70 mm, using the scull melting and casting technique. Table 1 shows the analyzed alloy ..... J. Kumpfert and C.H. Ward, in H. Buhl (ed.), Titanium. Aluminides ...
Materials Science and Engineering
A192/193
(1995) 465-473
Effect of microstructure on fatigue and tensile properties of the gamma TiAl alloy Ti-46.5A1-3.ONb-2.1Cr-0.2W J. Kumpfert”, Y.W. Kimb, D.M. Dimiduk” “DLR, Institute of Materials Research, Cologne, Germany b UES, Materials Research Division, Dayton, OH, USA c Wright-Laboratory, WLIMLLM Wright-Patterson Air Force Base, Dayton, OH, USA
Abstract The relationships between the microstructure and tensile and axial load-controlled fatigue properties of the alloy Ti-46.5Al-3.ONb-2.1Cr-0.2W (atomic per cent) have been studied. Two different microstructures, i.e. duplex (grain size, 20 pm) and fully lamellar (grain size, 300 ,um), were produced, through two-step forging and subsequent heat treatments, giving similar yield strengths at room temperature. The fracture strains at room temperature were about 1.1% and 2.9% for the materials with the fully lamellar and the duplex microstructures respectively. At 600 “C, the duplex material shows a 15% higher fatigue strength than that of the fully lamellar material. At this temperature, the gamma alloy of both microstructures reaches high ratios of the fatigue strength at 10’ cycles to the ultimate tensile strength (UTS), i.e. about 0.95. At 800 “C, the fully lamellar material exhibits higher fatigue strength values above lo5 cycles, and both microstructures result in a two-stage behavior, in contrast to the test at 600 “C. The second stage features the characteristic conventional fatigue behavior, with a broad amplitude stress range, while the first stage is characterized by a narrow band of fatigue stress levels near the UTS. The fracture modes for the duplex material showed a general trend from transgranular to intergranular failure with increasing temperature. For the fully lamellar material, a change from predominantly translamellar failure to a mixture of inter lamellar and translamellar failure was observed, resulting in a microscopically and macroscopically rough fracture surface. The strain rate sensitivity of the fully lamellar material was negligible in the temperature range tested. Keywords: Fatigue; Tensile properties;
Titanium;
Aluminium;
Niobium; Chromium
1. Introduction The need for stronger, lighter high-temperature structural aerospace alloys has led to the development of a group of titanium-based alloys with a high concentration of aluminum. T&Al-based alloys have shown high specific strength values and improved elastic moduli [ 11. However, since this group of alloys will not exceed the maximum service temperature of conventional titanium-based (e.g. Ti-1100, IMI 834) alloys to a sufficient amount to justify the disadvantages, current interest concentrates more on y-TiAl-based alloys [2,3]. Two-phase gamma alloys have received increasing interest since a better understanding of the fundamental and practical aspects of these aluminides led to advanced processing technologies, making them potentially viable engineering alloys [4,5]. Well-defined microstructures now make possible systematic investigations with reproducible results. The majority of alloy 0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved SD1 0921-5093(94)03263-A
and microstructural development studies have focused on ductility enhancement, deformation and fracture behavior, creep behavior, and toughness. Limited information is available on the fatigue behavior of these alloys with different uniform microstructures [6-S]. In the present investigation, the tensile and fatigue properties of a new two-phase gamma alloy, i.e. Ti-46.5A1-3.ONb-2.1Cr-0.2W, was investigated at room temperature and elevated temperatures.
2. Experimental details A two-phase gamma titanium alloy with the composition Ti-46.5A1-3.ONb-2.1Cr-0.2W (atomic per cent) was prepared into an ingot of 20 kg with diameter 70 mm, using the scull melting and casting technique. Table 1 shows the analyzed alloy chemistry, including interstitial elements. The ingot was given a hot isostatic
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Table 1 Alloy chemistry Alloying elements Ti at.% wt.%
Al
Interstitials
Nb
Cr
W
0
C
N
H
bal. 46.5 3.0 bal. 31.5 6.9
2.1 2.7
0.2 1.0
0.570
0.100
0.079
0.045
pressing (HIP) treatment at 1200 “C for 3 h under an argon pressure of 138 MPa to remove any porosity. Billets 120 mm long were cut from the ingot and forged isothermally at 1150 “C in two steps, with an intermediate annealing treatment, to yield pancakes 10 mm thick and 240 mm in diameter. Optical microscopy (OM) revealed non-uniform microstructures at the very edges and the center of the pancakes. Therefore, these regions were eliminated from further examinations. Differential thermal analysis (DTA) as well as metallographic observations were employed to determine the alpha transus temperature (T,). Heat treatment experiments to study microstructure evolution were conducted at various temperatures between 1260 and 1360 “C in the c( + y phase field, and the a phase field. For the annealing treatments, the samples were rapidly heated, by placing them into the furnace at the annealing temperature. The resulting microstructures were investigated using OM, backscattered electron imaging (BSEI) and transmission electron microscopy (TEM). Two microstructures, i.e. duplex (DP) and fully lamellar (FL), were selected for further investigations (Fig. 1). The specimens were annealed at 1280 “C for 3 h and at 1360 “C for 40 min to achieve the DP and the FL microstructures, respectively, followed by furnace cooling to 900 “C and then air cooling to room temperature. Finally, both microstructures were subjected to aging treatments at 900 “C for 6 h. The materials having the two different microstructures were tested for their tensile and fatigue properties. Tensile tests were performed in air at room temperature, 600 “C and 800 “C, at different strain rates between 1 x 10e4 s-’ and 5 x lo-” s-l. The tensile specimens had gauge dimensions of 3.3 mm (diameter) and 16.5 mm (length). Load-controlled fatigue tests were performed at 600 and 800 “C (in air) in the low and high cycle fatigue regimes. Standard, electropolished fatigue specimens with a minimum diameter of 4.7 mm were tested in the tension mode (R = 0.1) at a frequency of 25 Hz and a sinusoidal variation of load with time. The fracture modes of the tensile and fatigue specimens were investigated using a scanning electron microscope (Jeol840). Substructural changes resulting from fatigue deformation were
100 flrn
1
200 pm
1
Fig. 1. Polarized optical images of (a) the DP and (b) the FL microstructures.
studied by OM and TEM. The TEM foils were prepared by sectioning perpendicular to the load axis, grinding and dual jet electropolishing in a solution of 60% methanol, 35% ethylene glycol monobutyl ether and 5% perchloric acid at -40 “C and a potential of 2ov.
3. Results DTA and metallographic observations showed that the alphatransus temperature T, is 1320 f 5 “C. The part of the pancake used for experiments showed a homogeneous and fine-grained microstructure consisting of equiaxed y grains and a2 as a second phase. Slight striations were visible in the flow direction of the material, resulting from forging. Annealing experiments were conducted at various temperatures in the a + y and a phase fields, followed by furnace cooling at roughly 30 “C min- ‘. The volume fraction of the
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equiaxed y grains decreased with increasing annealing temperature. However, in the specimens annealed for 40 min, even at 40 “C, above T,, small equiaxed y grains were observed to be present at lamellar grain boundaries. The duplex (DP microstructure (Fig. l(a)) consists of fine equiaxed y grains and a low volume fraction of lamellar grains with an average size of 20 pm. Solution heat treatment at 1360 “C for 40 min resulted in a lamellar structure with grains 300 pm in diameter (Fig. l(b)). The lamellar spacing in this microstructurehereafter called FL-ranged from 0.5 to 2 pm, as observed by TEM. The lamellar grains in the DP microstructure exhibited lamellar spacings of the same order. The possible existence of a p phase in both microstructures in the non-lamellar a2 grains, resulting from the addition of tungsten [ 121 (as indicated by a third gray level in the BSEI results), was not verified. 3.1. Tensile tests The material with the DP microstructure was tested above and below the brittle-ductile transition (BDT) temperature, using different strain rates, while the FL material was tested below and at the BDT temperature. The results of the tensile tests are summarized in Table 2, and are shown in Figs. 2(a) and 2(b). At room temperature, both microstructures resulted in the same yield strength of about 464 MPa. The fracture strain at room temperature was about 1.1% for the FL material and 2.9% for the DP material. 800 “C is well above the
BDT temperature of the DP material; however, it is approximately equal to the BDT temperature of the FL material, which therefore exhibits higher strength values at 800 “C (Fig. 2(a)). Apparently, 600 “C is below the BDT temperatures for the alloys of both microstructures. At this temperature, the strain rate has very little effect on the tensile properties of the material of either microstructure. However, for the DP material at 800 “C, both the yield strength and the ultimate tensile strength (UTS) increase with increasing strain rate, while the fracture elongation shows a rapid drop from about 41% to 2.6% at strain rates higher than 1 x 10e4 SK’. For the FL material, the effect of the strain rate on the fracture elongation is gradual at 800 “C, with the fracture strain decreasing from 3.5% to 1.2% with increasing strain rate (Fig. 2(b)). The fracture morphologies after the tensile tests were quite different for the two microstructures. At room temperature, the DP specimens showed predominantly transgranular cleavage fracture over the entire fracture surface (Fig. 3(a)). In addition, crack nucleation sites were always located at the surface of the specimens for the DP material at all temperatures tested. At 8OO”C, the fracture morphology of the DP material depends on the applied strain rate. Low strain rates (1 X lo-” s-l) resulted in a ductile failure, with severe necking and the formation of dimples (Fig. 3(b)). However, high strain rates (5 x 10V2 s-l) dramatically reduce the ductility and result in predominantly intergranular fracture surfaces. However, no dynamic re-
Table 2 Tensile test results Microstructure
Temp. (“C)
Strain rate (s-‘1
DP
Room temp.
10-4
600 “C
UTS (MPa)
Elongation (%I
No. of specimens
463
579
2.9
1x
10-4 10-j 5x10-2
405 403 414
550 534 538
3.4 2.9 2.8
10-4 10-j 5 x 10-2
317 374 418
349 438 513
41.5 2.7 2.6
Room temp.
10-4
465
535
1.1
2x
600 “C
10-4 10-j 5x10-2
392 402 407
496 518 500
1.6 1.5 1.4
2x
10-4 10-3 5 x 10-2
370 362 375
497 490 464
3.5 1.7 1.2
3x IX 2x
800 “C
FL
800 “C
“The number of specimens used for each test.
2x
1x 2x 3x
lx 2x
1x 2x
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0 200
400 600 Temperature [Oc]
600
0
1000
200
400 600 Temperature PC]
600
1000
iai 600
- Duplex - 6OooC ??
?? -
T a z 400
-
200
-
8----b-----”
_ Fully-Lamellar ?? UTS
-
Y.S.
-
- 60&
#.-a
UTS
:
8---^--c9
‘s.
-
f c In
“; ;
-
20
_ El. -LLyw 1 o.3 strain rate [s-‘1
. Duplex - 600%
lo-? strain rate [s-l]
.,,..,
-
-0
1 o-5
lo”
: z c
.nn.,
_ Fully-Lamellar
.‘....,
10-l
n.,....,
.rrmu
- 600°C
0 1 o.3 strain rate [s-‘1
1 0.
’
1 o-3 strain rate Is.‘1
10-l
(b) Fig. 2. (a) Tensile properties of Ti-46.5A1-3.ONb-2.1Cr-0.2W with the DP and the FL microstructures vs. temperature at a strain rate of 1 x lo-” s-l. (b) Strain rate sensitivity of Ti-46.5Al-3.ONb-2.1Cr-0.2W with the DP and the FL microstructures at 600 and 800 “C.
crystallization was found for the DP material, even after 39% elongation at 800 “C. The FL specimens show predominantly transgranular cleavage fracture at room temperature, which consists of mainly translamellar failure, leading to relative smooth fracture surfaces (Fig. 4(a)). At 800 “C, however, the fracture surfaces of the FL specimens are macroscopically rough, with transgranular fracture consisting of both interlamellar and translamellar failure. In addition, a considerable amount of plastic deformation is observed in the y plates (Fig. 4(b)). The crack nucleation sites for the tensile tests are not as easy to determine as in the DP microstructure, and could not be clearly identified. 3.2. Fatigue tests The results of the load-controlled fatigue tests are shown in Figs. 5(a) and 5(b) for both microstructures at
600 “C and 800 “C. The open symbols in Figs. 5(a) and 5(b) represent data of specimens which had reached lo7 cycles and were tested afterwards at higher fatigue loads. The first value (first cycle) of each curve shows the UTS measured at a strain rate of 5 X 1O-2 s-l. At 600 “C, the fatigue curves of both microstructures are characterized by nearly flat S-N curves. At this temperature, the DP material showed a fatigue strength (FS) about lo%-15% higher in the range tested. An FL specimen fatigue tested at room temperature showed the same FS level as for the specimens tested at 600 “C (Fig. 5(a)). At 800 “C, both microstructures resulted (at a lower overall stress amplitude level), in a two-step fatigue behavior. The S-N curves in the high stress region are characterized by a gradual linear slope, followed by a steeper portion in the low stress region. The FL material yields an FS 20% higher at lo7 cycles in comparison with the DP material, In the high stress amplitude region, the material of either microstructure
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10 urn
!I
20
urn
10 w Fig. 4. Tensile fracture behavior of the FL microstructure:
Fig. 3. Tensile fracture behavior of the DP microstructure: (a) fracture surface failed at room temperature and 1 x 10m4s- ‘; (b) fracture surface failed at 800 “Cand 1 x lo-' s-l.
(a) fracture surface failed at room temperature and 1 x 10 -’ s _ ’; (b) fracture surface failed at 800 “C and 1 X 10mJ s- ‘.
exhibits comparable FS. Specimens which underwent 10’ cycles, and were tested again at a higher stress level, showed nearly the same FS as those tested at the higher stress alone. The ratio of the FS at lo7 cycles to the UTS (5 X 1o-2 s-1 ), and the ratio of the FS at 10 - ’ cycles to the yield strength (YS; 5 x lop2 s - ‘) are shown for both microstructures and temperatures in Table 3. While the YS values of the material of either microstructure reach about 80% of the UTS values at 600 and 800 “C, the FS/UTS values vary from more than 0.9 at 600°C to 0.46 and 0.66 for the DP and the FL material, respectively, at 800 “C. The typical fatigue fracture morphology for the DP material is shown in Fig. 6. Essentially the same fracture morphology is observed for the DP material under high and low applied stresses at 600 and 800 “C. In all cases, the crack initiation sites were located at the
surface, as was the case in the tensile tests. Transgranular cleavage-type fracture was predominant near the crack initiation sites (Fig. 6(a)), and intergranular fracture in the fast fracture regions. The crack initiation sites show an abundance of “river pattern” and of extensive transgranular “tongues” characteristic cleavage. For the FL specimens, the areas surrounding the crack initiation sites failed by fatigue crack growth at 800°C. At low applied stresses, these areas appear to extend into the interior of the specimen, as can be seen on the fracture surface, as a result of different levels of oxidation (Fig. 7(b)). Under high applied stresses, however, the fatigue crack growth area has a distinctly lower extension (Fig. 7(a)). The fracture mode is predominantly of the transgranular type, with interlamellar and translamellar failure (Fig. 8(a)). Some intergranular failure also occurs along the lamellar
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108
Cycles to Failure 64
Cycles to Failure (b)
Fig. 5. Load-controlled fatigue data for Ti-46SAl3.ONb-2.1Cr-0.2W with the DP and FL microstructures at (a) 600 “C and (b) 800 “C. RT, room temperature.
Table 3 Ratio between FS at lo7 cycles (FS (10’)) and UTS (5 x 10-l SK’), andratiobetweenYS(5x 10-*~-~)andUTS(Sx lo-?s-I) 600 “C
DP FL
800 “C
FS (1O’)/UTS
YSjUTS
FS (1O’)/UTS
YSjUTS
0.95 0.9
0.77 0.81
0.46 0.66
0.x1 0.81
grain boundaries where equiaxed y grains are present with a2 particles. This type of failure is observed predominantly in the high stress regime (Fig. 8(b)). Extensive delamination of the lamellae is observed in the failed fatigue specimens, as can be seen from longitudinal cross-sections (Fig. S(c)).
4. Discussion The results show that the microstructure has a marked effect on the elevated temperature fatigue properties of Ti-46.5A1-3.ONb-2.1Cr-0.2W, even though the YS remains constant. The ahoy exhibits a relatively small prior-a grain size (300 pm) after annealing for 40 min at T,+4O”C (1360°C). Much larger grain sizes (of more than 600 pm) are known to result from the a solution treatment, as reported for the alloys Ti-47AI-lCr-lV-2.6Nb [lO,l l] and
Fig. 6. Fatigue fracture morphology of the DP microstructure at (a) the crack initiation site and (b) the overload area. Tested at 600 “C in the low stress region.
Ti-46.6Al-2.7Nb-0.3Ta [ 121. Apparently, tungsten has a significant effect on the y-+ a transformation and grain growth kinetics. A powder alloy (Ti-48Al-1.9W) is known to show slower gram growth kinetics than for other alloys near y-TiAl[9], and a tungsten-rich b.c.c. p phase was reported to be present at prior-a grain boundaries in a Ti-49Al-2W alloy [ 131. More detailed work is necessary to confirm the tungsten effect in the current alloy, to study the situation with a much lower tungsten level. 4. I. Tensile properties The DP material than that for the FL perature range. The result from the finer order of magnitude)
shows a higher tensile ductility material over the entire test temincreased ductility is believed to grain size (about 20 ,um, or one of the DP microstructure, leading
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to more uniform yielding and to lower plastic incompatibilities at the grain boundaries during plastic deformation [ 10,141. The presence of fine lamellar grains may enhance the ductility, because they can be highly deformable when oriented in the soft mode to the loading direction [ 151. In single-phase microcrystalline y-TiAl, it ha s b een demonstrated that grain refining enhances the uniformity of deformation at room temperature, resulting in increased ductility [ 161. At elevated temperatures, grain boundary sliding can become active, especially in fine-grained material, resulting in a lowered BDT temperature for DP materials (Fig. 2(a)). Recent results attributed the low BDT temperature of a Ti-48Al-2Cr alloy with a cast DP microstructure to dynamic recrystallization of the y grains [ 171. Since dynamic recrystallization is a diffu-
471
Fig. 8. Fatigue fracture morphology of the FL microstructure tested at 800 “C (a) in the low and (b) high stress regimes; (c) longitudinal cross-section below the fracture surface from specimen in (a).
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sion-controlled mechanism, it might be severely restricted under high strain rate conditions. However, in this study, the DP material showed no evidence of dynamic recrystallization, even after 39% elongation at 800 “C. The change from transgranular crack initiation and ductile failure in the overload area at low strain rates to predominantly intergranular fracture at high strain rates indicates that, at high strain rates, the cohesive strength of grain boundaries is relatively lower than the cleavage strain of the matrix. In contrast to other investigations, the FL material in this investigation exhibits YS values comparable with that of the fine-grained DP microstructure below the BDT temperature [ 18,193. Recent results show that the strength of the FL structure increases very rapidly with decreasing grain size [ 10,191. The BDT temperature of the FL material has been determined to be about 8 10 “C [ 10,141. Furthermore, at strain rates of 1 X 1O- 3 s-l or below, the strength levels remain much higher than those of the DP microstructure above 8 10 “C [lo], 4.2. Fatigue properties At 600 “C and high strain rates, the material of either microstructure had FS values greater than 90% of their UTS values, and higher than the YS values for each microstructure. Such results imply that cyclic hardening may have been involved, as was previously reported for a DP material [6]. This is supported by the result that strain hardening in y-TiAl is enhanced by increasing the strain rate [20]. The higher FS values of the DP material are consistent with the higher tensile strengths of this structure, leading to a lower amount of fatigue damage, as a result of a lower plastic strain per cycle. The almost completely flat S-N curves for both microstructures at 600 “C indicate that crack initiation may determine the fatigue lives below the BDT temperature. Once a crack is formed, the propagation through the specimen will be rapid, as a result of the low ductility and the increased stress level. At 800 “C, the ratios of FS to UTS values after 10’ cycles drop to 0.46 and 0.66 for the DP and the FL materials respectively (Table 3). In the low stress regime, or for more than lo5 cycles to failure, the material of either microstructure shows conventional S-N curve behavior; for example, the fatigue live is determined by a broad stress amplitude range. In the low stress fatigue regime, the strains are primarily or totally elastic, and the strength level is expected to play a major role in the plastic behavior. Such expectations are inconsistent with the lower levels of tensile strengths observed for the FL material at 800 “C at a strain rate of 5 x 10e2 s-* (Table 2). The causes for the discrepancy are unknown at present; however, it
appears that crack propagation may become important in the low stress regime. Fracture surfaces of FL specimens tested in the high cycle fatigue regime exhibited an extended fatigue crack growth area around the initiation sites away from the specimen surface (Fig. 7(b)), extending the fatigue life. It is known that the FL microstructure is more resistant to crack propagation under monotonic and cyclic loading conditions [ 10,l 1,191. At high applied stress levels, the stable crack propagation area is considerably smaller and is located near the surface (Fig. 7(a)). The DP material always showed crack initiation at the surface for all stress levels and both temperatures tested. This behavior is different from fatigue tests in vacuum of Ti-6Al-4V (weight per cent), where the crack initiation sites shift from the surface to the interior with increasing stress ratio in the high cycle fatigue regime, independent of the microstructure [2 1,221. It is conceivable that a brittle surface layer may form on the specimens during testing, which overcompensates other effects leading to crack initiation at the surface. At a given test temperature, either 600 or 8OO”C, the fracture morphology of the tensile specimens tested at a high strain rate and the fatigue specimen tested at the same temperature are quite similar but not identical. The DP material showed intergranular crack initiation sites at high strain rates, whereas crack initiation is of a transgranular type in fatigue specimens. This indicates that, during fatigue above and below the BDT temperature, the cohesive strengths of grain boundaries are higher than the cleavage strength of the matrix for the DP material. In the case of the FL material, fatigue results in more abundant delamination in comparison with the tensile tests (Fig. 8(c)). It is well known that slip along the lamellar interface is much easier than is slip across the lamellar interface when the lamellae are oriented in the soft direction with respect to the loading direction [15,23,24]. For polycrystalline material, the grains with such orientations may be called “soft grains”. It appears that the delamination has something to do with the cyclic loading, through promoting easy slip along the lamellar interfaces in soft grains. In other words, the preconditioning may weaken the interface, making the opening up easier in the process zone during fracture. From the current results, it is not clear whether Ti-46SAl3.ONb-2.1Cr-0.2W has an endurance limit, or whether it behaves as aluminum alloys or high-strength steels which do not show such a limit. 5. Conclusions A DP microstructure (grain size, 20 pm) and an FL microstructure (grain size, 300 ym), formed in the
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Materials Science and Engineering,
isothermally forged, two-phase y alloy Ti-46SAl3.ONb-2.1Cr-0.2W, were investigated to correlate the microstructure with the tensile and tension-tension fatigue properties (R = 0.1) at room temperature, 600 “C and 800 “C. (1) At room temperature, the FL material showed the same YS as for the DP material, but the strains to failure were 1.1% and 2.9% for the FL and the DP materials respectively. The BDT temperatures were found to be about 600°C and about 800°C for Ti-46.5A1-3.ONb-2.1Cr-0.2W with the DP and the FL microstructures respectively. (2) Below the BDT temperature, essentially no strain rate sensitivity of tensile properties was observed for the material of either microstructure. At 800 “C, the DP material showed a very pronounced strain rate sensitivity, whereas the FL material at this temperature exhibited very little sensitivity to strain rate variations. (3) In the high cycle fatigue regime at 600 “C, the DP material exhibits a 15% higher fatigue strength than that of the FL material. However, at SOO”C, the FL material reaches fatigue strength values 20% higher for the cycles to failure above lo5 cycles. (4) The tensile fracture modes of the DP material showed a general trend from transgranular to intergranular failure with increasing temperature. For the FL material, a change from predominantly translamellar to a mixture of both interlamellar and translamellar failure and intergranular failure was observed with The DP microstructure increasing temperature. resulted in transgranular crack initiation sites near the surface in fatigue specimens tested at 600 and 800 “C, which is in contrast to the intergranular crack initiation observed in tensile specimens tested in the same temperature range. The FL material showed an extended fatigue crack growth area around the crack initiation site when the stress level was lowered. In this condition, pronounced delamination failure was observed.
Acknowledgments
The current work was performed while one author (JK) was in residence at Wright-Laboratory, WrightPatterson Air Force Base, Dayton, OH. The authors acknowledge the technical assistance from D. Maxwell of UDRF and from J. Henry of UES. One of the authors (YK) acknowledges support from the US Air Force Wright Laboratory, Materials Directorate, under Contract F3361.591-C-5663.
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References [ll J. Kumpfert, C.H. Ward, Y.T. Lee and M. Peters, Proc. Titanium ‘92, San Diego, CA, 1992, The Metallurgical Society of AIME, Warrendale, PA, 1994, p. 1685. [21 J. Kumpfert and C.H. Ward, in H. Buhl (ed.), Titanium Aluminides, Springer, Berlin, 1992, p. 73. C.H. Ward, ht. Mater. Rev., 38 (1993) 79. t:i Y-W. Kim, JOM, 41(1989) 24. Y-W. Kim and D.M. Dimiduk, JOM, 43 (1991) 40. t;i W.E. Dowling, W.T. Donlon and J.E. Allison, Proc. Hightemperature Ordered Intermetallic Alloys Iv Boston, MA, 1990, Materials Research Society, Pittsburgh, PA, 1991, p. 561. H. Jang and [71 D.S. Shih, S.-C. Huang, G.K. Starr, J.C. Chesnutt, Proc. Conf on Microstructure/Property Relationships in Titanium Aluminides and Alloys, Detroit, MI, 1991, The Metallurgical Society of AIME, Warrendale, PA, 1992,~. 135. [Sl S.M. Sastry and H.A. Lipsitt, Metall. Trans. A, 8 (1977) 299. [91 Y-W. Kim, Proc. 3rd Japan Int. SAMPE Symp., Chiba, 1993, Japan Chapter of SAMPE, Japan, 1994, p. 13 10. [lOI Y-W. Kim, Proc. High-Temperature Ordered Intermetallics Alloys Iv, Boston, MA, 1990, Materials Research Society, Pittsburgh, PA, 199 1, p. 777. [Ill S.L. Semiatin, R. Nekkanti, M.K. Alam and P.A. McQuay, Metall. Trans. A, 24 (1993) 1295. [121 J. Beddoes, L. Zhao and W. Wallace, Ser. Metall., 28 (1993) 383. [131 P.L. Martin, H.A. Lipsitt, N.F. Nuhfer and J.C. Williams, Proc. Titanium ‘80, Kyoto, 1980, The Metallurgical Society of AIME, Warrendale, PA, 1980, p. 1245. [141 Y-W. Kim and D.M. Dimiduk, Proc. Conf on Aspects of High Temperature Deformation and Fracture of Crystalline Materials, Nagoya, 1993, Institute of Metals, Japan, 1994. p. 373. [1X H. Inui, M.H. Oh, A. Nakamura and M. Yamaguchi, Acta Metall., 40 (1992) 309.5. 1161R. Imayev, V. Imayev and G. Salishchev, Ser. Metall., 2Y (1993)713. 1994. [171 R. Wagner, Personal communication, [lSl W.E. Dowling, B.D. Worth, J.E. Allison and J.W. Jones, Proc. Conf on Microstructure/Property Relationships in Titanium Aluminides and Alloys, Detroit, MI, 1990, The Metallurgical Society of AIME, Warrendale, PA, 1991, p. 123. [191 Y-W. Kim, Acra Metall., 40 ( 199 1) 112 1. [201 G.T. Gray, III, Proc. Conf on Microstructure/Property Relationships in Titanium Aluminides and Alloys, Detroit, MI, 1991, The Metallurgical Society of AIME, Warrendale, PA, 1991,~. 263. [211 G. Liitjering and A. Gysler, Proc. Ermiidungsverhalten metallischer Werkstoffe, Karlsruhe, 1984, DGM, Oberursel, 1984, p. 39. [221 S. Ada&i, L. Wagner and G. Lutjering, Proc. Titanium ‘84, Munich, 1984, DGM, Oberursel, 1985, p. 2 139. [231 P.M. Hazzledine, B.K. Kad and M.G. Mendiratta, Proc. Materials Research Society Spring Meeting, San Francisco, CA, 1993, Materials Research Society, Pittsburgh, PA, 1993, p. 725. 1241F. Appel, P.A. Beaven and R. Wagner, Acta Metall., 41 (1993) 1721.
