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Effects of microstructure on the deformation and fracture of γ-TiAl alloys
ELSEVIER
Materials
Science and Engineering
Al92/193
(1995) 519-533
Effects of microstructure on the deformation and fracture of y-TiAl alloys Young-Won Kim UES, Materials Research Division, 4401 Dayton-Xenia Rd., Dayton, OH 45432, USA
Abstract Deformation and fracture behavior of two-phase y-TiAl alloys were investigated under monotonic tension loading conditions for duplex and lamellar microstructural forms. The effects of microstructure on tensile properties and deformation-fracture behavior are analyzed for deformation temperatures below and above the brittle-ductile transition. The crack initiation toughness and associated strains near the crack tip are used to explain the inverse relationship between ductility and toughness observed at room temperature. Fracture resistance behavior and toughening mechanisms at room temperature are explained in terms of microstructure and deformation anisotropy. The competition between the effects of grain size and lamellar spacing or tensile and toughness properties is discussed. Keywords:
Deformation;
Fracture; Titanium;
Aluminium;
Alloys
1. Introduction
y-TiAl has attracted a great deal of attention from the aerospace community during the last eight years, and also recently from the automobile industry, because of its potentially attractive properties for high temperature structural applications such as low density, good oxidation and burn resistance, and high temperature strength retention [l-4]. Thanks to the extensive research and development activities in alloy development seen during this period, we have now several two-phase (TiAI+Ti,Al) y-TiAl alloys of engineering importance based on Ti-(4548)Al. Room temperature (RT) mechanical properties of twophase y titanium aluminides are known to depend strongly on microstructure [l-4]. The variations in microstructure that can be controlled in these alloys are numerous, but they exist in four broad groupings; that is near y, duplex (DP), near lamellar (NL), and fully lamellar (FL) microstructures [2,4,5]. DP and FL microstructures are two typical microstructures and have been subjected to the most investigation. A number of mechanical property measurements have been made on these microstructures [5-91, but the studies of fundamental understanding in deformation and fracture behavior have been concentrated mainly on single-phase or DP microstructures [ 10,l l] and on 0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved S.WI 092 I -5093(94)0327 1-8
polysynthetically twinned (PST) crystals [ 12,131. Only for the last two years has attention been drawn to the deformation and fracture behavior of two-phase alloys of engineering importance, particularly in FL conditions [ 14- 171. In general, FL microstructures consist of large lamellar grains resulting in poor tensile properties, and DP microstructures consist of fine grains which yield low toughness and high temperature creep resistance. During this period, it became clear that understanding the behavior and failure processes at RT, as well as high temperatures, is a fundamental basis for optimizing and designing the microstructures for improved and/or balanced properties. Our understanding in this area at RT is quite advanced at least although detailed observations and qualitatively, quantitative analyses remain to be made. Systematic investigations of the elevated temperature deformation and fracture processes have just begun [ 181. In this paper, our recent experimental data on tensile [ 181 and fracture toughness [ 18,191 properties at various temperatures are analyzed to correlate and combine them with previous results [5,14,16,17,20,21] in an effort to understand the whole failure process. Explanations and semiquantitative analyses are made on the dependence of tensile strength and ductility on grain size in lamellar materials, the inverse relationship
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1
Materials Science and Engineering Al 92/l 93 (1995) 519-533
between ductility and toughness, the effect of microstructure on the toughness in FL materials and the competition between grain size (GS) and lamellar spacing on the tensile and toughness properties. The detailed results and analysis for each aspect will be reported elsewhere.
2. Experimental details Ingots of four alloys, with nominal compositions (in atomic per cent) of Ti-47Al-lCr-lV-2SNb (alloy Gl) Ti-47.OA1-1.5Cr-0.5V-2.3Nb (alloy G8), Gi46.5Al-2Cr-3Nb-0.2W (alloy K5) and Ti-47.2Al1.5Cr-0.5Mn-2.6Nb-0.15B (alloy K7), were prepared by skull melting and casting, followed by hot isostatic pressing at 1260 “C for 3 h under an argon pressure of 173 MPa. Billets cut from the ingots were isothermally forged at 1150 “C. Alloy Gl billets were given a 88% height reduction by single-step forging [5] and other alloy billets a 9 1% reduction by two-step forging [ 18,221. Differential thermal analysis combined with metallographic observations were used to determine the a transus T, for each alloy on forged material. Various types of microstructures, DP, NL, FL, refined FL (RFL) and thermomechanically treated lamellar (TMTL), were controlled through specific recrystallization treatments, followed by controlled cooling and then an aging at 900 “C for 6 h. The recrystallization treatment producing a specific DP microstructure for each alloy was conducted at a temperature 40-70 “C below T,. FL microstructures for alloy Gl and G8 having various GSs (300-2600 pm) were produced through recrystallization treatments at temperatures between T, + 1 “C and T, + 35 “C, while RFL structures having a 300 pm GS in alloy K5 were controlled in a temperature range from T, + 15 “C to T, + 3 5 “C. TMTL microstructures having a 150 pm GS were generated in alloy K7 by recrystallization at temperatures between T, + 5 “C and T, + 40 “C. Lamellar spacing for each alloy lamellar microstructure was varied by employing alloy-microstructure-specific cooling methods and rates [18,23,24]. Tensile testing was conducted entirely in air. Tensile tests for data collection were conducted on asmachined specimens with a gauge diameter of 3.4 mm at temperatures between RT and 1000 “C at strain rates in the range from 1 X lo-” s-r to 1 X 10-l s-‘. RT tensile deformation behavior was observed on electropolished flat gauge section specimens. Fracture resistance behavior at RT was investigated for Gl and G8 alloys in vacuum on single-edge-notched (SEN) specimens with ion-milled surfaces in a scanning electron microscope equipped with a loading stage [ 14,161. Displacement rates (DRs) employed for these
tests ranged from 1 x 10-j 4 x 10-l mm s-l. Effective strains near the crack tip were calculated by measuring the displacements with respect to the crack tip. Fracture toughness values were measured in air using mechanically polished compact tension specimens under a loading rate of 1.5 MPa m”* s-l. Microstructures, deformation features and fracture surfaces were observed using light microscopes, the backscattered electron imaging technique, and scanning electron microscopy (SEM).
3. Results 3. I. Alloys and microstructures The chemical analyses of the forged materials for four alloys investigated are listed in Table 1. Also listed are various interstitial levels, with the total interstitial content ranging from 750 to 880 wtppm. The forging conditions-parameters, T,, heat treatment conditions and the resulting microstructures are listed in Table 2. Fig. 1 shows the DP microstructures of these alloys consisting of small (lo-40 pm) y grains with small amounts (2%-5%) of a,-Ti,Al plates-particles dispersed in the matrix. The finest DP microstructure (10-l 5 pm) was obtained in alloy K5 which contains an additional second phase, /?, imagined the brightest in Fig. lc. Fig. 2(a) shows lamellar microstructures consisting of large (above 300 pm) lamellar grains, typical of alloys Gl and G8, with small amounts of fine (5-15 pm) y grains along lamellar grain boundaries (GBs). The lamellar spacing for the FL materials was relatively coarse, ranging from 1.6 to 4.8 pm, because of the slow cooling rates employed for the microstructure. The RFL structure of alloy K5 consists of a 300 pm lamellar grains with fine (5-20 pm) GB 7-P grains (Fig. 2(b)). Th e p resence of /3 phase (Fig. l(c)) during heat treatment is important in controlling GS, which has been recently documented [24]. Because of relatively fast cooling rates that can be applied, the lamellar spacing for the RFL structure of alloy K5 ranged from 0.4 to 0.8 pm. The TMTL structure in alloy K7 consists of about 150 pm size lamellar colonies and ragged boundaries, with lamellar spacing ranging from 0.7 to 1.3 pm. Controlling such fine spacing is possible because the TMT microstructures in boron-containing alloys could be achieved at almost any cooling rates [ 181 and the RFL microstructures can also be obtained under widely ranging cooling rates. The TMT microstructure in non-boron-containing wrought alloys was recently discovered and patented [l&23,25]. The TMT microstructures have been recognized to be much more easily produced beneficially, when small amounts
Y.-W. Kim Table 1 Composition
of y-TiAl alloys developed
/
Materials Science und Engineering A1921193 (1995) 519-S33
521
for the experiments
Alloy
Alloy composition (at.%)
Interstitials (wt.ppm)
Cl G8 K5 K7
Ti-47.OAkl.OCr-0.9V-2.6Nb Ti-47.OA1- 1SCr-OSV-2.3Nb Ti-46.5Al-2.1 Cr-3.ONb-0.2W Ti-47.2AI- 1.SCr-O.SMn-2.hNb-0.15B
78O(O+C+N+H) 75O(O+C+N+H) 805(O+C+N+H)
MO(O+C+N+H)
Table 2 Processing and heat treatment conditions and resulting microstructures for the alloys in Table 1 Alloys
Forging’
Cl
SSF, 88%
1355
DP: T, - 70 to r, - 40 NL: T,-20 to T,- 10 FL: TcJ,+1 to T,+35
20-50 ym GS 30-l 50 pm GS 250-2600 pm GS
G8
TSF, 9 1%
1362
As for Gl
As for Gl
K5
TSF, 91%
1325
DP: T,I-50 NL: TI,- 10 FL: T,+ 15 to To+35
lo-15ymGS lo-120pm GS 200-400 pm RFL
K7
TSF, 9 I %
I365
DP: To-70 TMT: T, + 5 to T, + 40
“SSF. single-step
Recrystallization (“C)
temperatures
Microstructures
lo-15pmGS 80-200 pm GS
forging; TSF, two-step forging.
(0.05-0.3 at.%) of boron are added in wrought alloys, which results in 50-300 pm size colonies depending on the aluminium content, boron amount and heat treatment temperature and time [2,4,18]. Extensive studies on the refining and the controlling mechanism of TMT structures in wrought alloys are under way to explore their engineering significance [4,18].
3.2. Tensile tests RT tensile and fracture properties of all alloys are listed in Table 3 and tensile properties of Gl and K.5 at temperatures between RT and 1000 “C are listed in Table 4. Drastic differences exist in the temperature dependence of the tensile properties for the different microstructures as shown in Fig. 3, where plots are made on alloy K5 data in Table 4 and alloy Gl data from previous data in Ref. [22]. For the DP structures for these alloys, the UTS remains unchanged up to 800 “C while the YS decreases gradually over the same temperature range. Both strengths decrease rapidly at higher temperatures. The alloy Gl FL specimens with large grains exhibit fairly low strengths at RT, but superior strength retention at elevated temperatures, losing not more than 25% of the RT strength at 950 “C. Alloy K5 RFL microstructure also provides high
temperature strength retention, however, with much higher strengths than large-grained FL structures. RT strength for this RFL material is comparable with that of the DP material in spite of a grain size by a factor of 20. The DP-structured material exhibits higher ductility than the FL material throughout the temperature range examined. The brittle-ductile transition (BDT) temperature lies between 650 “C and 700 “C for the DP materials, around 800 “C for the RFL structure, and at 820 “C for the FL material with a 1100 pm GS. Tensile elongation for the large-grained ( 1100 pm) FL material is 0.8% at RT and 2.5% at 800 “C. RT tensile data in Table 3 are plotted in Fig. 4 as a function of microstructure type and GS. Both strength and ductility decreases as GS increases, especially within a class of microstructures. It is noted that the tensile strengths of lamellar materials (FL, TMTL and NL) having 100-300 pm grains are comparable to, or higher than, those of DP materials. NL material having about 100 pm grains shows an extraordinary combination of strength (near 700 MPa) and ductility (greater than 2.0%). The deformation and fracture behavior observed for the DP specimens is shown in Fig. 5. The RT fracture is dominantly transgranular cleavage-like failure, exhibiting river patterns (Fig. 5(a)) in spite of the 2.5% ductility and considerable deformation activity evident in
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Fig. 2. Optical micrographs of the lamellar structures: (a) alloys G 1 and G8 FL, (b) alloy K5 RFL and (c) alloy K7 TMTL microstructures.
I-15pm cig. 1. DP microstructures of (a) alloys G 1 and G8 under optical nicroscopy, (b) alloy K5 under optical microscopy, and (c) alloy 3 in backscattered electron imaging.
the material (Fig. l(a)). Fig. l(a), a micrograph taken from the area below the tensile fracture surfaces tested, shows planar-type deformation features evident in the DP specimens. This trend continues up to the BDT temperature. At 700 “C, the fracture is predominantly intermanular (Fig. 5(b)\ with a few cleavage events.
Y-W. Kim Table 3 Controlled
microstructures
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and lamellar spacing and corresponding
room temperature
523
tensile and toughness
properties
of the alloys
in Table 1 Alloys
Microstructure (pm)
Lamellar spacing (pm)
Alloy Cl G8
DP (20-40) NL(-80) FL (250-2600)
KS
DP(lO-15) NL (10-100) RFL ( = 300)
K7
DP(lO-15) TMT ( = 150)
UTS (MPa)
Elongation @)
Toughness (MPa rn’!‘)
Ref.
EPa)
557-601 702 588-300
2.5-3.8 2.8 1.2-0.5
K,,=K,,,=lO-12
1.6-S
420-485 511 508-280
[51 [51 [5,191
0.4 0.8
46.5 550 475 445
580 685 550 520
2.9 2.1 1.1 1.0
K ,,,=21.5 K max= 20.0
0.7 1.4
435 385 375
555 525 490
2.5 2.8 2.4
K max= K msx=
K,, = 13-21; K ma*= 20-38
K,,=K,,,=lO-11
[ 18,22,24] [ 18,22,24] [ 18,22,24]
18.5 17.0
YS, yield strength; UTS, ultimate tensile strength.
Table 4 Alloy K5 tensile properties at various temperatures and 2 X 10e4 s-’ [ 18,22,24] Microstructure (pm)
Property
Values at the following test temperatures
(“C)
RT
600
700
800
375 480 5.2
319 360 49
DP(lO-15)
YS(MPa) UTS (MPa) Elongation (%)
465 580 2.9
385 528 3.7
RFL ( = 300)
YS(MPa) UTS (MPa) Elongation (%)
473 557 1.2
408 525 1.7
More deformation is evident within grains; however, the planarity of deformation remains unchanged. At higher temperatures, specimens begin to neck with significant elongation. Tremendous amounts of deformation activity with multiple slip and twinning take place at 800 “C (Fig. 5(c)), resulting in an tensile elongation of 50% or more (Table 4 and Fig. 3) and almost entirely intergranular fracture. At still higher temperatures, the fracture surfaces show oxidation debris, and the longitudinal cross-sections reveal grain refinement as a result of dynamic recrystallization taking place, leading to enhanced ductility 17,181. The fracture at temperatures above 900 “C involved the formation of voids. FL specimens exhibit three characteristic fracture features at RT: delamination or interlamellar fracture (Figs. 6(a) and 6(b)), t ranslamellar fracture (Figs. 6(a), 6(b) and 6(c)), and stepwise fracture (Fig. 6(b)), depending on the lamellar orientation with respect to the crack path or stress axis. Translamellar fracture or cracking can occur in the crack arrester orientation
385 510 8.1
900
340 418 19.0
1000
278 284 > 20.00
(lamellar boundaries nearly perpendicular to the crack propagation direction (Fig. 6(c)) or in a more complex manner (Fig. 6(d)). In both cases, cleavage-type features with river patterns are often shown and traces of planar deformation bands parallel to the lamellar interfaces are frequently observed on the fracture surfaces. The planar deformation bands are observed in selective grains on the electropolished specimens during deformation (Fig. 7(a)). Translamellar deformation lines are observed occasionally during deformation, but they appear to form more easily during fracture, within the plastic wake or process zone of the crack (Fig. 7(b)). The elevated temperature fracture of lamellar specimens is characterized by translamellar fracture and delamination or secondary cracking, as observed after failure at 750 “C (Fig. 8(a)) and 800 “C (Fig. 8(b)). The delamination spacing is roughly 15 ,um, and the extent of splitting appears to increase with temperature, often leading to the interlamellar fracture of grains (Fig. 8(b)). Translamellar cleavage fracture was rarely
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600 -
on-
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Materials Science and Engineering
above 800 “C in any alloys. Intergranular observed fracture was observed occasionally in the alloy K5 RFL specimens (Fig. 8(b)), which appears to be related to the presence of GB y grains. In some cases, these GB grains appear to help to block or deviate the propagation of the delamination cracks (Fig. S(a)). At 900 “C, although not shown here, lamellar microstructures become unstable in the necked regions under the strain rates used, and they begin to recrystallize into fine (2-5 ,um) y grams. In this case, ductile fupture involving void formation starts to prevail. Twins were observed on the 900 “C fracture surface, as previously reported [7], indicating that twinning is still an important deformation mechanism at this temperature. However, lamellar microstructures can be unstable even at 800 “C when the tensile elongation becomes more than 8%, as was the case for tests conducted in vacuum at a strain rate of 1 x lo-” s-r [16].
?? mrS
Temperature (“C)
(a)
Al 92/l 93 (I 995) 519-533
3.3. Fracture toughness tests
0
600
600
200
1000
Temperature (“C)
W
Fig. 3. Tensile properties vs. test temperature in alloy Gl and K5 for DP and lamellar (FL and RFL) microstructures. Duplex I NL
l-9 1 IL
I
-
\
01
’ 0
Alloy Gl 0 Alloy K5 A
0 A
. A
Alloy K7
0
??
ility
\y
I 500
I 1000
700
I
I
I
n
I 1500
0
v
I 2000
600
{
2oo
100 2500
Grain Size (urn)
Fig. 4. RT tensile properties of y alloys as a function of microstructure and GS.
RT fracture toughness K,,, values of alloy Gl are lo- 12 MPa rn’iz for DP microstructures and range from 24 to 33 MPa ml’* for FL microstructures depending on GS (250-2000 ,um) (Table 3). The toughness values range from 20 to 22 MPa ml’* for alloy KS RFL material with a 300 pm GS and from 16 to 18 MPa ml’* for alloy K7 TMTL material with a 150 pm GS. The slight increase in toughness in each range appears to be related to the decrease in lamellar spacing. K-resistance curves of alloy Gl in Fig. 9(a) show that initiation toughness K,, and crack resistance behavior differ substantially between the DP and the large-grained FL structures [ 14,16,21]. K,, values for DP material are about 11 MPa ml’* and range from 16 to 21 MPa m1i2 for FL materials, depending on displacement rate. For FL structures, the initiation toughness may vary with lamellar orientation, but has not yet been measured systematically. No resistance to crack growth was observed in DP material (Fig. 9(a)), with K,, remaining unchanged with crack growth. For FL material, substantial resistance to crack propagation was observed, which resulted in saturation or maximum toughness K,,, values ranging from 25 to 42 MPa ml’*. For the tests reported here, on average, two FL grains are fractured by the crack in each specimen in the course of a 2 mm propagation distance (Fig. 9(a)). The variations in resistance were also related to the lamellar orientation with respect to the loading direction or the main crack direction. Some effects of DR on the resistance were observed. The specimens with high K,,, values ( 3 7 MPa m1/2 or higher) generally showed the main cracks growing translamellarly and
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Materials Science and En,+neering A19211 93 (1995) 519-533
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Fig. 5. ‘Tensile deformation and fracture of alloy K5 DP material showing (a) mostly transgranular cleave (SEM) at RT, (b) mostly intergranular fracture (SEM) at 600 “C, (c) deformation structure (optical microscopy) at 800 “C, and (d) dimple fracture (SEM) at 800 “C.
bridging ligaments failing interlamellarly by shear, as shown in Fig. 10. The displacements near the crack tip at were measured on the surface K=K,,=16MPam1~* of an alloy Gl FL SEN specimen having large lamellar grains (above 1000 pm) with laths inclined at 8 = 48” to the notch direction ( 8 = 0’) which is perpendicular to the loading direction [14]. The isostrain contours constructed using the effective strains calculated from the measured displacements show a large plastic zone
with a butterfly shape which is roughly symmetric with respect to the lamellar direction [26]. When the data of effective strains are plotted against radial distance r from the crack tip, a linear relationship exists on a double-logarithmic scale in a given direction, as shown in Fig. 9(b) for three directions, 8= O”, 45” and 90”. Two important aspects are obvious from Fig. 9(b). The effective strain is orientation dependent, and the maximum strain, or fracture strain, ranges from 10% to 2.5% depending on direction, with the largest straining
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Materials Science and Engineering AI921193 (1995) 519-533
Fig. 6. RT tensile fracture surfaces (SEM) of lamellar materials showing: (a) mostly translamellar fracture with a few interlamellar fracture events in alloy K5, (b) combination of interlamellar and stepwise fractures in alloy K5, (c) translamellar cleavage with river patterns and lamellar slip bands in alloy G8, and (d) complex translamellar cleavage-type fracture in alloy G8.
taking place in the soft orientation (0 = 0’). This indicates that under a triaxial loading condition the lamellar structure can be deformed to a large extent in any direction. This is in contrast for the DP structure which shows a small maximum effective strain of about 3% in front of the crack tip that was obtained from similar displacement measurements on a Gl DP specimen (Fig. 9(b)).
4. Discussion 4.1. Tensile deformation andfiacture The ductility of y alloys decreases, in general, with increasing GS, and so does tensile strength (Fig. 4). The ductility for lamellar structures decreases from 2.5% to 0.5% as the GS increases from 250 pm to 2500 ,um [5]
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Materials Science and Engineering A 192/I 93 (I 995) 519-533
3mm
Fig. 7. Alloy G8 FL flat specimen fractured at RT showing on the electropolished surface (a) deformation of some specific grains only and (b) plastic deformation structure in the crackwake or process zone.
4). The Hall-Petch relationship between YS and GS in FL microstructures exists with a large Hall-Petch constant Ic,,= 5 MPa m l/* for grain sizes from 250 to 2600 pm and yield strengths from 290 to 500 MPa [22], as shown in Fig. 1 l(a). This value is compared with those ( 1 .O-1.2 MPa ml’?) for fine y-DP microstructures [27-291 and that (0.5 MPa m’12) for PST crystals in a hard orientation [ 131 (Fig. 1 l(b)). This unusually large Hall-Petch constant is not readily explained, but may be related to the strong anisotropic flow stress behavior [3,12,13] of the lamellar structure. The gauge diameter to GS ratio may also play a role in the unusual relation. The ratio ranges from 1.3 to 14 for the FL specimens. Therefore, the specimens containing large grains have only a few grains within the gauge cross-section, with the result that one or two soft grains control the yielding process. An investigation is underway to clarify this uncertainty [ 181. (Fig.
527
15pm . ..-
Fig. 8. Elevated temperature tensile fracture (SEM) of lamellar structures showing (a) translamellar fracture with delamination or lamellar splitting and intergranular fracture of GB y grains at 750 “C in alloy Gl and (b) translamellar, interlamellar and intergranular fractures in alloy K5 at 800 “C.
The inverse relationship between ductility and GS (Fig. 4) was recently analyzed by assuming the presence of a microcrack comparable with the average GS and subsequently estimating the stress intensity of the crack and the resultant ductility due to crack propagation [ 141. However, the source of the assumed microcrack was not identified clearly and the role of general plasticity was not defined. On the contrary, the decreases in tensile strength and ductility with increasing GS in lamellar material can be explained using the characteristic anisotropy of tensile properties of lamellar structures [ 12,131, and assuming that the fracture is controlled by the crack nucleation process [30] involving the coalescence of
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FiTlAir 0
I 0
(a)
0.5
I
I
I
1.0
1.5
2.0
Aa
LO
03
r (mm)
Fig. 9. Fracture behavior of alloy Gl at RT in DP and FL conditions: (a) K-resistance curves and (b) effective strain field near crack tip measured at K,, = 16 MPa ml/* in the three directions with respect to the notch direction (8= 0”) which is perpendicular to the loading direction (LD). The effective local fracture strain of 15%-20% for the lamellar microstructure is compared with that of 3% for the DP structure.
dislocations under shear stress [ 3 11. In this process, the local magnified shear stress r,, at the pile-up is known to be a function of grain size d and applied stress r, and can be expressed as
Fig. 10. Compact tension specimen surface (electropolished) during RT testing showing (a) the major crack, a microcrack in the adjacent grain, a ligament area, and both interlamellar and translamellar slip-twin bands (optical microscopy) and (b) a magnified view (SEM) of the ligament area showing fine microcracks formed in the translamellar deformation direction [19,22].
tp = C, dra2 or, conveniently, aP = Cdoa2
(1)
where C, and C are constants. Using the relation (1) and the flow curve equation cs= KE”, the local flow curves for two different grain sizes d, and d2 with d, < d2 are constructed as shown in Fig. 12(a). Both materials fracture at the same local fracture stress oPr by initiating cracks but at different fracture strains elf and cZf, with elf > eZf. Using these fracture strains, the flow curve equation, and the relation o, = C,d- u2 (for given ur) derived from relation (l), the tensile stress-strain curves for the two materials having different GSs can be constructed as shown in Fig. 12(b). It is
then clear from Fig. 12(b) that the material having smaller grains shows not only higher flow stress but also greater fracture strain, resulting in its fracture stress calf being higher than that oa2f of large-grained material. Ductility is then primarily controlled by the amount of general yielding which will be greater for small grain sizes. In the FL material, cracking and subsequent fracture precede general yielding resulting in nonuniform deformation and low strains to failure [20,22,32]. Slip bands observed within the grains having easy orientations (Fig. 7(a)) are consistent with
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Materiuls Science and Engineering A 192/I93 (I 995) 519-533
d W)
3000 300 100 30 \\//I0 I 8001
- 800
-
EL
-
(Kim931
w - 600
- 400
- 200
(“makosh, 92)
0
400
800
1200
1600
(a) d-'" V"')
6oor-----l
(b)
Tensile Strain (%)
Fig. 11. (a) Hall-Petch relation obtained from the data of alloy Gl-G8-K5 FL materials is compared with those reported for other microstructural forms and (b) stress-strain curves of a TiAl PST crystal in various orientations which were drawn by rearranging the data of Ref. [ 121.
Fig. 12. Estimated flow curves of y alloys for two different grain sizes (d, < dZ) in two different stress conditions: (a) local stress-strain curves near the dislocation pile-up generated by the stress concentration ur and (b) the materials’ general tensile stress-strain curves under applied stress a,. For details, see text.
529
the deformation bands parallel to the lamellar directions observed on fracture surfaces in Figs. 6(c) and 6(d). This planar deformation leads to dislocation pileups nucleating microcracks, often in an adjacent grain (Fig. 13), which is usually preceded by either translamellar slip-twining or additional interlamellar deformation, and may result from plastic incompatibility stresses between grains. The microcrack usually forms in the translamellar direction (Fig. 13), although interlamellar microcracking should be possible depending on the lamellar orientation. While this plausible mechanism of plasticity-induced crack nucleation is now evident, the event has not been quantitatively tied to a stress or strain. Consequently, it is not yet entirely clear whether uniaxial tensile ductility and failure are a crack nucleation controlled event [31], or if propagation to a critical size is required. Tensile ductility is characterized by the BDT temperature which varies from 650 “C to 820 “C (Table 4 and Fig. 3) depending on alloy composition and microstructure. For a given alloy, the BDT temperature increases with GS or with decreasing RT ductility. It is also strain rate sensitive, especially for fine microstructures, as was observed in a DP microstructure [ 18,331. The enhanced ductility above the BDT temperature appears to be the direct result of increases in both twinning-slip activity and systems, as was discussed earlier. Some claims that dynamic recrystallization controls BDT temperature have yet to be confirmed. For DP microstructures, tensile strengths at temperatures below the BDT temperature depend on cooling rate and method [5,18,23,24] employed, with higher cooling rates yielding higher strength levels. The (72phase, whose morphology and distribution strongly control mechanical properties, is present in fine particle (instead of plate) forms with fast cooling rates. Finer lamellar spacings result in NL and FL microstructures with increasing cooling rate. Decreases in lamellar spacing appear to result in some increases in tensile strengths and ductility (Table 3), but this relation has not been quantitatively characterized. Much work is needed to quantify the relationship between cooling rate-scheme and mechanical properties. Cleavage and transgranular (translamellar and interlamellar) cracking are the dominant fracture modes of RT for both microstructures, although the fine DP material shows a mixture of transgranular and intergranular failure [ 17,20,22]. Cleavage-like brittle fracture was also observed in PST crystals even for failure occurring parallel to the lamellar interfaces with more than 20% plastic strains [ 12,341. At temperatures above BDT temperatures, intergranular fracture prevails in DP material and delamination and intergranular fracture become important for FL material [17,22,33]. In spite of the increased deformation
Y.-W. Kim
Fig. 13. Nucleation
I L
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Materials Science and Engineering Al 92/l 93 (1995) 519-533
of microcracks
I
I
at dislocation
pile-ups against grain boundaries
Spacing (pm)
I
Fully Lamellar
!
in FL materials.
50-
.?.
40.
(Expexted) ___________-----
,
4
1
,
0.44
,
0.25
1.5
2
AllDv
: 30
20
10 (150 W) 0 -0
...~‘~~~.l~~~.‘.~.~0.5 Lamellar
I
:
0
I
I
I
I
500
1000
1500
2000
IO
2500
1 Spacing,
XjlR
(pm-‘R)
Interrelations between fracture toughness, lamellar spacing d, in polycrystalline lamellar y alloys.
GS, and
Grain Size (pm)
Fig. 14. Inverse ductility-toughness K,, and K,,, relationships observed in alloys Gl and G8 but characteristic of y-TiAl alloys. Some data of alloys KS and K7 are plotted for comparison.
1
Duplex
1
Duplex
1
:&:, ;:,(-:‘, strain (a) Tensile Curves
Strain (b) Crack-Tip
Strain Local Yielding
Fig. 15. Comparison of general tensile yielding and crack tip local yielding behavior: (a) measured tensile stress-strain curves typical of DP and FL microstructures and (b) estimated local stress-strain flow behavior of the material in the near-tip plastic zone for these microstructures.
activity with temperature, transmission of slip and twinning across GBs or lamellar interfaces appear to remain difficult [22]. The apparent GB interface weakening with increasing temperature has yet to be elucidated.
4.2. Fracture toughness and toughening mechanism The K-resistance curves can be divided into two components; that is, the resistance to crack initiation or initiation toughness K,, and resistance to crack growth. K,, is essentially a measure of the plastic deformation before the crack propagation, from which the local tensile properties can be estimated within the plastic zone, as discussed in the next section. DP material exhibits small plastic strain near the onset of
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Materials Science and Engineering Al 92/l 93 (I 995) 519-533
crack extension and no resistance to crack propagation, whereas the FL structure yields large crack tip plastic strains and increased resistance to crack propagation (Fig. 9(a)) [ 14- 161. At RT, FL specimens show K,, = 16-21 MPa rn’i2 and K,, = 25-42 MPa rn’i2 and the DP specimens have KIc = K,,,,, = 11 MPa mii2 (Fig. 7(a)). The initiation toughness K,, is related to the strain measured near the crack tip (Fig. 9(b)). As shown, the fracture strain (at 0 = 0”) in the plastic zone can be 20% for FL material but about 3% for DP material. The occurrence of crack growth resistance is caused by the work expended mainly in failing the ligaments formed in the wake of crack [ 14,211. Extensive deformation activity, both interlamellar and translamellar, is seen in the ligament region in Fig. 10. This particular ligament is oriented for interlamellar shear failure linking the main crack and the microcrack which are both translamellar type. This crack orientation results in higher crack resistance than the other type of orientation, that is interlamellar main crack path with translamellar ligament failure [22]. Ultrafine microcracks are formed on the deformation planes during the translamellar deformation (Fig. 10(b)), and crack propagation can be shielded by these cracks and some other off-path cracks shown on the GB. Another, but less effective, source of growth resistance can be the increased deformation amounts due to the increased plastic zone size with crack growth. The sensitivity of toughness to displacement or loading rate increases with temperature; for example, K,, and K,,, values for the large-grained FL specimens at 800 “C were measured to be 35 MPa m1j2 and 79 MPa ml” respectively [16]. The absence of crack growth resistance behavior in DP structures is observed at least up to 600 “C and is due to the lack of bridging ligaments.
4.3. Ductility-toughness
relationship
The inverse relation between ductility and toughness (both K,, and K,,,), summarized in Fig. 14, can be explained by analyzing the bulk tensile behavior and the crack tip yielding behavior [22,32]. Tensile ductilities of polycrystalline material are primarily controlled by the amount of general yielding which will be greater at small grain sizes. In toughness testing of largegrained (above 500 pm) FL specimens, however, the crack tip plastic zone (estimated to be about 400 pm) encompasses only one or part of a grain in front of the crack tip, thereby confirming the deformation behavior to the characteristics of those few grains along the crack tip edge for thin specimens. K,c is then closely related to the intrinsic ductility of the lamellar structure and can be used to estimate the tensile stress-strain
531
curve within the plastic zone (22,321. The near-tip effective fracture strain in the lamellar structure in Fig. 9(b), about 20% at 0 = O”, is consistent with the tensile fracture strain directly measured on a TiAl PST crystal (Fig. 1 l(b)) [12,13]. In DP materials large plastic volumes are not sampled because of the small grain sizes. Fig. 15 compares the measured general tensile curves (Fig. 15(a)) and the stress-strain curves within the near-tip plastic zone (Fig. 15(b)) for coarse (larger than plastic zone size) FL and fine (lo-40 pm) DP materials [22,32]. The stress-strain curves in the plastic zone were estimated using the following relation [35] between K,, and local stress-strain behavior: Klc2 = CnEo
Y
EIf
(2)
where C is a constant, IZ the local work hardening coefficient, E the elastic modulus, oy the material’s YS, and qf the local true fracture strain which corresponds to the maximum effective strain measured near the crack tip (Fig. 9(b)). The K,, values used for the estimation were 10.5 MPa m’i2 and 16 MPa m’i2 respectively for the DP and the FL materials (see Fig. 9(a)). For FL materials with grain sizes greater than 500 pm, the plastic zone sizes rP at the onset of crack initiation range from 400 to 500 ym as calculated using the relation rP= a( K/o,)~, and it is thus contained within a grain in front of the crack tip. In this case, the local fracture strain, about 15%, was taken as an averaged value from Fig. 9(b). For the DP material, the material within the plastic zone (which is estimated to be 300 pm in size and contains many randomly oriented grains) is fairly isotropic and represents the bulk material with a fracture strain of about 3%, and thus the near-tip stress-strain curve is essentially identical to the tensile curve (Fig. 15(b)). The results in Fig. 15(b) clearly show that the socalled inverse ductility-toughness relationship is caused simply by the difference in sampling between tensile testing and toughness testing. Within the plastic zone, the material obeys the classical relation between tensile properties and initiation toughness expressed by relation (2). 4.4. Competition between grain size and lamellar spacing in toughness As was described in the previous section and shown in Fig. 14, the fracture toughness K,c and K,,, increases with lamellar GS up to about 500 pm, but appears independent of GS for larger sizes. In this constant-toughness region against GS, however, it was found that the fracture resistance increases with decreasing lamellar spacing II, as was measured in
532
Y.- W. Kim
1
Materials Science and Engineering
alloys Gl and G8 specimens in a il, range from 4.8 to 1.6 pm [19]. Fig. 16, where the data in Ref. [19] are selectively plotted, shows the Hall-Petch relation between toughness (especially K,,) and lamellar spacing, as follows: K Ic =K 0 +k;l L -r/l
(3)
where K, may be defined as the intrinsic K,, value and k is a Hall-Petch constant. The increase has been explained in terms of increased stresses required for translamellar slip-twinning [19]. The flow stresses of the lamellar structure in the hard orientation are higher than those of polycrystalline lamellar materials and increase to much higher levels with decreasing lamellar spacing (Fig. 1 l(a)). Interestingly, K,, of the FL material with large lamellar spacing becomes the same as that of DP material, that is about 10 MPa ml/*, when the lamellar spacing reaches about 10 pm (Fig. 16). This suggests that the lamellar structure with coarse lamellar spacing may behave like a fine-grained material by losing the anisotropy in deformation. In this case, the lamellar interfaces and domain boundaries are expected to act as GBs effectively pinning the dislocations. When the GS is smaller than the plastic zone size, then the GB effect on the toughness will begin to prevail over the lamellar spacing effect. Although not totally quantitative, this prediction is evidenced from the test results (Table 3) for K5 and K7 alloy compact specimens containing respectively 300 pm and 150 pm lamellar grains (Fig. 16). In this GS range, the variation in lamellar spacing between 0.4 and 1.3 pm does not alter the toughness K,,, appreciably, indicating that the Stroh fracture stress within the plastic zone only slightly depends on the variation in lamellar spacing. The large drop in toughness in this GS range, compared with that in the large-grain range, is thus mainly the GS effect. Systematic studies are underway to quantify the competition between GS and lamellar spacing using large, thick specimens [ 181, and the results will be reported elsewhere.
5. Conclusions RT and elevated temperature deformation and fracture behavior of y-TiAl alloys was investigated under tensile and compact-tension loading conditions. Tensile properties strongly depend on microstructure and temperature. FL structures yield lower ductility and strength, but greater strength retention at high temperatures, and higher BDT temperatures than DP structures. A Hall-Petch relationship with an unusually high Hall-Petch constant was observed to exist between YS and GS for FL structures; however, the
Al 92/l 93 (I 995) 519-533
mechanism for this phenomenon is unclear. The dependence of ductility on GS is analyzed to be limited by the formation of strain incompatibility induced microcracks. For both microstructures, fracture is preceded by plastic deformation followed by cleavagelike failure. Fracture toughness is also a strong function of microstructure and temperature, with higher toughness achieved in FL structures and also at higher temThe high fracture toughness for peratures. large-grained FL material is explained in terms of the characteristic of the lamellar structure observed at the crack tip, that is large plasticity and the ability to form ligaments. For a given large-grained lamellar material, fracture toughness is a function of lamellar spacing with Hall-Petch relation. The inverse the ductility-toughness relationship is explained by comparing the general tensile yielding and the near-tip local plasticity measured at the onset of crack initiation. DP materials fail transgranularly at RT and intergranularly at high temperatures. FL materials show predominantly transgranular fracture over a range of temperatures as long as the microstructures are stable. Decreasing GS or GB y phase, however, promotes intergranular fracture at elevated temperatures.
Acknowledgments
The author appreciates the technical discussions with Drs. D.M. Dimiduk and KS. Chan and the technical assistance from Ms. Sonya Boone. Most of the work in the paper was conducted under Contract F33615-91-C-5663, US Air Force Wright Laboratory, Materials Directorate.
References Y.-W. Kim, J. Met., 41(7) (1989) 24-30. Y.-W. Kimand D.M.Dimiduk,J. Met., 43(8)( 1991)40-47. M. Yamaguchi and H. Inui, in Structural Intermetalks, TMS, Pittsburgh, PA, 1993, pp. 127-142. Y.-W.Kim,J. Met., 46(7)(1994)30-39,7. Y.-W.Kim,Acta Metall. Mater., 40(6)(1992) 1121-1134. D.S. Shih, S.-C. Huang, G.K. Starr, H. Jang and J.C. Chesnutt, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property
Relationships
in Titanium
Aluminides
and
A/loys,TMS, Pittsburgh, PA, 1991, pp. 135-148. S. Krishnamurthy and Y.-W. Kim, in Y.-W. Kim and R.R. Relationships in Boyer (eds.), Microstructure/Property Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991,~~. 149-163. W.E. Dowling, B.D. Worth, J.E. Allison and J.W. Jones, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991, pp. 123-134. A.W. James and P. Bowen, Mater. Sci. Erg. A, 153 (1992) 486-492.
Y.-W. Kim [lo] [ 1 I] [12] [ 131
[ 141 [15] [16]
[ 171 [ 181 [ 191 [20] [21] [22]
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H.A. Lipsitt, D. Schechtman and R. Schafrik, Metafl. Trans. A, 6(1975) 1991-1997. S.-C. Huang and E.L. Hall, Metall. Trans. A, 22 (1991) 427-438. H. Inui, M.H. Oh and M. Yamaguchi, Acta Metall. Mater., 40(1992)3095-3104. Y. Umakoshi, T. Nakano and T. Yamane, Mater. Sci. Eng. A, I52(1992)81-88. K. Chan and Y.-W. Kim, Metall. Trans. A, 23 (1992) 1663-1677. H.E. Deve, A.G. Evans and D.S. Shih, Acta Metall. Mater., 40(1992) 1259-1265. K. Chan and Y.-W. Kim, Metall. Trans. A, 24 (1993) 113-125. Y.-W. Kim and D.M. Dimiduk, Mater. Res. Sot. Symp. Proc., 288( 1993) 671-677. Y.-W. Kim, research in progress, 1992-l 994. K. Chan and Y.-W. Kim, Acta Metall, Mater., in press. G. Malakondaiah, Y.-W. Kim and T. Nicholas, Ser. Metall., 3U( 1994) 939-944. K. Chan and Y.-W. Kim, Metall. Trans. A, 25 (1994) 1217-1228. Y.-W. Kim and D.M. Dimiduk, Proc. JlMIS-7 on High Temperature Deformation and Fracture, Japan Institute for Metals, Sendai, 1993, pp. 373-382.
533
[23] Y.-W. Kim and D.M. Dimiduk. US Patent 5,226,985, July 13, 1993. [24] Y.-W. Kim and D.M. Dimiduk, filed for a US patent, 22 April 1994. [25] P. McQuay, D.M. Dimiduk and Y.-W. Kim, filed for a US patent, May 1993. (261 Y.-W. Kim and K.S. Chan, unpublished results, 1994. [27] S.-C. Huang and D.S. Shih, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991, pp. 105-122. (281 P. McQuay, M.S. Thesis, Wright State University, OH, 1992. [29] C. Koeppe, A. Bartels, J. Seeger and H. Mecking, Metal/. Trans. A, 24(1993) 1795-1806. [30] A.N. Stroh, Philos. Mag., 46 (1955) 968-972. [31] A.N. Stroh, Proc. R. Sot. London Ser. A, 233 (1954) 404. [32] Y.-W. Kim, in Intermetallic Compounds, Proc. 3rd Jpn. Int. SAMPE Symp., Chiba, 1993, pp. 3 lo- 13 17. [33] i;$;,pfert, Y.-W. Kim and D.M. Dimiduk, Mater. Sci. Eng. [34] M.H. Oh, H. Inui and M. Yamaguchi, Acta Metall. Mater., #I (7) (1993) 1939-1949. [35] J. Langford, D.L. Davidson, K.S. Chan and G.R. Leverant, AFOSR Rep., Southwest Research Institute, May 1989.
Materials
Science and Engineering
Al92/193
(1995) 519-533
Effects of microstructure on the deformation and fracture of y-TiAl alloys Young-Won Kim UES, Materials Research Division, 4401 Dayton-Xenia Rd., Dayton, OH 45432, USA
Abstract Deformation and fracture behavior of two-phase y-TiAl alloys were investigated under monotonic tension loading conditions for duplex and lamellar microstructural forms. The effects of microstructure on tensile properties and deformation-fracture behavior are analyzed for deformation temperatures below and above the brittle-ductile transition. The crack initiation toughness and associated strains near the crack tip are used to explain the inverse relationship between ductility and toughness observed at room temperature. Fracture resistance behavior and toughening mechanisms at room temperature are explained in terms of microstructure and deformation anisotropy. The competition between the effects of grain size and lamellar spacing or tensile and toughness properties is discussed. Keywords:
Deformation;
Fracture; Titanium;
Aluminium;
Alloys
1. Introduction
y-TiAl has attracted a great deal of attention from the aerospace community during the last eight years, and also recently from the automobile industry, because of its potentially attractive properties for high temperature structural applications such as low density, good oxidation and burn resistance, and high temperature strength retention [l-4]. Thanks to the extensive research and development activities in alloy development seen during this period, we have now several two-phase (TiAI+Ti,Al) y-TiAl alloys of engineering importance based on Ti-(4548)Al. Room temperature (RT) mechanical properties of twophase y titanium aluminides are known to depend strongly on microstructure [l-4]. The variations in microstructure that can be controlled in these alloys are numerous, but they exist in four broad groupings; that is near y, duplex (DP), near lamellar (NL), and fully lamellar (FL) microstructures [2,4,5]. DP and FL microstructures are two typical microstructures and have been subjected to the most investigation. A number of mechanical property measurements have been made on these microstructures [5-91, but the studies of fundamental understanding in deformation and fracture behavior have been concentrated mainly on single-phase or DP microstructures [ 10,l l] and on 0921-5093/95/$9.50 0 1995 - Elsevier Science S.A. All rights reserved S.WI 092 I -5093(94)0327 1-8
polysynthetically twinned (PST) crystals [ 12,131. Only for the last two years has attention been drawn to the deformation and fracture behavior of two-phase alloys of engineering importance, particularly in FL conditions [ 14- 171. In general, FL microstructures consist of large lamellar grains resulting in poor tensile properties, and DP microstructures consist of fine grains which yield low toughness and high temperature creep resistance. During this period, it became clear that understanding the behavior and failure processes at RT, as well as high temperatures, is a fundamental basis for optimizing and designing the microstructures for improved and/or balanced properties. Our understanding in this area at RT is quite advanced at least although detailed observations and qualitatively, quantitative analyses remain to be made. Systematic investigations of the elevated temperature deformation and fracture processes have just begun [ 181. In this paper, our recent experimental data on tensile [ 181 and fracture toughness [ 18,191 properties at various temperatures are analyzed to correlate and combine them with previous results [5,14,16,17,20,21] in an effort to understand the whole failure process. Explanations and semiquantitative analyses are made on the dependence of tensile strength and ductility on grain size in lamellar materials, the inverse relationship
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between ductility and toughness, the effect of microstructure on the toughness in FL materials and the competition between grain size (GS) and lamellar spacing on the tensile and toughness properties. The detailed results and analysis for each aspect will be reported elsewhere.
2. Experimental details Ingots of four alloys, with nominal compositions (in atomic per cent) of Ti-47Al-lCr-lV-2SNb (alloy Gl) Ti-47.OA1-1.5Cr-0.5V-2.3Nb (alloy G8), Gi46.5Al-2Cr-3Nb-0.2W (alloy K5) and Ti-47.2Al1.5Cr-0.5Mn-2.6Nb-0.15B (alloy K7), were prepared by skull melting and casting, followed by hot isostatic pressing at 1260 “C for 3 h under an argon pressure of 173 MPa. Billets cut from the ingots were isothermally forged at 1150 “C. Alloy Gl billets were given a 88% height reduction by single-step forging [5] and other alloy billets a 9 1% reduction by two-step forging [ 18,221. Differential thermal analysis combined with metallographic observations were used to determine the a transus T, for each alloy on forged material. Various types of microstructures, DP, NL, FL, refined FL (RFL) and thermomechanically treated lamellar (TMTL), were controlled through specific recrystallization treatments, followed by controlled cooling and then an aging at 900 “C for 6 h. The recrystallization treatment producing a specific DP microstructure for each alloy was conducted at a temperature 40-70 “C below T,. FL microstructures for alloy Gl and G8 having various GSs (300-2600 pm) were produced through recrystallization treatments at temperatures between T, + 1 “C and T, + 35 “C, while RFL structures having a 300 pm GS in alloy K5 were controlled in a temperature range from T, + 15 “C to T, + 3 5 “C. TMTL microstructures having a 150 pm GS were generated in alloy K7 by recrystallization at temperatures between T, + 5 “C and T, + 40 “C. Lamellar spacing for each alloy lamellar microstructure was varied by employing alloy-microstructure-specific cooling methods and rates [18,23,24]. Tensile testing was conducted entirely in air. Tensile tests for data collection were conducted on asmachined specimens with a gauge diameter of 3.4 mm at temperatures between RT and 1000 “C at strain rates in the range from 1 X lo-” s-r to 1 X 10-l s-‘. RT tensile deformation behavior was observed on electropolished flat gauge section specimens. Fracture resistance behavior at RT was investigated for Gl and G8 alloys in vacuum on single-edge-notched (SEN) specimens with ion-milled surfaces in a scanning electron microscope equipped with a loading stage [ 14,161. Displacement rates (DRs) employed for these
tests ranged from 1 x 10-j 4 x 10-l mm s-l. Effective strains near the crack tip were calculated by measuring the displacements with respect to the crack tip. Fracture toughness values were measured in air using mechanically polished compact tension specimens under a loading rate of 1.5 MPa m”* s-l. Microstructures, deformation features and fracture surfaces were observed using light microscopes, the backscattered electron imaging technique, and scanning electron microscopy (SEM).
3. Results 3. I. Alloys and microstructures The chemical analyses of the forged materials for four alloys investigated are listed in Table 1. Also listed are various interstitial levels, with the total interstitial content ranging from 750 to 880 wtppm. The forging conditions-parameters, T,, heat treatment conditions and the resulting microstructures are listed in Table 2. Fig. 1 shows the DP microstructures of these alloys consisting of small (lo-40 pm) y grains with small amounts (2%-5%) of a,-Ti,Al plates-particles dispersed in the matrix. The finest DP microstructure (10-l 5 pm) was obtained in alloy K5 which contains an additional second phase, /?, imagined the brightest in Fig. lc. Fig. 2(a) shows lamellar microstructures consisting of large (above 300 pm) lamellar grains, typical of alloys Gl and G8, with small amounts of fine (5-15 pm) y grains along lamellar grain boundaries (GBs). The lamellar spacing for the FL materials was relatively coarse, ranging from 1.6 to 4.8 pm, because of the slow cooling rates employed for the microstructure. The RFL structure of alloy K5 consists of a 300 pm lamellar grains with fine (5-20 pm) GB 7-P grains (Fig. 2(b)). Th e p resence of /3 phase (Fig. l(c)) during heat treatment is important in controlling GS, which has been recently documented [24]. Because of relatively fast cooling rates that can be applied, the lamellar spacing for the RFL structure of alloy K5 ranged from 0.4 to 0.8 pm. The TMTL structure in alloy K7 consists of about 150 pm size lamellar colonies and ragged boundaries, with lamellar spacing ranging from 0.7 to 1.3 pm. Controlling such fine spacing is possible because the TMT microstructures in boron-containing alloys could be achieved at almost any cooling rates [ 181 and the RFL microstructures can also be obtained under widely ranging cooling rates. The TMT microstructure in non-boron-containing wrought alloys was recently discovered and patented [l&23,25]. The TMT microstructures have been recognized to be much more easily produced beneficially, when small amounts
Y.-W. Kim Table 1 Composition
of y-TiAl alloys developed
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Materials Science und Engineering A1921193 (1995) 519-S33
521
for the experiments
Alloy
Alloy composition (at.%)
Interstitials (wt.ppm)
Cl G8 K5 K7
Ti-47.OAkl.OCr-0.9V-2.6Nb Ti-47.OA1- 1SCr-OSV-2.3Nb Ti-46.5Al-2.1 Cr-3.ONb-0.2W Ti-47.2AI- 1.SCr-O.SMn-2.hNb-0.15B
78O(O+C+N+H) 75O(O+C+N+H) 805(O+C+N+H)
MO(O+C+N+H)
Table 2 Processing and heat treatment conditions and resulting microstructures for the alloys in Table 1 Alloys
Forging’
Cl
SSF, 88%
1355
DP: T, - 70 to r, - 40 NL: T,-20 to T,- 10 FL: TcJ,+1 to T,+35
20-50 ym GS 30-l 50 pm GS 250-2600 pm GS
G8
TSF, 9 1%
1362
As for Gl
As for Gl
K5
TSF, 91%
1325
DP: T,I-50 NL: TI,- 10 FL: T,+ 15 to To+35
lo-15ymGS lo-120pm GS 200-400 pm RFL
K7
TSF, 9 I %
I365
DP: To-70 TMT: T, + 5 to T, + 40
“SSF. single-step
Recrystallization (“C)
temperatures
Microstructures
lo-15pmGS 80-200 pm GS
forging; TSF, two-step forging.
(0.05-0.3 at.%) of boron are added in wrought alloys, which results in 50-300 pm size colonies depending on the aluminium content, boron amount and heat treatment temperature and time [2,4,18]. Extensive studies on the refining and the controlling mechanism of TMT structures in wrought alloys are under way to explore their engineering significance [4,18].
3.2. Tensile tests RT tensile and fracture properties of all alloys are listed in Table 3 and tensile properties of Gl and K.5 at temperatures between RT and 1000 “C are listed in Table 4. Drastic differences exist in the temperature dependence of the tensile properties for the different microstructures as shown in Fig. 3, where plots are made on alloy K5 data in Table 4 and alloy Gl data from previous data in Ref. [22]. For the DP structures for these alloys, the UTS remains unchanged up to 800 “C while the YS decreases gradually over the same temperature range. Both strengths decrease rapidly at higher temperatures. The alloy Gl FL specimens with large grains exhibit fairly low strengths at RT, but superior strength retention at elevated temperatures, losing not more than 25% of the RT strength at 950 “C. Alloy K5 RFL microstructure also provides high
temperature strength retention, however, with much higher strengths than large-grained FL structures. RT strength for this RFL material is comparable with that of the DP material in spite of a grain size by a factor of 20. The DP-structured material exhibits higher ductility than the FL material throughout the temperature range examined. The brittle-ductile transition (BDT) temperature lies between 650 “C and 700 “C for the DP materials, around 800 “C for the RFL structure, and at 820 “C for the FL material with a 1100 pm GS. Tensile elongation for the large-grained ( 1100 pm) FL material is 0.8% at RT and 2.5% at 800 “C. RT tensile data in Table 3 are plotted in Fig. 4 as a function of microstructure type and GS. Both strength and ductility decreases as GS increases, especially within a class of microstructures. It is noted that the tensile strengths of lamellar materials (FL, TMTL and NL) having 100-300 pm grains are comparable to, or higher than, those of DP materials. NL material having about 100 pm grains shows an extraordinary combination of strength (near 700 MPa) and ductility (greater than 2.0%). The deformation and fracture behavior observed for the DP specimens is shown in Fig. 5. The RT fracture is dominantly transgranular cleavage-like failure, exhibiting river patterns (Fig. 5(a)) in spite of the 2.5% ductility and considerable deformation activity evident in
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Fig. 2. Optical micrographs of the lamellar structures: (a) alloys G 1 and G8 FL, (b) alloy K5 RFL and (c) alloy K7 TMTL microstructures.
I-15pm cig. 1. DP microstructures of (a) alloys G 1 and G8 under optical nicroscopy, (b) alloy K5 under optical microscopy, and (c) alloy 3 in backscattered electron imaging.
the material (Fig. l(a)). Fig. l(a), a micrograph taken from the area below the tensile fracture surfaces tested, shows planar-type deformation features evident in the DP specimens. This trend continues up to the BDT temperature. At 700 “C, the fracture is predominantly intermanular (Fig. 5(b)\ with a few cleavage events.
Y-W. Kim Table 3 Controlled
microstructures
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and lamellar spacing and corresponding
room temperature
523
tensile and toughness
properties
of the alloys
in Table 1 Alloys
Microstructure (pm)
Lamellar spacing (pm)
Alloy Cl G8
DP (20-40) NL(-80) FL (250-2600)
KS
DP(lO-15) NL (10-100) RFL ( = 300)
K7
DP(lO-15) TMT ( = 150)
UTS (MPa)
Elongation @)
Toughness (MPa rn’!‘)
Ref.
EPa)
557-601 702 588-300
2.5-3.8 2.8 1.2-0.5
K,,=K,,,=lO-12
1.6-S
420-485 511 508-280
[51 [51 [5,191
0.4 0.8
46.5 550 475 445
580 685 550 520
2.9 2.1 1.1 1.0
K ,,,=21.5 K max= 20.0
0.7 1.4
435 385 375
555 525 490
2.5 2.8 2.4
K max= K msx=
K,, = 13-21; K ma*= 20-38
K,,=K,,,=lO-11
[ 18,22,24] [ 18,22,24] [ 18,22,24]
18.5 17.0
YS, yield strength; UTS, ultimate tensile strength.
Table 4 Alloy K5 tensile properties at various temperatures and 2 X 10e4 s-’ [ 18,22,24] Microstructure (pm)
Property
Values at the following test temperatures
(“C)
RT
600
700
800
375 480 5.2
319 360 49
DP(lO-15)
YS(MPa) UTS (MPa) Elongation (%)
465 580 2.9
385 528 3.7
RFL ( = 300)
YS(MPa) UTS (MPa) Elongation (%)
473 557 1.2
408 525 1.7
More deformation is evident within grains; however, the planarity of deformation remains unchanged. At higher temperatures, specimens begin to neck with significant elongation. Tremendous amounts of deformation activity with multiple slip and twinning take place at 800 “C (Fig. 5(c)), resulting in an tensile elongation of 50% or more (Table 4 and Fig. 3) and almost entirely intergranular fracture. At still higher temperatures, the fracture surfaces show oxidation debris, and the longitudinal cross-sections reveal grain refinement as a result of dynamic recrystallization taking place, leading to enhanced ductility 17,181. The fracture at temperatures above 900 “C involved the formation of voids. FL specimens exhibit three characteristic fracture features at RT: delamination or interlamellar fracture (Figs. 6(a) and 6(b)), t ranslamellar fracture (Figs. 6(a), 6(b) and 6(c)), and stepwise fracture (Fig. 6(b)), depending on the lamellar orientation with respect to the crack path or stress axis. Translamellar fracture or cracking can occur in the crack arrester orientation
385 510 8.1
900
340 418 19.0
1000
278 284 > 20.00
(lamellar boundaries nearly perpendicular to the crack propagation direction (Fig. 6(c)) or in a more complex manner (Fig. 6(d)). In both cases, cleavage-type features with river patterns are often shown and traces of planar deformation bands parallel to the lamellar interfaces are frequently observed on the fracture surfaces. The planar deformation bands are observed in selective grains on the electropolished specimens during deformation (Fig. 7(a)). Translamellar deformation lines are observed occasionally during deformation, but they appear to form more easily during fracture, within the plastic wake or process zone of the crack (Fig. 7(b)). The elevated temperature fracture of lamellar specimens is characterized by translamellar fracture and delamination or secondary cracking, as observed after failure at 750 “C (Fig. 8(a)) and 800 “C (Fig. 8(b)). The delamination spacing is roughly 15 ,um, and the extent of splitting appears to increase with temperature, often leading to the interlamellar fracture of grains (Fig. 8(b)). Translamellar cleavage fracture was rarely
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above 800 “C in any alloys. Intergranular observed fracture was observed occasionally in the alloy K5 RFL specimens (Fig. 8(b)), which appears to be related to the presence of GB y grains. In some cases, these GB grains appear to help to block or deviate the propagation of the delamination cracks (Fig. S(a)). At 900 “C, although not shown here, lamellar microstructures become unstable in the necked regions under the strain rates used, and they begin to recrystallize into fine (2-5 ,um) y grams. In this case, ductile fupture involving void formation starts to prevail. Twins were observed on the 900 “C fracture surface, as previously reported [7], indicating that twinning is still an important deformation mechanism at this temperature. However, lamellar microstructures can be unstable even at 800 “C when the tensile elongation becomes more than 8%, as was the case for tests conducted in vacuum at a strain rate of 1 x lo-” s-r [16].
?? mrS
Temperature (“C)
(a)
Al 92/l 93 (I 995) 519-533
3.3. Fracture toughness tests
0
600
600
200
1000
Temperature (“C)
W
Fig. 3. Tensile properties vs. test temperature in alloy Gl and K5 for DP and lamellar (FL and RFL) microstructures. Duplex I NL
l-9 1 IL
I
-
\
01
’ 0
Alloy Gl 0 Alloy K5 A
0 A
. A
Alloy K7
0
??
ility
\y
I 500
I 1000
700
I
I
I
n
I 1500
0
v
I 2000
600
{
2oo
100 2500
Grain Size (urn)
Fig. 4. RT tensile properties of y alloys as a function of microstructure and GS.
RT fracture toughness K,,, values of alloy Gl are lo- 12 MPa rn’iz for DP microstructures and range from 24 to 33 MPa ml’* for FL microstructures depending on GS (250-2000 ,um) (Table 3). The toughness values range from 20 to 22 MPa ml’* for alloy KS RFL material with a 300 pm GS and from 16 to 18 MPa ml’* for alloy K7 TMTL material with a 150 pm GS. The slight increase in toughness in each range appears to be related to the decrease in lamellar spacing. K-resistance curves of alloy Gl in Fig. 9(a) show that initiation toughness K,, and crack resistance behavior differ substantially between the DP and the large-grained FL structures [ 14,16,21]. K,, values for DP material are about 11 MPa ml’* and range from 16 to 21 MPa m1i2 for FL materials, depending on displacement rate. For FL structures, the initiation toughness may vary with lamellar orientation, but has not yet been measured systematically. No resistance to crack growth was observed in DP material (Fig. 9(a)), with K,, remaining unchanged with crack growth. For FL material, substantial resistance to crack propagation was observed, which resulted in saturation or maximum toughness K,,, values ranging from 25 to 42 MPa ml’*. For the tests reported here, on average, two FL grains are fractured by the crack in each specimen in the course of a 2 mm propagation distance (Fig. 9(a)). The variations in resistance were also related to the lamellar orientation with respect to the loading direction or the main crack direction. Some effects of DR on the resistance were observed. The specimens with high K,,, values ( 3 7 MPa m1/2 or higher) generally showed the main cracks growing translamellarly and
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Fig. 5. ‘Tensile deformation and fracture of alloy K5 DP material showing (a) mostly transgranular cleave (SEM) at RT, (b) mostly intergranular fracture (SEM) at 600 “C, (c) deformation structure (optical microscopy) at 800 “C, and (d) dimple fracture (SEM) at 800 “C.
bridging ligaments failing interlamellarly by shear, as shown in Fig. 10. The displacements near the crack tip at were measured on the surface K=K,,=16MPam1~* of an alloy Gl FL SEN specimen having large lamellar grains (above 1000 pm) with laths inclined at 8 = 48” to the notch direction ( 8 = 0’) which is perpendicular to the loading direction [14]. The isostrain contours constructed using the effective strains calculated from the measured displacements show a large plastic zone
with a butterfly shape which is roughly symmetric with respect to the lamellar direction [26]. When the data of effective strains are plotted against radial distance r from the crack tip, a linear relationship exists on a double-logarithmic scale in a given direction, as shown in Fig. 9(b) for three directions, 8= O”, 45” and 90”. Two important aspects are obvious from Fig. 9(b). The effective strain is orientation dependent, and the maximum strain, or fracture strain, ranges from 10% to 2.5% depending on direction, with the largest straining
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Fig. 6. RT tensile fracture surfaces (SEM) of lamellar materials showing: (a) mostly translamellar fracture with a few interlamellar fracture events in alloy K5, (b) combination of interlamellar and stepwise fractures in alloy K5, (c) translamellar cleavage with river patterns and lamellar slip bands in alloy G8, and (d) complex translamellar cleavage-type fracture in alloy G8.
taking place in the soft orientation (0 = 0’). This indicates that under a triaxial loading condition the lamellar structure can be deformed to a large extent in any direction. This is in contrast for the DP structure which shows a small maximum effective strain of about 3% in front of the crack tip that was obtained from similar displacement measurements on a Gl DP specimen (Fig. 9(b)).
4. Discussion 4.1. Tensile deformation andfiacture The ductility of y alloys decreases, in general, with increasing GS, and so does tensile strength (Fig. 4). The ductility for lamellar structures decreases from 2.5% to 0.5% as the GS increases from 250 pm to 2500 ,um [5]
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3mm
Fig. 7. Alloy G8 FL flat specimen fractured at RT showing on the electropolished surface (a) deformation of some specific grains only and (b) plastic deformation structure in the crackwake or process zone.
4). The Hall-Petch relationship between YS and GS in FL microstructures exists with a large Hall-Petch constant Ic,,= 5 MPa m l/* for grain sizes from 250 to 2600 pm and yield strengths from 290 to 500 MPa [22], as shown in Fig. 1 l(a). This value is compared with those ( 1 .O-1.2 MPa ml’?) for fine y-DP microstructures [27-291 and that (0.5 MPa m’12) for PST crystals in a hard orientation [ 131 (Fig. 1 l(b)). This unusually large Hall-Petch constant is not readily explained, but may be related to the strong anisotropic flow stress behavior [3,12,13] of the lamellar structure. The gauge diameter to GS ratio may also play a role in the unusual relation. The ratio ranges from 1.3 to 14 for the FL specimens. Therefore, the specimens containing large grains have only a few grains within the gauge cross-section, with the result that one or two soft grains control the yielding process. An investigation is underway to clarify this uncertainty [ 181. (Fig.
527
15pm . ..-
Fig. 8. Elevated temperature tensile fracture (SEM) of lamellar structures showing (a) translamellar fracture with delamination or lamellar splitting and intergranular fracture of GB y grains at 750 “C in alloy Gl and (b) translamellar, interlamellar and intergranular fractures in alloy K5 at 800 “C.
The inverse relationship between ductility and GS (Fig. 4) was recently analyzed by assuming the presence of a microcrack comparable with the average GS and subsequently estimating the stress intensity of the crack and the resultant ductility due to crack propagation [ 141. However, the source of the assumed microcrack was not identified clearly and the role of general plasticity was not defined. On the contrary, the decreases in tensile strength and ductility with increasing GS in lamellar material can be explained using the characteristic anisotropy of tensile properties of lamellar structures [ 12,131, and assuming that the fracture is controlled by the crack nucleation process [30] involving the coalescence of
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FiTlAir 0
I 0
(a)
0.5
I
I
I
1.0
1.5
2.0
Aa
LO
03
r (mm)
Fig. 9. Fracture behavior of alloy Gl at RT in DP and FL conditions: (a) K-resistance curves and (b) effective strain field near crack tip measured at K,, = 16 MPa ml/* in the three directions with respect to the notch direction (8= 0”) which is perpendicular to the loading direction (LD). The effective local fracture strain of 15%-20% for the lamellar microstructure is compared with that of 3% for the DP structure.
dislocations under shear stress [ 3 11. In this process, the local magnified shear stress r,, at the pile-up is known to be a function of grain size d and applied stress r, and can be expressed as
Fig. 10. Compact tension specimen surface (electropolished) during RT testing showing (a) the major crack, a microcrack in the adjacent grain, a ligament area, and both interlamellar and translamellar slip-twin bands (optical microscopy) and (b) a magnified view (SEM) of the ligament area showing fine microcracks formed in the translamellar deformation direction [19,22].
tp = C, dra2 or, conveniently, aP = Cdoa2
(1)
where C, and C are constants. Using the relation (1) and the flow curve equation cs= KE”, the local flow curves for two different grain sizes d, and d2 with d, < d2 are constructed as shown in Fig. 12(a). Both materials fracture at the same local fracture stress oPr by initiating cracks but at different fracture strains elf and cZf, with elf > eZf. Using these fracture strains, the flow curve equation, and the relation o, = C,d- u2 (for given ur) derived from relation (l), the tensile stress-strain curves for the two materials having different GSs can be constructed as shown in Fig. 12(b). It is
then clear from Fig. 12(b) that the material having smaller grains shows not only higher flow stress but also greater fracture strain, resulting in its fracture stress calf being higher than that oa2f of large-grained material. Ductility is then primarily controlled by the amount of general yielding which will be greater for small grain sizes. In the FL material, cracking and subsequent fracture precede general yielding resulting in nonuniform deformation and low strains to failure [20,22,32]. Slip bands observed within the grains having easy orientations (Fig. 7(a)) are consistent with
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d W)
3000 300 100 30 \\//I0 I 8001
- 800
-
EL
-
(Kim931
w - 600
- 400
- 200
(“makosh, 92)
0
400
800
1200
1600
(a) d-'" V"')
6oor-----l
(b)
Tensile Strain (%)
Fig. 11. (a) Hall-Petch relation obtained from the data of alloy Gl-G8-K5 FL materials is compared with those reported for other microstructural forms and (b) stress-strain curves of a TiAl PST crystal in various orientations which were drawn by rearranging the data of Ref. [ 121.
Fig. 12. Estimated flow curves of y alloys for two different grain sizes (d, < dZ) in two different stress conditions: (a) local stress-strain curves near the dislocation pile-up generated by the stress concentration ur and (b) the materials’ general tensile stress-strain curves under applied stress a,. For details, see text.
529
the deformation bands parallel to the lamellar directions observed on fracture surfaces in Figs. 6(c) and 6(d). This planar deformation leads to dislocation pileups nucleating microcracks, often in an adjacent grain (Fig. 13), which is usually preceded by either translamellar slip-twining or additional interlamellar deformation, and may result from plastic incompatibility stresses between grains. The microcrack usually forms in the translamellar direction (Fig. 13), although interlamellar microcracking should be possible depending on the lamellar orientation. While this plausible mechanism of plasticity-induced crack nucleation is now evident, the event has not been quantitatively tied to a stress or strain. Consequently, it is not yet entirely clear whether uniaxial tensile ductility and failure are a crack nucleation controlled event [31], or if propagation to a critical size is required. Tensile ductility is characterized by the BDT temperature which varies from 650 “C to 820 “C (Table 4 and Fig. 3) depending on alloy composition and microstructure. For a given alloy, the BDT temperature increases with GS or with decreasing RT ductility. It is also strain rate sensitive, especially for fine microstructures, as was observed in a DP microstructure [ 18,331. The enhanced ductility above the BDT temperature appears to be the direct result of increases in both twinning-slip activity and systems, as was discussed earlier. Some claims that dynamic recrystallization controls BDT temperature have yet to be confirmed. For DP microstructures, tensile strengths at temperatures below the BDT temperature depend on cooling rate and method [5,18,23,24] employed, with higher cooling rates yielding higher strength levels. The (72phase, whose morphology and distribution strongly control mechanical properties, is present in fine particle (instead of plate) forms with fast cooling rates. Finer lamellar spacings result in NL and FL microstructures with increasing cooling rate. Decreases in lamellar spacing appear to result in some increases in tensile strengths and ductility (Table 3), but this relation has not been quantitatively characterized. Much work is needed to quantify the relationship between cooling rate-scheme and mechanical properties. Cleavage and transgranular (translamellar and interlamellar) cracking are the dominant fracture modes of RT for both microstructures, although the fine DP material shows a mixture of transgranular and intergranular failure [ 17,20,22]. Cleavage-like brittle fracture was also observed in PST crystals even for failure occurring parallel to the lamellar interfaces with more than 20% plastic strains [ 12,341. At temperatures above BDT temperatures, intergranular fracture prevails in DP material and delamination and intergranular fracture become important for FL material [17,22,33]. In spite of the increased deformation
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Fig. 13. Nucleation
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of microcracks
I
I
at dislocation
pile-ups against grain boundaries
Spacing (pm)
I
Fully Lamellar
!
in FL materials.
50-
.?.
40.
(Expexted) ___________-----
,
4
1
,
0.44
,
0.25
1.5
2
AllDv
: 30
20
10 (150 W) 0 -0
...~‘~~~.l~~~.‘.~.~0.5 Lamellar
I
:
0
I
I
I
I
500
1000
1500
2000
IO
2500
1 Spacing,
XjlR
(pm-‘R)
Interrelations between fracture toughness, lamellar spacing d, in polycrystalline lamellar y alloys.
GS, and
Grain Size (pm)
Fig. 14. Inverse ductility-toughness K,, and K,,, relationships observed in alloys Gl and G8 but characteristic of y-TiAl alloys. Some data of alloys KS and K7 are plotted for comparison.
1
Duplex
1
Duplex
1
:&:, ;:,(-:‘, strain (a) Tensile Curves
Strain (b) Crack-Tip
Strain Local Yielding
Fig. 15. Comparison of general tensile yielding and crack tip local yielding behavior: (a) measured tensile stress-strain curves typical of DP and FL microstructures and (b) estimated local stress-strain flow behavior of the material in the near-tip plastic zone for these microstructures.
activity with temperature, transmission of slip and twinning across GBs or lamellar interfaces appear to remain difficult [22]. The apparent GB interface weakening with increasing temperature has yet to be elucidated.
4.2. Fracture toughness and toughening mechanism The K-resistance curves can be divided into two components; that is, the resistance to crack initiation or initiation toughness K,, and resistance to crack growth. K,, is essentially a measure of the plastic deformation before the crack propagation, from which the local tensile properties can be estimated within the plastic zone, as discussed in the next section. DP material exhibits small plastic strain near the onset of
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crack extension and no resistance to crack propagation, whereas the FL structure yields large crack tip plastic strains and increased resistance to crack propagation (Fig. 9(a)) [ 14- 161. At RT, FL specimens show K,, = 16-21 MPa rn’i2 and K,, = 25-42 MPa rn’i2 and the DP specimens have KIc = K,,,,, = 11 MPa mii2 (Fig. 7(a)). The initiation toughness K,, is related to the strain measured near the crack tip (Fig. 9(b)). As shown, the fracture strain (at 0 = 0”) in the plastic zone can be 20% for FL material but about 3% for DP material. The occurrence of crack growth resistance is caused by the work expended mainly in failing the ligaments formed in the wake of crack [ 14,211. Extensive deformation activity, both interlamellar and translamellar, is seen in the ligament region in Fig. 10. This particular ligament is oriented for interlamellar shear failure linking the main crack and the microcrack which are both translamellar type. This crack orientation results in higher crack resistance than the other type of orientation, that is interlamellar main crack path with translamellar ligament failure [22]. Ultrafine microcracks are formed on the deformation planes during the translamellar deformation (Fig. 10(b)), and crack propagation can be shielded by these cracks and some other off-path cracks shown on the GB. Another, but less effective, source of growth resistance can be the increased deformation amounts due to the increased plastic zone size with crack growth. The sensitivity of toughness to displacement or loading rate increases with temperature; for example, K,, and K,,, values for the large-grained FL specimens at 800 “C were measured to be 35 MPa m1j2 and 79 MPa ml” respectively [16]. The absence of crack growth resistance behavior in DP structures is observed at least up to 600 “C and is due to the lack of bridging ligaments.
4.3. Ductility-toughness
relationship
The inverse relation between ductility and toughness (both K,, and K,,,), summarized in Fig. 14, can be explained by analyzing the bulk tensile behavior and the crack tip yielding behavior [22,32]. Tensile ductilities of polycrystalline material are primarily controlled by the amount of general yielding which will be greater at small grain sizes. In toughness testing of largegrained (above 500 pm) FL specimens, however, the crack tip plastic zone (estimated to be about 400 pm) encompasses only one or part of a grain in front of the crack tip, thereby confirming the deformation behavior to the characteristics of those few grains along the crack tip edge for thin specimens. K,c is then closely related to the intrinsic ductility of the lamellar structure and can be used to estimate the tensile stress-strain
531
curve within the plastic zone (22,321. The near-tip effective fracture strain in the lamellar structure in Fig. 9(b), about 20% at 0 = O”, is consistent with the tensile fracture strain directly measured on a TiAl PST crystal (Fig. 1 l(b)) [12,13]. In DP materials large plastic volumes are not sampled because of the small grain sizes. Fig. 15 compares the measured general tensile curves (Fig. 15(a)) and the stress-strain curves within the near-tip plastic zone (Fig. 15(b)) for coarse (larger than plastic zone size) FL and fine (lo-40 pm) DP materials [22,32]. The stress-strain curves in the plastic zone were estimated using the following relation [35] between K,, and local stress-strain behavior: Klc2 = CnEo
Y
EIf
(2)
where C is a constant, IZ the local work hardening coefficient, E the elastic modulus, oy the material’s YS, and qf the local true fracture strain which corresponds to the maximum effective strain measured near the crack tip (Fig. 9(b)). The K,, values used for the estimation were 10.5 MPa m’i2 and 16 MPa m’i2 respectively for the DP and the FL materials (see Fig. 9(a)). For FL materials with grain sizes greater than 500 pm, the plastic zone sizes rP at the onset of crack initiation range from 400 to 500 ym as calculated using the relation rP= a( K/o,)~, and it is thus contained within a grain in front of the crack tip. In this case, the local fracture strain, about 15%, was taken as an averaged value from Fig. 9(b). For the DP material, the material within the plastic zone (which is estimated to be 300 pm in size and contains many randomly oriented grains) is fairly isotropic and represents the bulk material with a fracture strain of about 3%, and thus the near-tip stress-strain curve is essentially identical to the tensile curve (Fig. 15(b)). The results in Fig. 15(b) clearly show that the socalled inverse ductility-toughness relationship is caused simply by the difference in sampling between tensile testing and toughness testing. Within the plastic zone, the material obeys the classical relation between tensile properties and initiation toughness expressed by relation (2). 4.4. Competition between grain size and lamellar spacing in toughness As was described in the previous section and shown in Fig. 14, the fracture toughness K,c and K,,, increases with lamellar GS up to about 500 pm, but appears independent of GS for larger sizes. In this constant-toughness region against GS, however, it was found that the fracture resistance increases with decreasing lamellar spacing II, as was measured in
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alloys Gl and G8 specimens in a il, range from 4.8 to 1.6 pm [19]. Fig. 16, where the data in Ref. [19] are selectively plotted, shows the Hall-Petch relation between toughness (especially K,,) and lamellar spacing, as follows: K Ic =K 0 +k;l L -r/l
(3)
where K, may be defined as the intrinsic K,, value and k is a Hall-Petch constant. The increase has been explained in terms of increased stresses required for translamellar slip-twinning [19]. The flow stresses of the lamellar structure in the hard orientation are higher than those of polycrystalline lamellar materials and increase to much higher levels with decreasing lamellar spacing (Fig. 1 l(a)). Interestingly, K,, of the FL material with large lamellar spacing becomes the same as that of DP material, that is about 10 MPa ml/*, when the lamellar spacing reaches about 10 pm (Fig. 16). This suggests that the lamellar structure with coarse lamellar spacing may behave like a fine-grained material by losing the anisotropy in deformation. In this case, the lamellar interfaces and domain boundaries are expected to act as GBs effectively pinning the dislocations. When the GS is smaller than the plastic zone size, then the GB effect on the toughness will begin to prevail over the lamellar spacing effect. Although not totally quantitative, this prediction is evidenced from the test results (Table 3) for K5 and K7 alloy compact specimens containing respectively 300 pm and 150 pm lamellar grains (Fig. 16). In this GS range, the variation in lamellar spacing between 0.4 and 1.3 pm does not alter the toughness K,,, appreciably, indicating that the Stroh fracture stress within the plastic zone only slightly depends on the variation in lamellar spacing. The large drop in toughness in this GS range, compared with that in the large-grain range, is thus mainly the GS effect. Systematic studies are underway to quantify the competition between GS and lamellar spacing using large, thick specimens [ 181, and the results will be reported elsewhere.
5. Conclusions RT and elevated temperature deformation and fracture behavior of y-TiAl alloys was investigated under tensile and compact-tension loading conditions. Tensile properties strongly depend on microstructure and temperature. FL structures yield lower ductility and strength, but greater strength retention at high temperatures, and higher BDT temperatures than DP structures. A Hall-Petch relationship with an unusually high Hall-Petch constant was observed to exist between YS and GS for FL structures; however, the
Al 92/l 93 (I 995) 519-533
mechanism for this phenomenon is unclear. The dependence of ductility on GS is analyzed to be limited by the formation of strain incompatibility induced microcracks. For both microstructures, fracture is preceded by plastic deformation followed by cleavagelike failure. Fracture toughness is also a strong function of microstructure and temperature, with higher toughness achieved in FL structures and also at higher temThe high fracture toughness for peratures. large-grained FL material is explained in terms of the characteristic of the lamellar structure observed at the crack tip, that is large plasticity and the ability to form ligaments. For a given large-grained lamellar material, fracture toughness is a function of lamellar spacing with Hall-Petch relation. The inverse the ductility-toughness relationship is explained by comparing the general tensile yielding and the near-tip local plasticity measured at the onset of crack initiation. DP materials fail transgranularly at RT and intergranularly at high temperatures. FL materials show predominantly transgranular fracture over a range of temperatures as long as the microstructures are stable. Decreasing GS or GB y phase, however, promotes intergranular fracture at elevated temperatures.
Acknowledgments
The author appreciates the technical discussions with Drs. D.M. Dimiduk and KS. Chan and the technical assistance from Ms. Sonya Boone. Most of the work in the paper was conducted under Contract F33615-91-C-5663, US Air Force Wright Laboratory, Materials Directorate.
References Y.-W. Kim, J. Met., 41(7) (1989) 24-30. Y.-W. Kimand D.M.Dimiduk,J. Met., 43(8)( 1991)40-47. M. Yamaguchi and H. Inui, in Structural Intermetalks, TMS, Pittsburgh, PA, 1993, pp. 127-142. Y.-W.Kim,J. Met., 46(7)(1994)30-39,7. Y.-W.Kim,Acta Metall. Mater., 40(6)(1992) 1121-1134. D.S. Shih, S.-C. Huang, G.K. Starr, H. Jang and J.C. Chesnutt, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property
Relationships
in Titanium
Aluminides
and
A/loys,TMS, Pittsburgh, PA, 1991, pp. 135-148. S. Krishnamurthy and Y.-W. Kim, in Y.-W. Kim and R.R. Relationships in Boyer (eds.), Microstructure/Property Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991,~~. 149-163. W.E. Dowling, B.D. Worth, J.E. Allison and J.W. Jones, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991, pp. 123-134. A.W. James and P. Bowen, Mater. Sci. Erg. A, 153 (1992) 486-492.
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[ 141 [15] [16]
[ 171 [ 181 [ 191 [20] [21] [22]
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H.A. Lipsitt, D. Schechtman and R. Schafrik, Metafl. Trans. A, 6(1975) 1991-1997. S.-C. Huang and E.L. Hall, Metall. Trans. A, 22 (1991) 427-438. H. Inui, M.H. Oh and M. Yamaguchi, Acta Metall. Mater., 40(1992)3095-3104. Y. Umakoshi, T. Nakano and T. Yamane, Mater. Sci. Eng. A, I52(1992)81-88. K. Chan and Y.-W. Kim, Metall. Trans. A, 23 (1992) 1663-1677. H.E. Deve, A.G. Evans and D.S. Shih, Acta Metall. Mater., 40(1992) 1259-1265. K. Chan and Y.-W. Kim, Metall. Trans. A, 24 (1993) 113-125. Y.-W. Kim and D.M. Dimiduk, Mater. Res. Sot. Symp. Proc., 288( 1993) 671-677. Y.-W. Kim, research in progress, 1992-l 994. K. Chan and Y.-W. Kim, Acta Metall, Mater., in press. G. Malakondaiah, Y.-W. Kim and T. Nicholas, Ser. Metall., 3U( 1994) 939-944. K. Chan and Y.-W. Kim, Metall. Trans. A, 25 (1994) 1217-1228. Y.-W. Kim and D.M. Dimiduk, Proc. JlMIS-7 on High Temperature Deformation and Fracture, Japan Institute for Metals, Sendai, 1993, pp. 373-382.
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[23] Y.-W. Kim and D.M. Dimiduk. US Patent 5,226,985, July 13, 1993. [24] Y.-W. Kim and D.M. Dimiduk, filed for a US patent, 22 April 1994. [25] P. McQuay, D.M. Dimiduk and Y.-W. Kim, filed for a US patent, May 1993. (261 Y.-W. Kim and K.S. Chan, unpublished results, 1994. [27] S.-C. Huang and D.S. Shih, in Y.-W. Kim and R.R. Boyer (eds.), Microstructure/Property Relationships in Titanium Aluminides and Alloys, TMS, Pittsburgh, PA, 1991, pp. 105-122. (281 P. McQuay, M.S. Thesis, Wright State University, OH, 1992. [29] C. Koeppe, A. Bartels, J. Seeger and H. Mecking, Metal/. Trans. A, 24(1993) 1795-1806. [30] A.N. Stroh, Philos. Mag., 46 (1955) 968-972. [31] A.N. Stroh, Proc. R. Sot. London Ser. A, 233 (1954) 404. [32] Y.-W. Kim, in Intermetallic Compounds, Proc. 3rd Jpn. Int. SAMPE Symp., Chiba, 1993, pp. 3 lo- 13 17. [33] i;$;,pfert, Y.-W. Kim and D.M. Dimiduk, Mater. Sci. Eng. [34] M.H. Oh, H. Inui and M. Yamaguchi, Acta Metall. Mater., #I (7) (1993) 1939-1949. [35] J. Langford, D.L. Davidson, K.S. Chan and G.R. Leverant, AFOSR Rep., Southwest Research Institute, May 1989.