consolidation of monolithic MoSi 2 were employed in the fabrication of the MoSi2/Nb composite. The resulting composite microstructure is shown in Fig. 2. The matrix microstructure (Fig. 2(a)) is similar to that of the monolithic MoSi, (Fig. 1) which had a density of about 97%. However, in contrast, the Nb reinforcements exhibit an almost fully dense structure (Fig. 2(b)) with some incidence of microcracking. A thick (approximately 40 /~m) reaction zone was observed between the Nb reinforcements and the MoSi 2 matrix (Fig. 2(c)). This has been shown to consist of layers of NbsSi 3 and (Mo,Nb)~Si3 [24]. The actual layered interface compositions obtained by semi-quantitative EDAX techniques are shown in Table 2. The Ti-48A! was procured from Nuclear Metals, Inc., Concord, MA, in the form of - 35 mesh (average diameter, 5 0 0 x 10-6m) powder. The powder was produced by the plasma rotating electrode process. The monolithic Ti-48AI powder was canned in T i - 6 A I - 4 V and extruded at a die temperature of 1343 °C with a reduction ratio of 14:1. This resulted in a duplex microstructure with 18 ,um a 2 + 7 lamellar packets and 16 /~m equiaxed }, grains. The Ti-48Al + 20 vol.% TiNb composite was produced by mixing - 35 mesh Nb and Ti-48Al powder, and processing under the same conditions that were employed for monolithic Ti-48AI. This resulted in a duplex matrix microstructure and non-uniform elongated strips of TiNb, as shown in Fig. 3. A small reaction layer was also observed at the interface between the TiNb and the Ti-48Al matrix (Fig. 3). This has been shown by previous workers [14] to consist predominantly of a(2). EDAX compositions of the TiNb
~
. ..
,,
..-
2
~.
o
. . ~, : ~
:r,-
,~.
, .o'
,
)
.~, -.~.
z~r:.t..,,
.;.%
•
---
j4~/¢
o ~. , ¢x.;: .
.
,,.~.~.
..°~.." ~ , ~ . , ~
_
•
.... 1 1
"#"
,~
.
;" . -
TABLE 1. Chemical composition in weight per cent MoSi, Si Fe Nb O C Mo
Ti-48AI 34.99 0.041 < 0.05 0.54 0.O4 Balance
AI Si H N O S "I'i
34.37 0.024 0.001 0.005 0.095 O.OO8 Balance
Fig. 1. Scanning electron photomicrograph showing 97% dense structure of MoSi2 with equiaxed grains and MoSi~Si3 and SiO, particles.
W. O. Soboye]o, S. M. L. Sast~'
,,,¢
/
Effects of ductile phase reinforcement
97
~,L
(b)
lain
Fig. 3. Photomicrographs of y-based titanium aluminides taken under polarized light: (a) monolithic Ti-48AI, and (b) Ti-48AI/ 20 vol.% TiNb and composite.
TABLE 3. Chemical compositions of components of as-extruded Ti-48AI/20vol.%TiNb composite determined by EDAX
L~J lOI.tm
(c)
Fig. 2. Scanning electron photomicrographs showing typical features of the MoSi,/Nb composite: (a) matrix structure. (b) niobium structure, and :el structure of layered Nb-MoSi, interfttce~.
Region
Matrix Interface Reinforcement
Composition (at.%) Ti
AI
Nb
48.5 48.1 57.5
51.5 51.5 15.3
-0.4 27.2
TABLE 2. Chemical compositions fin atomic per cent) of components of MoSi2/2(Ivol.%Nb composite determined by EDAX Region
Nb reinforcement Internal interface Middle interface External interface MoSi, matrix
Composition (at.%) Mo
Si
0.8 1.8 37.2 29.3
0.8 41.6 39.9 70.7
Nb 100.0
98.4 56.6 22.9
particles, interfacial regions, and the T i - 4 8 A 1 matrix are shown in Table 3. Note that the actual interfacial composition observed in this study indicates the presence of a(2) and 7 phases at the interface. It is also important to note here that monolithic/matrix MoSi, were not extruded for direct comparison since the facilities required for extrusion at the required temperatures are not currently available.
98
W. O. Soboyejo, S. M. L. Sastry /
3. Experimental procedures Bend tests were performed on MoSi 2 and the Nb/ MoSi 2 composite at temperatures between 25 and 1500 °C. The tests were conducted on flexure specimens with square (6.35 mm x 6.35 mm) cross-sections, and the specimens were loaded monotonically to failure under three-point bend conditions. The ductileto-brittle transition temperatures (DBTT) for MoSi 2 and Nb/MoSi 2 were thus established from the resulting bend strength data. Similarly, the tensile properties of Ti-48Al and the Ti-48Al/TiNb composite were measured at temperatures between 25 and 982 °C, and DBTT were established for the monolithic and composite y-based alloys. Duplicate tests were conducted on smooth cylindrical button-headed specimens with a gage diameter of approximately 3.18 mm. The specimens were loaded monotonically to failure at a strain rate of 5 x 10-4 S - 1 The fatigue-crack growth behavior of the monolithic alloys and the composites were investigated using 25.4 mm long single edge notched (SEN) bend specimens with square (6.35 mm x 6.35 mm) cross-sections. The tests were conducted in laboratory air under threepoint bend constant amplitude loading after precracking under far-field compression [7, 8] loading. A cyclic frequency of 10 Hz was employed. Crack growth was monitored using a high resolution (2.5 /~m) telescope connected to a video monitoring unit, and an initial load-increasing scheme was employed until crack growth was detected after about 106 cycles. The load range (AP=P(max)-P(min)) was then maintained constant until the end of the test. The tests were stopped prior to specimen fracture. Metallographic examination of the interaction of the crack path with the local microstructure was then carried out, and the micrographs were used to determine the deflection angles. Room temperature fracture toughness tests were also conducted on SEN specimens that were precracked under far-field compression loading [7, 8]. 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 t/2 s - i.
Effects of ductile phase reinforcement
reinforcement-matrix interface and the constrained plastic deformation of the ductile phase during loading to failure. However, in the absence of more detailed analysis, the use of simple rule-of-mixtures is explored in the current investigation. The improvement in toughness due to ductile phase reinforcement can be assessed in to ways, i.e. in terms of G the fracture energy, or in terms of the fracture toughness K o. Improvements in G due to ductile phase reinforcement can be estimated using composite mechanics. This yields [26-28] Gc=(1 -f)Gm +fG r
(2a)
or It*
G~=(1 -f)Gm +f J o(u)du
(2b)
tl
where G m and Gr are the respective fracture energies of the matrix and reinforcement, o(u)is the polynomial expression relating stress to strain in the reinforcement, u is the crack opening displacement (COD) and u* is the COD at the point of fracture. Rearranging eqn. (2b) yields
age = G¢- Gm+ f j o(u)du
(3)
0
Equation (3) represents the toughness increase due to plastic stretching and necking down of the ductile phase reinforcement to a point when u = u*. Further non-dimensional analysis [26-28] reveals that AG¢ is given by AGc =foyrX
(4)
where oy is the yield stress of the ductile phase reinforcements, r is the characteristic reinforcement radius, and X is a dimensionless work of rupture parameter which is given by [28] u*/r
X =f J (O/Oy)d(u/r)
(5)
0
4. Micromechanical modeling The strength oc of a ductile phase reinforced composite system can be estimated using simple rule-ofmixtures. This yields for constant strain conditions
oc=Om(1-f)+ for
(1)
where Om is the matrix strength, f is the reinforcement volume fraction, and or is the reinforcement strength. Note, however, that this simple approach neglects the
where o is the applied stress. Typical values for X are 1.5 for 7-TiAI reinforced with 20 vol.% TiNb and 6.3 for MoSi 2 reinforced with 20 voi.% Nb [ 12, 24]. The fracture toughness can also be expressed in terms of the stress intensity factor at the point of fracture instability, i.e. K o or KIc. For a ductile phase reinforced composite, the toughening components due to potential crack bridging and deflection mechanisms can be assessed using existing analytical models [27, 29]. The toughening due to crack bridging can be assessed using the following expression due to
W. O. Soboyejo, S. M. L. Sastry
/
Budiansky et al. [27]: 2h-
K K,~
- 1 + 2(21~)'i2[foy(LIKm'i2l
m
~
1 cos2(~t/2) .
.
.
.
(6)
.
(7)
The combined effects of crack bridging and deflection can thus be assessed by applying the principle of superposition [30, 31 ]. This yields K = 2b2,1Km
99
TABLE 4. Effects of temperature on the compression strengths of MoSi, and MoSi~ + 20vol.%Nb
where )ch is the toughening ratio due to crack bridging, K is the actual stress intensity factor at the tip of the bridged crack, Km is the matrix stress intensity factor/ toughness, and L is the characteristic length of the reinforcement. Similarly, the toughening due to crack deflection through an angle ~t, 2d, can be estimated using the following expression due to Suresh [29]: K 2,,= K,,
Effects of ductile phase reinforcement
(8)
Equation (8) can be used to estimate the K o or Kjc of a ductile phase reinforced composite in which crack bridging and deflection mechanisms are effective. However, it is important to study the crack paths in the composite system before assuming that crack-tip shielding occurs by the above mechanisms. For the systems discussed in the current paper, the analysis of crack bridging and deflection will be shown to be important in the estimation of the effective driving force for crack growth. However, for other systems, crack-tip blunting by the ductile phase may also contribute to crack-tip shielding [30, 31], and its effects may have to be analyzed when crack growth occurs across the ductile phase.
5. Results and discussion 5.1. M o S i 2 b a s e d systems
The bend strengths of the hot isostatically pressed (HIPed) MoSi2 and MoSi2/Nb composite are compared in Table 4. Both the monolithic MoSi., and the MoSi2/ Nb composite had very high strengths up to 1250 °C, and the strengths of the composite were lower than those of the matrix, as expected from simple rule-ofmixtures. Significant softening was observed in both systems above 1250°C. This is consistent with the results from previous studies on monolithic MoSi, single crystals [32]. The retention of strength in the matrix up to the DBTT has been attributed to a KearWilsdorf type cross-slip mechanism similar to that observed in Ll(2) ordered alloys. The ductility improvement in single crystals of MoSi2 is also associated with increased mobility and climb potential of
Temperature (°C)
MoSi 2 (MPa (ksi))
MoSi 2 + 20vol.%Nb (MPa (ksi))
1050 115(1 1250 1350 1450
594 (86.2) 493 (71.4) 31)5 (44.3) 222 (32.2) 65 (9.4)
513 392 286 156 25
(74.4) (56.9) (41.5) (22.7) (3.6)
(100) and (110) ordinary dislocations in addition to {110}(331) and/or {013}(331) slip which occur both above and below the DBTT. The measured strength of the MoSi2/Nb composite at room temperature (583 MPa) is underpredicted by simple rule-of-mixtures to be 519 MPa. However, it is important to note that the accuracy of simple rule-ofmixture predictions is limited by matrix and powder contamination by interstitiais, porosity, and failure to account for the reinforcement-matrix interactions. More detailed analysis of the constraint associated with the deformation of the Nb reinforcements is also needed for improved strength predictions. Nevertheless, the use of simple rule-of-mixtures provides some useful insights into relative differences between the strengths of the monolithic MoSi, and the MoSiffNb composite. It is of interest to examine the effects of ductile phase reinforcement on the fracture toughness of the MoSiffNb composite. The monolithic MoSi 2 had a fracture toughness of 4.6 MPa m"-" compared with a measured toughness of 5.7 MPa m ~;2 obtained for the composite. This represents a modest increase of about 24% in the fracture toughness due to ductile phase reinforcement. Such a modest increase in fracture toughness appears at first glance to be inconsistent with previous reports of toughness improvement where toughness was characterized by the work of rupture WR [24]. However, although previous studies predicted and measured a work of rupture improvement of about 335% in MoSi2 reinforced with 20 vol.% Nb, they did not consider the fracture toughness in terms of the stress intensity factor at the point of fracture instability Ko •
The increase in WR and Ko in MoSi 2 + 20 vol.% Nb due to ductile phase reinforcement can be predicted using composite mechanics. The increase in WR is given by AG, in eqn. (4). By substituting f = 0 . 2 , oy = 200 MPa, r = 80/~m, and X = 6.3 into eqn. (4), the increase in A G c reported in ref. 24 is predicted to be 202 J m -2. The composite fracture energy G,. is thus computed from eqn. (2) to be 262 J m 2, since the
1O0
W.O. Soboye]o, S. M. L. Sastry /
Effectsof ductile phase reinforcement
Fig. 4. Stable fatigue precrack obtained by far-field compression cyclic loading.
~K O~i./~m. ) 0.5
10 -2
1 i i [ I I
I
I
5 I I I i t
I
10-~
I 10-4
10 -a
10_5 I ~ 10 ~
Fatigue crack growth rata, 10 -5 da/dN 10 ~
da/dN
!
; 10 -7
(in./cycle)
I lO s
(mm/cycle) 10_7
f l0-~
10~ 10 -9
,
0.5
,
,
,
~
,
,
i
,
1 5 Suess intensity range. AK (MP~/-m-)
,
,
~
~
%
. *~1~'-" l i ~
Lif
:ll
)10 ~10 10
Fig. 5. Fatigue crack growth rate data obtained for MoSi2/Nb composite at room temperature (R=Km,JKm~=O.I)(from ref. 25).
matrix fracture energy Gm is approximately 60 J m-2 Similarly, the increase in K o can be predicted for the MoSi2/Nb composite using composite mechanics. The prediction involves the superpositiop of toughening components due to crack bridging (L = 2 r = 160/~m, ;tb=l.19 ) and deflection ( ¢ = 2 0 °, 2d=1.03 ). The respective toughening ratios, 2b and ;tj, are then substituted into eqn. (8) to obtain predictions of K o. The predicted fracture toughness of 5.7 MPa m 1/2 is in exact agreement with the measured fracture toughness of 5.7 MPa m 1/2. The differences between the degree of toughening indicated by changes in WR and K o can therefore be rationalized using existing composite mechanics analyses. Note that the potential effects of SiO2 contamination are not assessed in the above analysis. Attempts to obtain stable fatigue crack growth rate data from precracked SEN specimens were unsuccessful under three-point loading in spite of successful precracking of the monolithic material under far-field compression (Fig. 4), i.e. stable crack growth was not observed in the monolithic MoSi 2 for A K levels with Kmax < Ktc, and the SEN specimens of the monolithic
Fig. 6. Photomicrographs showing the interaction of the fatigue crack with the microstructure of the MoSi2/Nb composite.
MoSi 2 fractured under cyclic loading when Kmax>t K o. However, stable fatigue crack growth was observed in the MoSi2/Nb composite after precracking in compression (Figs. 5 and 6). The occurrence of subcritical fatigue crack growth in the MoSi2/Nb composite is attributed to the inelastic strains that are developed in the relatively brittle layered interfaces (Figs. 2(c) and 2(d)). The intermetallic compounds in these interfacial
W. O. Soboyejo, S. M. L. Sastry
/
regions can provide the "'weak link" for initial fracture and damage accumulation under cyclic loading conditions. It is also possible that the matrix MoSi 2 is further embrittled by interstitial elements which are not revealed by the EDAX analysis shown in Table 2. Stable subcritical fatigue crack growth in the MoSi2/ Nb composite occurs at stress intensity factor ranges as low as 1.7 MPa m l.'-~,and the crack growth rates can be expressed as a function of A K, the stress intensity factor range, using a simple Paris power-law da/d N vs. AK relationship. However, as in ceramics [32, 33] the power-law exponent is very high (approximately 14) compared with typical values between 2 and 4 in conventional metallic materials [34]. Limited crack-particle interactions were observed under cyclic loading, and the cracks generally tended to avoid the Nb particles by deflection around the layered interfaces between the Nb particles and the MoSi~ matrix (Fig. 6). Nevertheless, the crack-tip shielding provided by crack deflection [29] was not very significant since the deflection angles under cyclic loading conditions were typically between 45 ° and 60 °. Such deflection angles result in deflection toughening ratios ;t~ between 15% and 25% over an equivalent linear crack (see eqn. (7)). The absence of crack bridging under cyclic loading (Fig. 6)also partly explains the limited resistance of the MoSi~ matrix to fatigue crack growth. Crack-tip shielding contributions from bridging and plastic stretching of Nb particles can be significant for material with 20 vol.% Nb, as discussed earlier. The differences between the resistance of the MoSi2/Nb composite to crack growth under monotonic and cyclic loading can therefore be rationalized by considering (i) the effects of loading and microstructure on the crack path, (ii) the effects of crack path on crack-tip shielding, and (iii) the effects of inelastic strains due to the accumulation of damage at the MoSi~-Nb interfaces during cyclic loading. The current results therefore suggest that improved fracture toughness, K~,., in the model MoSi2/Nb composite does not necessarily imply improved fatigue crack growth resistance. The results also show that improvements in the work of rupture computed by integrating the expression relating reinforcement stress
Effects of ductile phase reinfi>rcement
101
to strain do not necessarily imply similar levels of fracture toughness improvement. Durability and damage tolerance assessment in the MoSi2/Nb composites therefore requires a careful analysis of both fracture toughness and fatigue behavior. It is also encouraging to note that the strength and fracture toughness of the MoSi2/Nb composite can be estimated using existing analyses, and that the MoSi 2 and MoSi2/Nb model composite systems examined here appear to have the potential for 1250-1350°C applications owing to their exceptional elevatedtemperature strength retention. However, further system optimization is required to improve the oxidative stability and damage tolerance capability of MoSi2 and its composites for structural applications. 5.2. Duplex y-based titanium aluminide systems The tensile strengths of the as-extruded Ti-48AI and Ti-48AI-20vol.%TiNb are compared in Table 5. Both the monolithic and composite materials retained their relatively high strengths at temperatures up to 815 °C. Significant degradation in strength and increased plasticity was observed at 982 °C where the composite had a plastic elongation to failure of 84.5°/,,. Reasons for the large increase in ductility are not fully understood at present. However, it is important to note that previous studies have attributed the increased ductility of the monolithic y alloys to the activation of multiple slip systems [35, 36] and transitions in the micromechanisms of fracture that occur at about 650-700 °C 17, 34, 35] and about 900-982 °C [7]. The predicted rule-of-mixture room temperature strength of the Ti-48AI/TiNb composite is 563 MPa (Om=594 MPa [7] and o~=440 MPa [15]). This is within 10% of the measured value of 523 MPa, in spite of failure to account for the reinforcement-matrix interface and constraint effects associated with thc deformation of the TiNb reinforcements. The fracture toughness K~ of the as-extruded Ti-48AI was increased from approximately 15 MPa m ~'-' to approximately 20 MPa m ~'-" by ductile phase reinforcement with 20 vol.% TiNb. This modest increase was predicted by the superposition of the combined effects of crack bridging and deflection
TABLE 5. ('omparison of the tensile properties of as-extruded T i - 4 8 A I and T i - 4 8 A I ~- 20 vol.% TiNb Temperature (°C~
25 700 815 982
Yield stress (MPa)
Ultimate tensile stress (MPa)
Plastic elongation to failure (%1
y
y+TiNb
y
y +TiNb
I,
y -TiNb
594 483 476 262
--394 117
686 579 586 317
623 461 464 132
1.7 2.8 6.(I 20.0
--2.4 84.5
102
W. O. Soboyejo, S. M. L. Sastry /
described in Section 4. By substituting in the matrix fracture toughness and the respective toughening ratios due to bridging (2,= 1.3) and deflection (2j= 1.03) into eqn. (8), the fracture toughness of the Ti-48AI/ 20vol.%TiNb composite is predicted to be 20.1 MPa m u2. This is in excellent agreement with the measured fracture toughness of 20 MPa mU2 in spite of failure to account for the strong a2 interface in the analysis. Note that f = 0 . 2 , Oy=440 MPa, L = 1 0 0 0 # m (average length of the TiNb reinforcements which was equated with the bridging length) and 4 = 2 0 ° (based on measurements of crack angles at the sides of specimens). Also, the good agreement between the measured and predicted fracture toughness values confirms the important role of crack bridging reported by previous workers [11-15] for monotonic loading conditions. As observed in the MoSi2/Nb composite, the improvements in the fracture toughness Ktc due to ductile phase reinforcement were modest compared with work of rupture WR or A G c values reported by previous workers [12]. By substituting in f = 0 . 2 , % = 440 MPa, r = 25 /~m (average width of the TiNb reinforcements), and X = 1.5 into eqn. (4), AGc or WR is estimated to be 44 J m- 2 in the Ti-48AI/TiNb composite. This is comparable with the matrix toughness and the toughness of polycrystalline oxides [ 12]. However, the increase of about 100% in AGc or WR is significantly less than the increase of 33% observed in the Ktc of Ti-48AI after ductile phase reinforcement with 20 vol.% TiNb. The use of WR or A G c as toughness parameters in ductility phase reinforced systems must therefore be done with caution since they clearly appear to exaggerate the extent of toughening in cracked structures where Kt~ values provide a better
Effectsof ductile phase reinforcement
characterization of the fracture critical conditions in brittle materials. The fatigue crack growth rate data for the monolithic Ti-48AI are compared with those of the Ti-48AI/TiNb composite in Fig. 7. As for the MoSi2/ Nb composite, the fatigue crack growth resistance of the composite was worse than that of the monolithic matrix alloy. This was due partly to the absence of crack-bridging under cyclic loading conditions, as shown in Fig. 8 which reveals the transgranular fatigue crack path across the matrix and the TiNb reinforcements. Similarly, the contributions to shielding from crack deflection are small owing to the relatively small angles of deflection (see eqn. (7) and Fig. 8). The
Stress intensity factor range. ~,K ( k s i ~ ) 1 10 20 304050 10_4 10-31~-k- t i i , ,Inll ~''
•
u
i0_ s
Fatigue-crack growth rate. da/dN 10-4
Fatigue-crack growth rate. da]dN
(ram/cycle)
10 ~ (in./cycle) 10-5
I
I
I
I l I I l I
I
I
I
I I0-~
10 20 30 4050 Sta'ess intensity factor range. AK (MPaq'm)
Fig. 7. Comparison of fatigue crack growth rate data obtained for monolithic Ti-48AI and the Ti-48AI/TiNb composite, e, Ti-48AI + 20 vol.% TiNb (as-extruded--no TiNb strip across crack path); t~, Ti-48AI-20 vol.% TiNb (as-extruded--with TiNB strips across crack path); , Ti-48AI (as-extruded).
Fig. 8. Photomicrographshowingthe interaction of the fatigue crack path with the microstructure of the Ti-48AI/TiNb composite.
~ O. Soboyejo, S. M. L. Sastry
/
fatigue crack growth rates in the T i - 4 8 A I / T i N b composite must therefore represent the "composite" fatigue resistance of the T i N b reinforcements, the ct 2 interface, and the T i - 4 8 A I matrix since the crack growth rates lie between those obtained for T i N b [14] and the T i - 4 8 A I matrix [37]. However, the partitioning of strains and the accumulation in the different sections of the composite cannot be explained so easily since no premature cracking was observed in the T i N b (the c o m p o n e n t with the poorest fatigue resistance) strips ahead of the crack-tip. T h e r e is clearly need for m o r e work in this area. Nevertheless, it is possible to conclude from a practical standpoint that for the two model composite systems examined in this paper, i m p r o v e d toughness due to ductile phase reinforcement does not imply improved resistance to fatigue crack growth.
6. Conclusions T h e effects of ductile phase reinforcement on the mechanical behavior of T i - 4 8 A I and MoSi, were investigated using T i - 4 8 A I / T i N b and MoSi2/Nb model systems. T h e following conclusions were reached. (1) T h e increase in fracture toughness, K o or Kt~, due to ductile phase reinforcement ( 2 4 % - 3 3 % ) is significantly less than the increase in the work of rupture (100°/,,-337%) due to ductile phase reinforcement. Claims of toughness i m p r o v e m e n t in ductile phase reinforced systems may therefore be misleading when toughness is characterized by the work of rupture. (2) T h e i m p r o v e m e n t s in fracture toughness, K o or K~. observed under monotonic loading are associated with a degradation in fatigue crack growth resistance since crack bridging does not occur under cyclic loading conditions. Extrinsic toughening by ductile phase reinh)rcement may therefore degrade damage tolerance in structures subjected to cyclic loading. (3) T h e loss of strength associated with ductile phase reinforced systems can be estimated using simple rule-of-mixtures at temperatures below the DBTT. Also, the D B T T occurs at similar temperatures in the matrix and composite systems, although the extent of softening is greater in the ductile phase reinforced composites.
Acknowledgments T h e authors would like to acknowledge useful discussions with Professor R. O. Ritchie, Dr. K. T. Venkateswara Rao, and Dr. P. J. Meschter. Apprecia-
Effects of ductile phase reinforcement
103
tion is also extended to D. J. Deuser, B. A. Abbott, J. D. Keyes, J. J. Evans, II, J. E. O'Neal and R. J. Lederich for their technical assistance. T h e work on MoSi~ and MoSi, + 20 vol.°/,, Nb was supported by the US Air Force Office of Scientific Research under Contract No. F 4 9 6 2 0 - 9 0 - C - 0 0 3 0 with Dr. Alan H. Rosenstein as Contract Monitor. T h e work on T i - 4 8 A I and T i - 4 8 A I + 20 vol.% T i N b was conducted under the McDonnell Douglas Independent Research and Development Program.
References 1 H.A. Lipsitt, Titanium aluminides--future turbine materials. In S. M. Allen, R. M. Pelloux and R. Widner (eds.), Advanced Itigh- Temperature AIloys: Processing and Properties, American Society for Metals, Metals Park, OH, 1985, pp. 157-164. 2 Y.-W.Kim, J. Met., 41 (1989)24-30. 3 J. B. Berkowitz-Mattuck, M. Rossetti and D. W. Lee, Metall. Trans., / (1971)) 479-483. 4 P. J. Mcschter and D. S. Schwartz. J. Met., 41 (11:(19891 52-55. 5 E. Fitzer, O. Rubisch, J. Schlichting and I. Sewdas, Sci. Ceram., 18 (1973) 6. 6 J. Schlichting, High Temp. High lYessures, 10 (1978) 241-269. 7 W. O. Soboyejo, D. S. Schwartz and S. M. L. Sastry, Metall. Trans. A, 23 (1992) 2039-2059. 8 W.O. Soboyejo, in Barlan and Dickson (eds.), Fatigue Crack growth in Gamma-based Titanium Aluminides, I¥oc. I'A TIGUE '93, An Int. Con[. on Fatigue, Montreal ('anada, 1993, EMAS, Warley, UK, 1993, vol. II. 1031 - 1(136. 9 Y.-W.Kim and I). M. Dimiduk, J. Met.. 43 ¢8! ( 1991 ', 40-47. 10 D. E. Larsen, M. L. Adams, S. I,. Kampe, L. Christodoughi and J. I). Bryant, Scripta Metall. ;~huer.. 24 ( 1990', 851. I 1 ('. K. Elliott, G. R. Odctte, G. E. l,ucas and J. W. Shcckherd. in F. D. l,cmkey, A. G. Evans, S. G. Fishman and J. R. Strife (cds. I, Iligh "l~,mperature/ltigh-Pe~ornumce ('omposites. Mater. Res. Soc. ,~[vmp. I¥oc., 120 ( 1988 ) 95. 12 1t. E. I)cve, A. G. Evans, G. R. Odcttc, R. Mehrabian. M. L. Emiliani and R. J. tlecht, Acta Metull. Mater., 38 11990) 1491. 13 G. R. Odcne, H. E. l)cve, C. K. Elliott, A. Harigowa and G. E. l,ucas, in R. J. Arsenault, R. Y. Lin, G. P. Martins and S. G. Fishman (cds.), Interfaces in ('eramic Metal ('omposites. TMS-AIME, Warrcndale, PA, 1990, p. 443. 14 K. T. Vcnkateswara Rao, G. R. Odcttc and R. O. Ritchic, Acta Metall. Mater, 40 (1990) 353-36 I. 15 W.O. Soboyejo. K. T. Venkateswara Rao, S. M. L. Sastry and R. O. Ritchie, Metall. Trans. A, 24 (1993) 585-600. 16 F. I). Gac and J. J. Pctrovic. J. Am. (erum..~)c., 68(1985', ('-200. 17 W. S. Gibbs, J. J. Petrovic and R. E. Honnell, ('emm. Eng. Sci. Proc., g~(198 7 ) 645. 18 D. H. Carter, J. J. Petrovic, R. E. Honnell and W. S. Gibbs, Ceram. Eng. &'i. Proc., 10(1989) 1121. 19 D.H. Carter and G. F. Hurley, J. Am. ("eram. Sot., 70 ( 1987'1 C-79-8. 20 J.-M. Yang, W. Kai and S. M. Jeng, Scripta MemlL, 23 (1989) 1953-1958.
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/
21 J.J. Petrovic and R. E. Honnell, J. Mater Sci. Lett., 9(1990) 1083-1084. 22 J. J. Petrovic and R. E. Honncll, J. Mater. &q., 25 (1990) 4453-4456. 23 J.-M. Yang and S. M. Jeng, J. Mater Res., 6 ( 1991 ) 1-9. 24 T, C. Lu, A. G. Evans. R. J. Hecht and R. Mehrabian, Acta Metall. Mater., 39 ( 1991 ) 1853-1862. 25 K. T. Venkateswara Rao, W. O. Soboyejo and R. O. Ritchie, Metall. Trans. A, 23 (1992) 2249-2257. 26 A. G. Evans and R. M. McMeeking, Acta Metall., 34 (1986) 24-35. 27 B. Budiansky, J. C. Amazigo and A. G. Evans, J. Mech. Phys. Solids, 36 (1988) 167-188. 28 M. E Ashby, F. J. Blunt and M. Bannister, Acta Metall., 37 (1989) 1847-1857.
Lffects of ductile phase reinforcement 29 30 31 32 33 34 35 36 37
S. Suresh, Metall. Trans. A, 16 (1985) 249-260. K.S. Chan, Metall. Trans. A, 21 (1990) 2687-2699. W.O. Soboyejo, Metall. Trans. A, 23 ( 1991 ) 1737-1750. R. H. Dauskardt. D. B. Marshall and R. O. Ritchie, J. Am. Ceram. Sot., Z~'(1990) 893-903. R. H. Dauskardt and R. O. Ritchie, Closed Loop, 27(1989) 7-17. R.O. Ritchie, Int. Met. Rev., 20 (1979) 205-230. S. M. L. Sastry and H. A. Lipsitt, Metall. Trans. A, 8 (1977) 299-308. S.-C. Huang and E. L. Hall, Metall. Trans. A, 22 (1991) 427-439. W. O. Soboyejo, P. B. Aswath and J. E. Deffeyes, Mater. Sci. Eng., A138 (1991) 95-101.
95
An investigation of the effects of ductile phase reinforcement on the mechanical behavior of advanced high temperature intermetallics* W. O. Soboyejo Department of Materials Science and Engineering, The Ohio State University, 116 West 19th A venue, Columbus, OH 43210-1179 (USA)
S. M. L. Sastry Department of Mechanical Engineering, Washington University, 1 Brookings Drive, St. LouL~, MO 63130 (USA) (Received January 27, 1993; in revised form April ! 5, 1993)
Abstract The effects of ductile phase reinforcement on the mechanical behavior of Ti-48AI reinforced with 20 vol.% TiNb and MoSi 2 reinforced with 20 vol.% Nb are reported. Ductile phase reinforcement is shown to promote improved fracture toughness and a reduction in fatigue crack growth resistance. The role of crack/tip shielding by bridging and deflection mechanisms is also modeled, and the effects of temperature on the bend and tensile strengths are assessed.
1. Introduction
Recent efforts to develop hypersonic vehicles have stimulated research on advanced high temperature structural intermetallics for potential applications in the temperature range between 700 and 1500°C. In particular, y-based titanium aluminides have been considered for intermediate temperature (700-1000°C) applications [1, 2], while MoSi2-based systems have been proposed for higher temperature uses in the temperature range between 1000 and 1350°C [3-6]. The selection of these systems has been due to their relatively low density (compared with Ni-based superalloys), and strength retention in the stated temperature regimes [1-6]. Unfortunately, however, both the Y titanium aluminide and MoSi 2 based systems are limited by their room-temperature damage tolerance, and there is a clear need for improved fracture toughness and fatigue resistance in both systems. There are essentially two approaches for engineering damage tolerance improvements in y-based titanium aluminides and MoSi2. The first approach relies on intrinsic modification via alloying and variations in thermomechanical processing schedules, and this has been used successfully to engineer toughness *Paper presented at the ASM Symposium on High-Temperature Aluminides and Intermetallics, San Diego, CA, September 16-19, 1991. 0921-5093/93/$6.00
improvements and in y-based titanium aluminides [7-10]. The second approach utilizes brittle or ductile reinforcements to increase fracture toughness via crack-tip shielding mechanisms [11-24]. The latter "composite" approach has been applied, with varying degrees of success, to the extrinsic toughening of ybased titanium aluminides [ 11 - 15] and MoSi 2 [ 15-25]. In particular, ductile phase reinforcement of y alloys with Nb or TiNb has been shown to promote significant (200%-300%) improvements in the work of rupture WR (a measure of toughness), and modest improvements (24%) in the fracture toughness K o due to crack-tip shielding by crack bridging mechanisms [11-13]. Similar improvements in WR and K o have been reported for MoSi 2 reinforced with Nb [24, 25]. Brittle reinforcement of MoSi2 with 20 vol.% SiC whiskers [19] has also been shown to improve the fracture toughness from about 5.3 to 8.2 MPa m ~/2, while particulate reinforcement with brittle TiC has been shown to have little effect on fracture toughness
[2o[. In the current paper, we report the results of recent efforts to engineer improvements in the damage tolerance of a y-based titanium aluminide alloy (Ti-48AI) and MoSi 2 (all compositions are quoted in atomic per cent unless stated otherwise). The effects of ductile phase reinforcement on strength and fracture toughness are rationalized using micromechanics models. Ductile phase reinforcement is also shown to © 1993 - Elsevier Sequoia. All fights reserved
96
W. O. Soboyejo, S. M. I,. Sastry
/
degrade fatigue crack growth resistance, in spite of the toughening that is achieved largely via plastic deformation of the ductile phases and the deflection of cracks around the reinforcements under monotonic loading. The use of linear superposition in the prediction of the combined effects of multiple toughening mechanisms is also elucidated.
2.
Material
The MoSi 2 powder (99.5°/,,, - 3 2 5 mesh) was obtained from Cerac, Milwaukee, WI, and the Nb powder ( - 35 to 80 mesh) was procured from Nuclear Metals, Inc., Concord, MA. The composition of the powders is given in Table 1. The Cerac powder was found to be molybdenum rich relative to stoichiometric M o S i 2 which contains 36.92 wt.% Si. M o S i 2 powders were poured into an Nb can, degassed and sealed by electron beam welding. The can was then hot isostatically pressed at 1700 °C for 4 h under 207 MPa argon pressure. The resulting compact was then cooled slowly (10 h) to room temperature to minimize thermal gradients and cracking. The equiaxed microstructure of the resulting MoSi, compact is shown in Fig. 1. The mean linear intercept grain size was found to be 14/~m, which is consistent with the size of the original powder. This indicates that very limited grain growth occurred during hot isostatic pressing. The density of the material was approximately 97% of the theoretical density, and this is manifested by the incidence of "dark" pores on the photomicrograph shown in Fig. 1. Silicon and SiO, particles (Fig. 1) were also revealed in the monolithic MoSi 2 using energy dispersive X-ray (EDAX) spectroscopy. EDAX analysis revealed that many of the apparent w)ids contained remnants of these particles which were pulled out during metallographic specimen preparation. The Nb-reinforced MoSi2 was produced by phase blending and hot isostatic pressing of nominal 20 voi.% spherical niobium particles (approximately 250/~m in diameter) and MoSi: powder. No SiO2 contamination
Effects o f ductile phase reinforcement
was detected in the Nb by EDAX analysis, and the same processing schedules that were used for the consolidation of monolithic MoSi 2 were employed in the fabrication of the MoSi2/Nb composite. The resulting composite microstructure is shown in Fig. 2. The matrix microstructure (Fig. 2(a)) is similar to that of the monolithic MoSi, (Fig. 1) which had a density of about 97%. However, in contrast, the Nb reinforcements exhibit an almost fully dense structure (Fig. 2(b)) with some incidence of microcracking. A thick (approximately 40 /~m) reaction zone was observed between the Nb reinforcements and the MoSi 2 matrix (Fig. 2(c)). This has been shown to consist of layers of NbsSi 3 and (Mo,Nb)~Si3 [24]. The actual layered interface compositions obtained by semi-quantitative EDAX techniques are shown in Table 2. The Ti-48A! was procured from Nuclear Metals, Inc., Concord, MA, in the form of - 35 mesh (average diameter, 5 0 0 x 10-6m) powder. The powder was produced by the plasma rotating electrode process. The monolithic Ti-48AI powder was canned in T i - 6 A I - 4 V and extruded at a die temperature of 1343 °C with a reduction ratio of 14:1. This resulted in a duplex microstructure with 18 ,um a 2 + 7 lamellar packets and 16 /~m equiaxed }, grains. The Ti-48Al + 20 vol.% TiNb composite was produced by mixing - 35 mesh Nb and Ti-48Al powder, and processing under the same conditions that were employed for monolithic Ti-48AI. This resulted in a duplex matrix microstructure and non-uniform elongated strips of TiNb, as shown in Fig. 3. A small reaction layer was also observed at the interface between the TiNb and the Ti-48Al matrix (Fig. 3). This has been shown by previous workers [14] to consist predominantly of a(2). EDAX compositions of the TiNb
~
. ..
,,
..-
2
~.
o
. . ~, : ~
:r,-
,~.
, .o'
,
)
.~, -.~.
z~r:.t..,,
.;.%
•
---
j4~/¢
o ~. , ¢x.;: .
.
,,.~.~.
..°~.." ~ , ~ . , ~
_
•
.... 1 1
"#"
,~
.
;" . -
TABLE 1. Chemical composition in weight per cent MoSi, Si Fe Nb O C Mo
Ti-48AI 34.99 0.041 < 0.05 0.54 0.O4 Balance
AI Si H N O S "I'i
34.37 0.024 0.001 0.005 0.095 O.OO8 Balance
Fig. 1. Scanning electron photomicrograph showing 97% dense structure of MoSi2 with equiaxed grains and MoSi~Si3 and SiO, particles.
W. O. Soboye]o, S. M. L. Sast~'
,,,¢
/
Effects of ductile phase reinforcement
97
~,L
(b)
lain
Fig. 3. Photomicrographs of y-based titanium aluminides taken under polarized light: (a) monolithic Ti-48AI, and (b) Ti-48AI/ 20 vol.% TiNb and composite.
TABLE 3. Chemical compositions of components of as-extruded Ti-48AI/20vol.%TiNb composite determined by EDAX
L~J lOI.tm
(c)
Fig. 2. Scanning electron photomicrographs showing typical features of the MoSi,/Nb composite: (a) matrix structure. (b) niobium structure, and :el structure of layered Nb-MoSi, interfttce~.
Region
Matrix Interface Reinforcement
Composition (at.%) Ti
AI
Nb
48.5 48.1 57.5
51.5 51.5 15.3
-0.4 27.2
TABLE 2. Chemical compositions fin atomic per cent) of components of MoSi2/2(Ivol.%Nb composite determined by EDAX Region
Nb reinforcement Internal interface Middle interface External interface MoSi, matrix
Composition (at.%) Mo
Si
0.8 1.8 37.2 29.3
0.8 41.6 39.9 70.7
Nb 100.0
98.4 56.6 22.9
particles, interfacial regions, and the T i - 4 8 A 1 matrix are shown in Table 3. Note that the actual interfacial composition observed in this study indicates the presence of a(2) and 7 phases at the interface. It is also important to note here that monolithic/matrix MoSi, were not extruded for direct comparison since the facilities required for extrusion at the required temperatures are not currently available.
98
W. O. Soboyejo, S. M. L. Sastry /
3. Experimental procedures Bend tests were performed on MoSi 2 and the Nb/ MoSi 2 composite at temperatures between 25 and 1500 °C. The tests were conducted on flexure specimens with square (6.35 mm x 6.35 mm) cross-sections, and the specimens were loaded monotonically to failure under three-point bend conditions. The ductileto-brittle transition temperatures (DBTT) for MoSi 2 and Nb/MoSi 2 were thus established from the resulting bend strength data. Similarly, the tensile properties of Ti-48Al and the Ti-48Al/TiNb composite were measured at temperatures between 25 and 982 °C, and DBTT were established for the monolithic and composite y-based alloys. Duplicate tests were conducted on smooth cylindrical button-headed specimens with a gage diameter of approximately 3.18 mm. The specimens were loaded monotonically to failure at a strain rate of 5 x 10-4 S - 1 The fatigue-crack growth behavior of the monolithic alloys and the composites were investigated using 25.4 mm long single edge notched (SEN) bend specimens with square (6.35 mm x 6.35 mm) cross-sections. The tests were conducted in laboratory air under threepoint bend constant amplitude loading after precracking under far-field compression [7, 8] loading. A cyclic frequency of 10 Hz was employed. Crack growth was monitored using a high resolution (2.5 /~m) telescope connected to a video monitoring unit, and an initial load-increasing scheme was employed until crack growth was detected after about 106 cycles. The load range (AP=P(max)-P(min)) was then maintained constant until the end of the test. The tests were stopped prior to specimen fracture. Metallographic examination of the interaction of the crack path with the local microstructure was then carried out, and the micrographs were used to determine the deflection angles. Room temperature fracture toughness tests were also conducted on SEN specimens that were precracked under far-field compression loading [7, 8]. 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 t/2 s - i.
Effects of ductile phase reinforcement
reinforcement-matrix interface and the constrained plastic deformation of the ductile phase during loading to failure. However, in the absence of more detailed analysis, the use of simple rule-of-mixtures is explored in the current investigation. The improvement in toughness due to ductile phase reinforcement can be assessed in to ways, i.e. in terms of G the fracture energy, or in terms of the fracture toughness K o. Improvements in G due to ductile phase reinforcement can be estimated using composite mechanics. This yields [26-28] Gc=(1 -f)Gm +fG r
(2a)
or It*
G~=(1 -f)Gm +f J o(u)du
(2b)
tl
where G m and Gr are the respective fracture energies of the matrix and reinforcement, o(u)is the polynomial expression relating stress to strain in the reinforcement, u is the crack opening displacement (COD) and u* is the COD at the point of fracture. Rearranging eqn. (2b) yields
age = G¢- Gm+ f j o(u)du
(3)
0
Equation (3) represents the toughness increase due to plastic stretching and necking down of the ductile phase reinforcement to a point when u = u*. Further non-dimensional analysis [26-28] reveals that AG¢ is given by AGc =foyrX
(4)
where oy is the yield stress of the ductile phase reinforcements, r is the characteristic reinforcement radius, and X is a dimensionless work of rupture parameter which is given by [28] u*/r
X =f J (O/Oy)d(u/r)
(5)
0
4. Micromechanical modeling The strength oc of a ductile phase reinforced composite system can be estimated using simple rule-ofmixtures. This yields for constant strain conditions
oc=Om(1-f)+ for
(1)
where Om is the matrix strength, f is the reinforcement volume fraction, and or is the reinforcement strength. Note, however, that this simple approach neglects the
where o is the applied stress. Typical values for X are 1.5 for 7-TiAI reinforced with 20 vol.% TiNb and 6.3 for MoSi 2 reinforced with 20 voi.% Nb [ 12, 24]. The fracture toughness can also be expressed in terms of the stress intensity factor at the point of fracture instability, i.e. K o or KIc. For a ductile phase reinforced composite, the toughening components due to potential crack bridging and deflection mechanisms can be assessed using existing analytical models [27, 29]. The toughening due to crack bridging can be assessed using the following expression due to
W. O. Soboyejo, S. M. L. Sastry
/
Budiansky et al. [27]: 2h-
K K,~
- 1 + 2(21~)'i2[foy(LIKm'i2l
m
~
1 cos2(~t/2) .
.
.
.
(6)
.
(7)
The combined effects of crack bridging and deflection can thus be assessed by applying the principle of superposition [30, 31 ]. This yields K = 2b2,1Km
99
TABLE 4. Effects of temperature on the compression strengths of MoSi, and MoSi~ + 20vol.%Nb
where )ch is the toughening ratio due to crack bridging, K is the actual stress intensity factor at the tip of the bridged crack, Km is the matrix stress intensity factor/ toughness, and L is the characteristic length of the reinforcement. Similarly, the toughening due to crack deflection through an angle ~t, 2d, can be estimated using the following expression due to Suresh [29]: K 2,,= K,,
Effects of ductile phase reinforcement
(8)
Equation (8) can be used to estimate the K o or Kjc of a ductile phase reinforced composite in which crack bridging and deflection mechanisms are effective. However, it is important to study the crack paths in the composite system before assuming that crack-tip shielding occurs by the above mechanisms. For the systems discussed in the current paper, the analysis of crack bridging and deflection will be shown to be important in the estimation of the effective driving force for crack growth. However, for other systems, crack-tip blunting by the ductile phase may also contribute to crack-tip shielding [30, 31], and its effects may have to be analyzed when crack growth occurs across the ductile phase.
5. Results and discussion 5.1. M o S i 2 b a s e d systems
The bend strengths of the hot isostatically pressed (HIPed) MoSi2 and MoSi2/Nb composite are compared in Table 4. Both the monolithic MoSi., and the MoSi2/ Nb composite had very high strengths up to 1250 °C, and the strengths of the composite were lower than those of the matrix, as expected from simple rule-ofmixtures. Significant softening was observed in both systems above 1250°C. This is consistent with the results from previous studies on monolithic MoSi, single crystals [32]. The retention of strength in the matrix up to the DBTT has been attributed to a KearWilsdorf type cross-slip mechanism similar to that observed in Ll(2) ordered alloys. The ductility improvement in single crystals of MoSi2 is also associated with increased mobility and climb potential of
Temperature (°C)
MoSi 2 (MPa (ksi))
MoSi 2 + 20vol.%Nb (MPa (ksi))
1050 115(1 1250 1350 1450
594 (86.2) 493 (71.4) 31)5 (44.3) 222 (32.2) 65 (9.4)
513 392 286 156 25
(74.4) (56.9) (41.5) (22.7) (3.6)
(100) and (110) ordinary dislocations in addition to {110}(331) and/or {013}(331) slip which occur both above and below the DBTT. The measured strength of the MoSi2/Nb composite at room temperature (583 MPa) is underpredicted by simple rule-of-mixtures to be 519 MPa. However, it is important to note that the accuracy of simple rule-ofmixture predictions is limited by matrix and powder contamination by interstitiais, porosity, and failure to account for the reinforcement-matrix interactions. More detailed analysis of the constraint associated with the deformation of the Nb reinforcements is also needed for improved strength predictions. Nevertheless, the use of simple rule-of-mixtures provides some useful insights into relative differences between the strengths of the monolithic MoSi, and the MoSiffNb composite. It is of interest to examine the effects of ductile phase reinforcement on the fracture toughness of the MoSiffNb composite. The monolithic MoSi 2 had a fracture toughness of 4.6 MPa m"-" compared with a measured toughness of 5.7 MPa m ~;2 obtained for the composite. This represents a modest increase of about 24% in the fracture toughness due to ductile phase reinforcement. Such a modest increase in fracture toughness appears at first glance to be inconsistent with previous reports of toughness improvement where toughness was characterized by the work of rupture WR [24]. However, although previous studies predicted and measured a work of rupture improvement of about 335% in MoSi2 reinforced with 20 vol.% Nb, they did not consider the fracture toughness in terms of the stress intensity factor at the point of fracture instability Ko •
The increase in WR and Ko in MoSi 2 + 20 vol.% Nb due to ductile phase reinforcement can be predicted using composite mechanics. The increase in WR is given by AG, in eqn. (4). By substituting f = 0 . 2 , oy = 200 MPa, r = 80/~m, and X = 6.3 into eqn. (4), the increase in A G c reported in ref. 24 is predicted to be 202 J m -2. The composite fracture energy G,. is thus computed from eqn. (2) to be 262 J m 2, since the
1O0
W.O. Soboye]o, S. M. L. Sastry /
Effectsof ductile phase reinforcement
Fig. 4. Stable fatigue precrack obtained by far-field compression cyclic loading.
~K O~i./~m. ) 0.5
10 -2
1 i i [ I I
I
I
5 I I I i t
I
10-~
I 10-4
10 -a
10_5 I ~ 10 ~
Fatigue crack growth rata, 10 -5 da/dN 10 ~
da/dN
!
; 10 -7
(in./cycle)
I lO s
(mm/cycle) 10_7
f l0-~
10~ 10 -9
,
0.5
,
,
,
~
,
,
i
,
1 5 Suess intensity range. AK (MP~/-m-)
,
,
~
~
%
. *~1~'-" l i ~
Lif
:ll
)10 ~10 10
Fig. 5. Fatigue crack growth rate data obtained for MoSi2/Nb composite at room temperature (R=Km,JKm~=O.I)(from ref. 25).
matrix fracture energy Gm is approximately 60 J m-2 Similarly, the increase in K o can be predicted for the MoSi2/Nb composite using composite mechanics. The prediction involves the superpositiop of toughening components due to crack bridging (L = 2 r = 160/~m, ;tb=l.19 ) and deflection ( ¢ = 2 0 °, 2d=1.03 ). The respective toughening ratios, 2b and ;tj, are then substituted into eqn. (8) to obtain predictions of K o. The predicted fracture toughness of 5.7 MPa m 1/2 is in exact agreement with the measured fracture toughness of 5.7 MPa m 1/2. The differences between the degree of toughening indicated by changes in WR and K o can therefore be rationalized using existing composite mechanics analyses. Note that the potential effects of SiO2 contamination are not assessed in the above analysis. Attempts to obtain stable fatigue crack growth rate data from precracked SEN specimens were unsuccessful under three-point loading in spite of successful precracking of the monolithic material under far-field compression (Fig. 4), i.e. stable crack growth was not observed in the monolithic MoSi 2 for A K levels with Kmax < Ktc, and the SEN specimens of the monolithic
Fig. 6. Photomicrographs showing the interaction of the fatigue crack with the microstructure of the MoSi2/Nb composite.
MoSi 2 fractured under cyclic loading when Kmax>t K o. However, stable fatigue crack growth was observed in the MoSi2/Nb composite after precracking in compression (Figs. 5 and 6). The occurrence of subcritical fatigue crack growth in the MoSi2/Nb composite is attributed to the inelastic strains that are developed in the relatively brittle layered interfaces (Figs. 2(c) and 2(d)). The intermetallic compounds in these interfacial
W. O. Soboyejo, S. M. L. Sastry
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regions can provide the "'weak link" for initial fracture and damage accumulation under cyclic loading conditions. It is also possible that the matrix MoSi 2 is further embrittled by interstitial elements which are not revealed by the EDAX analysis shown in Table 2. Stable subcritical fatigue crack growth in the MoSi2/ Nb composite occurs at stress intensity factor ranges as low as 1.7 MPa m l.'-~,and the crack growth rates can be expressed as a function of A K, the stress intensity factor range, using a simple Paris power-law da/d N vs. AK relationship. However, as in ceramics [32, 33] the power-law exponent is very high (approximately 14) compared with typical values between 2 and 4 in conventional metallic materials [34]. Limited crack-particle interactions were observed under cyclic loading, and the cracks generally tended to avoid the Nb particles by deflection around the layered interfaces between the Nb particles and the MoSi~ matrix (Fig. 6). Nevertheless, the crack-tip shielding provided by crack deflection [29] was not very significant since the deflection angles under cyclic loading conditions were typically between 45 ° and 60 °. Such deflection angles result in deflection toughening ratios ;t~ between 15% and 25% over an equivalent linear crack (see eqn. (7)). The absence of crack bridging under cyclic loading (Fig. 6)also partly explains the limited resistance of the MoSi~ matrix to fatigue crack growth. Crack-tip shielding contributions from bridging and plastic stretching of Nb particles can be significant for material with 20 vol.% Nb, as discussed earlier. The differences between the resistance of the MoSi2/Nb composite to crack growth under monotonic and cyclic loading can therefore be rationalized by considering (i) the effects of loading and microstructure on the crack path, (ii) the effects of crack path on crack-tip shielding, and (iii) the effects of inelastic strains due to the accumulation of damage at the MoSi~-Nb interfaces during cyclic loading. The current results therefore suggest that improved fracture toughness, K~,., in the model MoSi2/Nb composite does not necessarily imply improved fatigue crack growth resistance. The results also show that improvements in the work of rupture computed by integrating the expression relating reinforcement stress
Effects of ductile phase reinfi>rcement
101
to strain do not necessarily imply similar levels of fracture toughness improvement. Durability and damage tolerance assessment in the MoSi2/Nb composites therefore requires a careful analysis of both fracture toughness and fatigue behavior. It is also encouraging to note that the strength and fracture toughness of the MoSi2/Nb composite can be estimated using existing analyses, and that the MoSi 2 and MoSi2/Nb model composite systems examined here appear to have the potential for 1250-1350°C applications owing to their exceptional elevatedtemperature strength retention. However, further system optimization is required to improve the oxidative stability and damage tolerance capability of MoSi2 and its composites for structural applications. 5.2. Duplex y-based titanium aluminide systems The tensile strengths of the as-extruded Ti-48AI and Ti-48AI-20vol.%TiNb are compared in Table 5. Both the monolithic and composite materials retained their relatively high strengths at temperatures up to 815 °C. Significant degradation in strength and increased plasticity was observed at 982 °C where the composite had a plastic elongation to failure of 84.5°/,,. Reasons for the large increase in ductility are not fully understood at present. However, it is important to note that previous studies have attributed the increased ductility of the monolithic y alloys to the activation of multiple slip systems [35, 36] and transitions in the micromechanisms of fracture that occur at about 650-700 °C 17, 34, 35] and about 900-982 °C [7]. The predicted rule-of-mixture room temperature strength of the Ti-48AI/TiNb composite is 563 MPa (Om=594 MPa [7] and o~=440 MPa [15]). This is within 10% of the measured value of 523 MPa, in spite of failure to account for the reinforcement-matrix interface and constraint effects associated with thc deformation of the TiNb reinforcements. The fracture toughness K~ of the as-extruded Ti-48AI was increased from approximately 15 MPa m ~'-' to approximately 20 MPa m ~'-" by ductile phase reinforcement with 20 vol.% TiNb. This modest increase was predicted by the superposition of the combined effects of crack bridging and deflection
TABLE 5. ('omparison of the tensile properties of as-extruded T i - 4 8 A I and T i - 4 8 A I ~- 20 vol.% TiNb Temperature (°C~
25 700 815 982
Yield stress (MPa)
Ultimate tensile stress (MPa)
Plastic elongation to failure (%1
y
y+TiNb
y
y +TiNb
I,
y -TiNb
594 483 476 262
--394 117
686 579 586 317
623 461 464 132
1.7 2.8 6.(I 20.0
--2.4 84.5
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W. O. Soboyejo, S. M. L. Sastry /
described in Section 4. By substituting in the matrix fracture toughness and the respective toughening ratios due to bridging (2,= 1.3) and deflection (2j= 1.03) into eqn. (8), the fracture toughness of the Ti-48AI/ 20vol.%TiNb composite is predicted to be 20.1 MPa m u2. This is in excellent agreement with the measured fracture toughness of 20 MPa mU2 in spite of failure to account for the strong a2 interface in the analysis. Note that f = 0 . 2 , Oy=440 MPa, L = 1 0 0 0 # m (average length of the TiNb reinforcements which was equated with the bridging length) and 4 = 2 0 ° (based on measurements of crack angles at the sides of specimens). Also, the good agreement between the measured and predicted fracture toughness values confirms the important role of crack bridging reported by previous workers [11-15] for monotonic loading conditions. As observed in the MoSi2/Nb composite, the improvements in the fracture toughness Ktc due to ductile phase reinforcement were modest compared with work of rupture WR or A G c values reported by previous workers [12]. By substituting in f = 0 . 2 , % = 440 MPa, r = 25 /~m (average width of the TiNb reinforcements), and X = 1.5 into eqn. (4), AGc or WR is estimated to be 44 J m- 2 in the Ti-48AI/TiNb composite. This is comparable with the matrix toughness and the toughness of polycrystalline oxides [ 12]. However, the increase of about 100% in AGc or WR is significantly less than the increase of 33% observed in the Ktc of Ti-48AI after ductile phase reinforcement with 20 vol.% TiNb. The use of WR or A G c as toughness parameters in ductility phase reinforced systems must therefore be done with caution since they clearly appear to exaggerate the extent of toughening in cracked structures where Kt~ values provide a better
Effectsof ductile phase reinforcement
characterization of the fracture critical conditions in brittle materials. The fatigue crack growth rate data for the monolithic Ti-48AI are compared with those of the Ti-48AI/TiNb composite in Fig. 7. As for the MoSi2/ Nb composite, the fatigue crack growth resistance of the composite was worse than that of the monolithic matrix alloy. This was due partly to the absence of crack-bridging under cyclic loading conditions, as shown in Fig. 8 which reveals the transgranular fatigue crack path across the matrix and the TiNb reinforcements. Similarly, the contributions to shielding from crack deflection are small owing to the relatively small angles of deflection (see eqn. (7) and Fig. 8). The
Stress intensity factor range. ~,K ( k s i ~ ) 1 10 20 304050 10_4 10-31~-k- t i i , ,Inll ~''
•
u
i0_ s
Fatigue-crack growth rate. da/dN 10-4
Fatigue-crack growth rate. da]dN
(ram/cycle)
10 ~ (in./cycle) 10-5
I
I
I
I l I I l I
I
I
I
I I0-~
10 20 30 4050 Sta'ess intensity factor range. AK (MPaq'm)
Fig. 7. Comparison of fatigue crack growth rate data obtained for monolithic Ti-48AI and the Ti-48AI/TiNb composite, e, Ti-48AI + 20 vol.% TiNb (as-extruded--no TiNb strip across crack path); t~, Ti-48AI-20 vol.% TiNb (as-extruded--with TiNB strips across crack path); , Ti-48AI (as-extruded).
Fig. 8. Photomicrographshowingthe interaction of the fatigue crack path with the microstructure of the Ti-48AI/TiNb composite.
~ O. Soboyejo, S. M. L. Sastry
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fatigue crack growth rates in the T i - 4 8 A I / T i N b composite must therefore represent the "composite" fatigue resistance of the T i N b reinforcements, the ct 2 interface, and the T i - 4 8 A I matrix since the crack growth rates lie between those obtained for T i N b [14] and the T i - 4 8 A I matrix [37]. However, the partitioning of strains and the accumulation in the different sections of the composite cannot be explained so easily since no premature cracking was observed in the T i N b (the c o m p o n e n t with the poorest fatigue resistance) strips ahead of the crack-tip. T h e r e is clearly need for m o r e work in this area. Nevertheless, it is possible to conclude from a practical standpoint that for the two model composite systems examined in this paper, i m p r o v e d toughness due to ductile phase reinforcement does not imply improved resistance to fatigue crack growth.
6. Conclusions T h e effects of ductile phase reinforcement on the mechanical behavior of T i - 4 8 A I and MoSi, were investigated using T i - 4 8 A I / T i N b and MoSi2/Nb model systems. T h e following conclusions were reached. (1) T h e increase in fracture toughness, K o or Kt~, due to ductile phase reinforcement ( 2 4 % - 3 3 % ) is significantly less than the increase in the work of rupture (100°/,,-337%) due to ductile phase reinforcement. Claims of toughness i m p r o v e m e n t in ductile phase reinforced systems may therefore be misleading when toughness is characterized by the work of rupture. (2) T h e i m p r o v e m e n t s in fracture toughness, K o or K~. observed under monotonic loading are associated with a degradation in fatigue crack growth resistance since crack bridging does not occur under cyclic loading conditions. Extrinsic toughening by ductile phase reinh)rcement may therefore degrade damage tolerance in structures subjected to cyclic loading. (3) T h e loss of strength associated with ductile phase reinforced systems can be estimated using simple rule-of-mixtures at temperatures below the DBTT. Also, the D B T T occurs at similar temperatures in the matrix and composite systems, although the extent of softening is greater in the ductile phase reinforced composites.
Acknowledgments T h e authors would like to acknowledge useful discussions with Professor R. O. Ritchie, Dr. K. T. Venkateswara Rao, and Dr. P. J. Meschter. Apprecia-
Effects of ductile phase reinforcement
103
tion is also extended to D. J. Deuser, B. A. Abbott, J. D. Keyes, J. J. Evans, II, J. E. O'Neal and R. J. Lederich for their technical assistance. T h e work on MoSi~ and MoSi, + 20 vol.°/,, Nb was supported by the US Air Force Office of Scientific Research under Contract No. F 4 9 6 2 0 - 9 0 - C - 0 0 3 0 with Dr. Alan H. Rosenstein as Contract Monitor. T h e work on T i - 4 8 A I and T i - 4 8 A I + 20 vol.% T i N b was conducted under the McDonnell Douglas Independent Research and Development Program.
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