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Plastic anisotropy and fatigue of TiAl PST crystals: a review
S0966-9795(96)OOOIZ-X
lmermetullics 4 (1996) S65-S75 " 1996 Elsevier Science Limited Printedin Great Britain. All rightsreserved 0966-97951961$15,00
ELSEVIER
Plastic anisotropy and fatigue of TiAl PST crystals: a review Y. Umakoshi, H. Y. Yasuda & T. Nakano Department of Materials Science and Engineering, Faculty of Engineering, Osaka University, 2-/ Yamada-Oka, Suita, Osaka 565. Japan (Received I November 1995: accepted 27 December 1995)
The cyclic deformation and fracture behaviour of TiAI PST crystals are reviewed. Anisotropy in cyclic hardening. fracture and fatigue life of the crystals deformed at 4> 0 and 45°. where 4> is the angle between the loading axis and lamellar planes. are discussed focusing on the different deformation substructures in '}'-domains. The effects of deformation temperature. addition of third elements and lamel1ar structure on fatigue life are also described. Copyright © 1996 Elsevier Science Ltd
=
Key words: A. titanium aluminides, based on TiAI, B. anisotropy. fatigue resistance and crack growth.
1 INTRODUCTION
lamellar spacing, domain size and the structure of lamellar and domain boundaries: the slip mode of the a2-phase was responsible for the
The discovery of a combination of high strength and good ductility in two-phase Ti-rich TiAI alloys has accelerated attempts to develop the alloys as potential high-temperature structural materials for use in airframes, car bodies, automobile engine valves and turbo rotors.':' Upon solidification of Ti-rich TiAI, a lamellar structure composed of y-matrix and a small amount of thin arplates forms during the a-to y-phase transformation. Since mechanical properties such as strength, ductility, fracture toughness and fatigue are known to be very sensitive to the orientation and microstructure of lamellae, knowledge of the effect of lamellar structure on plastic behaviour of TiAI is necessary to improve these properties. The successful achievement of polysynthetically twinned (PST) crystals, in which a single set of lamellae is unidirectionally aligned with no grain boundaries, has provided clear information on the effect of lamellar structure."? In monotonic deformation, the y-phase in equilibrium with the a2-phase in two-phase TiAI showed good ductility with low yield stress, but lamellar boundaries containing a2-phase-which were often accompanied by a loss of ductility-were required for strengthening. The anisotropy in strength, ductility and fracture behaviour was closely related to the angle between the loading axis and lamellar planes,
anisotropy.r" Since industrial materials are often subjected to repeatedly applied stress and strain, knowledge of the plastic behaviour under cyclic loading conditions is required for practical use in addition to monotonic deformation. The fatigue life, crack growth and fracture behaviour of polycrystalline TiAI alloys have been examined and good fatigue properties have been demonstrated at room and high temperatures.P:" The fatigue life of metals and alloys is controlled by the initiation of a microcrack and its propagation accompanied by an intergranular and/or transgranular crack path. Crack initiation in intermetallic compounds is known to play an important role in fatigue life because of their low ductility.P''" However, the lamellar boundaries of TiAI alloys often act as an effective barrier to the propagation of a crack. Similar to that in monotonic deformation, remarkable progress has been made in understanding the effect of lamellar structure on the plastic behaviour of TiAI under cyclic loading using TiAI PST crystals.v"?' This article reviews recent advances in our understanding of anisotropy and the mechanism of plastic and fracture behaviour of TiAI alloys under cyclic loading, concentrating on recent results on TiAI PST crystals. S65
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2 LOADING AXES AND SCHMID FACTOR OF y-DOMAINS The plastic behaviour of TiAI PST crystals is well known to exhibit a strong dependence on orientation c/J, where c/J is the angle between the loading axis and lamellar planes. Deformation in easy mode occurs only on (I I I) planes in the ')'-matrix parallel to the lamellar boundaries, and good ductility with low yield stress is shown at c/J = 45°.6.22 For hard orientations at c/J =0 and/or 90°, lamellar boundaries act as an effective barrier to the motion of dislocations, resulting in high yield stress but a loss of ductility. The angle c/J was chosen to be 0 and 45° to examine the effect of hard and easy mode on the cyclic deformation and fatigue properties of TiAI PST crystals. At c/J = 90° specimens were broken within a few cycles under push/pull loading because no slip systems other than pyramidal slip were activated in the a2-plates and the a2-plates hardly deformed. Figure 1 shows the stress axes of the arphase and each domain in the y-matrix of PST crystals c/J = 0 and 45°. The loading axis of the a2-phase occupies one position for each specimen, but since there are six types of ')'-domain in consideration of the matrix and twinning relation, the corresponding stress axes in each domain of the ')'-matrix are plotted as AM-CT in Fig. I(a), where the suffixes M and T represent the matrix and twinning relationship to each other. To examine the effect of deformation modes of twinning and slip in each domain, the loading axis is rotated on a (l I l) lamellar plane at c/J = 0 as shown in Fig. I(c). When the rotation angle (X) between the stress axis and the [I I 2] plane in AM domains changes from 0 to 30°, the combination of equivalent
domains on the basis of stress factors for the dominant slip and twinning systems is alternated. The two domains in the following parentheses are equivalent: (AM' AT). (BM• BT) and (CM, CT) at X =0; (AM' AT), (BM, CT) and (BT, CM) at X = 15°; (AM' BT), (BM, ~) and (CM, AT) at X = 30°. In particular, high stress factors for twinning exist in domains of (B M • ~), (CM, CT) at X =0, (B M , CT) at X = 15° and (BM , CT) at X = 30°. Therefore, as X varies at c/J = 0, the number of domains in which twinning would be activated can be changed to examine the effect of deformation twins on cyclic deformation and fatigue properties.i-" When x-value is not indicated in this paper, X =0 is chosen. 3 CYCLIC HARDENING AND FATIGUE LIFE UNDER A TOTAL STRAIN CONTROLLED CONDITION In monotonic deformation, the strong anisotropy of the plastic behaviour of TiAl PST crystals was found to depend on the angle c/J. In soft deformation mode at c/J = 45°. shear deformation occurs predominantly on (11 I) planes in the ')'-matrix parallel to the lamellar planes, giving low yield stress and adequate elongation. In contrast, slip and defonnation twins occur on {Ill} planes crossing the lamellar planes in hard deformation at c/J = 0 or 90°, resulting in high yield stress but poor ductility. Figure 2 shows cyclic hardening curves of TiAl PST crystals containing 49·1 at% Al and 50·8 at% AI with c/J = 0 and 45° at different total strain amplitudes. 19 The volume fraction of arphase, the mean lamellar spacing and the mean domain size of Ti-49·1 at% Al and Ti-50·8 at% Al PST crystals are 8 and 2 vol%, 0·75 and 1·13 J..l.m, and Loading axis
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Fig. 1. Stereographic projections of loading axes of TiAI PST crystals: (a) 0 and 6 show loading axes for each domain in the y-phase at c/J = 0 with X = 0 and at c/J = 45°. respectively; (b) loading axes in the arphase at c/J = 0 (0) and 45° (6); (c) schematic drawing of loading axis for changing x-value.
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Fig. 2. Cyclic hardening curves for Ti-49 ·1 at% Al and Ti-50·g at'}'o Al PST crystals deformed at ¢ = 0 with X == 0 and at ¢ = 45° at various stra in amplitudes (.1e). Deformation was performed at room temperature in air.
23 and 64 J,Lm. respectively. If compared at the same total strain. the stress amplitudes are about three times greater at f/> = 0 than at f/> = 45° on the basis of different interactions of moving dislocations at the lamellar boundaries. The cyclic hardening behaviour depended strongly on the orientation and the strain amplitude. At f/> = 0 the stress amplitude increased with the number of cycles and cyclic hardening became noticeable as the strain amplitude rose. In the range of strain amplitudes less than ±O·3% the stress amplitude was saturated after an initial sharp increase giving no failure of specimens until I X 104 cycles. while an increase in the strain amplitude to more than ±O·4% produced greater cyclic hardening and a short fatigue life. In contrast. at f/> = 45° the stress amplitude (which increased with increasing strain amplitude) was saturated in an early stage after slight cyclic hardening and no noticeable hardening was observed even when the specimen broke. Anisotropy in cyclic hardening between specimens deformed at f/> = 0 and 45° occurred similarly in both Ti-49·1 at% Al and Ti-50·g at% Al crystals. but there was a difference in fatigue life. Figure 3 shows the fatigue lives of Ti-49·1 at% Al and Ti 50·8 at% AI PST crystals with cP = 0 and 45° at different strain amplitudes." A decrease in strain amplitude and refinement of lamellae prolonged fatigue life. At large applied total strains. specimens with f/> = 45° showed a more
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Fig. 3. Fatigue life of Ti-49 ·1 at%AI and Ti-50·8 at%AI PST crystals cyclically deformed with ¢ = 0 and 45° at different total strain amplitudes.
longer fatigue life than those with cP =O. while the reverse was true at low strain amplitude. In monotonic deformation a great difference was noticed between the yield stresses of specimens with cP = 0 and 45°. At the same total strain amplitude the specimen with cP = 0 showed a narrow hysteresis loop extending to the stress axis between applied stress and total strain. while at f/> = 45° there was a wide and orthorhombic-shaped loop expanding to the strain axis. If the yield stress increases. the elastic region in the hysteresis loop is extended. resulting in reduction of the plastic strain at the same total strain. The plastic strain given in specimens at cP = 45° for each cycle was larger than that at cP =O. The larger plastic strain at cP = 45° may be responsible for the shorter fatigue life at cP = 45° than at cP = 0 in the range of low strain amplitudes. The x-value also affects the cyclic hardening behaviour and fatigue life because activated slip and twinning systems in each domain vary with the change of this value. Details are given in the next section. 4 ANISOTROPY IN DEFORMATION SUBSTRUCTURE OF CYCLICALLY DEFORMED TiAI PST CRYSTALS
The strong anisotropy of the cyclic behaviour of TiAI PST crystals at cP = 0 and 45° is closely related to the deformation substructure which is driven by the difference in the roles of oiy lamellar and vv domain boundaries; the boundaries act as a barrier to the motion of dislocations and the propagation of deformation twins on {I II} planes parallel or across the lamellar planes.
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At cP = 45° dislocations andlor twins were mainly observed on (I I I) planes parallel to the lamellar planes. The density of dislocations and twins depended strongly on the type of ;,-domain. Focusing on the difference in deformation substructure, the domains in Fig. I(a) can be classified into T(twin)-type and D(dislocation)-type domains. Deformation twins and dislocations were predominantly observed in T-type and D-type domains, respectively, as shown in Fig. 4. Considering the analogous crystallographic relationship and the deformation substructure among these six types of domain, (AM and AT) and (BM, BT, CM and CT) belong to T-type and D-type domains, respectively. In monotonic deformation of Ti-rich TiAI alloys, ~< I I OJ-type ordinary dislocations and deformation twins (which are equivalently activated) are preferentially more operative than <101]type superlattice' dislocations.P:" During cyclic deformation twinning occurs in the initial stage, but it does not directly contribute to further plastic behaviour since it cannot perform the reversible motion in tension and compression. In D-type domains, a high density of dislocations, loops and debris-which were uniformly distributed on the (I I I) planes-were observed as shown in Fig. 4(b). Although some segments of [0 j I] and ~[I I 2] superlattice dislocations were observed, a large number of dislocations and almost all loops and debris were composed of ![I j 0] ordinary dislocations. At cP = 45° dislocations and twins, predominantly activated on the (II I) planes in the ;,-phase parallel to the lamellar planes, were homogeneously distributed on the planes and their motion may be interrupted at the "11"1 domain boundaries, which did not act as effectively as did the lamellar boundaries at cP = O. ~[I j 0] ordinary dislocations are dominant in D-type domains but not in neighbouring different T-type domains
(a)
(b)
Fig. 4. Deformation substructure of Ti-49·1 at% Al PST crystals cyclically deformed at room temperature to 1000 cycles !!t rb = 45°_and Je = ±Q·3%: (a) Ar
because of the low stress factor for (I 1 1)1[1 j 0] slip. Shear displacement induced by (l 1 I)~[l j 0] slip in a domain is parallel to that induced by (l I 1)type dislocations were observed in the arphase with further cyclic deformation andlor at a large strain amplitude of Je = ±O·5%. In T-type (AM and AT) domains a large number of deformation twins could be seen on {I I I} planes crossing the lamellar boundaries, accompanied by a low density of dislocations. In contrast, in V-type (BM, BT, CM and CT) domains, !< I I Ol-type ordinary dislocations were predominantly activated and there were few deformation twins. Since the vv twin and aiy lamellar boundaries are an effective barrier to
Plastic anisotropy and fatigue of TiAI PST crystals
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the motion of dislocations and the propagation of twins, ~< I IOJ-type ordinary dislocations performed the to-and-fro motion in the same domain during cyclic loading. During this motion, they tangled with each other and formed a vein-like structure with high residual stress, resulting in strong cyclic hardening. The increase in number of V-type domains was responsible for the sharp cyclic hardening and short fatigue life. Since twinning is directional and occurs in only one direction, twins can be activated only under tension (positive loading) or compression (negative loading) depending on the type of domain. In the positive half-cycle during cyclic loading, twins can be activated in (BM • Br• eM and CT ) domains, while they can be activated in (AM and AT) domains in the negative half-cycle. In Tvtype (AM and Ar) domains both deformation twins and UI 10] ordinary dislocations could be operative. but the motion was interrupted at the yly domain boundaries since their Burgers vector was not parallel to the lamellar planes. In addition, the Schmid factor for the activation of ~II 10] dislocations in T-type domains was smaller than that for UI 10] dislocations in V-type (BM , Br• eM and c,.) domains. Therefore. twins were predominantly activated in the negative half-cycle accompanied by a small
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amount of hard mode dislocations, and high stress was required to deform the T-type domains. Since deformation twins were induced at an early stage of fatigue and their to-and-fro motion did not occur easily, the density of twins did not increase remarkably during cyclic loading, resulting in a weak contribution to cyclic'hardening. The deformation substructure in each domain varies with the X-value. The combination of V-type and T-type domains can be changed with different X-values. Figure 6 shows schematic pictures of the x-dependence of deformation substructure in cyclically deformed TiAI PST crystals at cP = O. As the number of V-type domains increases, the stress amplitude of PST crystals at the initial stage of fatigue decreases and the cyclic hardening becomes pronounced. The fatigue life is related to the cyclic hardening behaviour, with rapid cyclic hardening shortening the fatigue life. During cyclic loading, plastic deformation is concentrated in V-type domains by the to-and-fro motion of dislocations. When V-type domains harden sufficiently, deformation in T-type domains is accelerated under high applied stress and the deformation concentrates near the twins, which act as nucleation and propagation sites for slip andlor twinning and aid the development of surface steps. The heavy surface steps may play an important role in the initiation of a microcrack. Highly dense dislocations pile up near the lamellar and domain boundaries forming high residual stress which may also have an important function as a trigger for crack initiation. Thus, the increase in number of V-type domains accelerates cyclic hardening and is detrimental to fatigue life.
5 PLASTIC BEHAVIOUR OF TiAI PST CRYSTALS FATIGUED UNDER CONSTANT APPLIED STRESS Since the six types of y-domain and u2-phase in TiAI PST crystals show a significant difference in deformation mode and strength, their plastic behaviour depends strongly on the applied stress under cyclic loading. The effect of the deformation mode of y-domains and urphase on fatigue behaviour can be clearly recognised by a constant loading fatigue test. Figure 7 shows the hysteresis loops between applied stress and total strain of TiAI PST crystals cyclically deformed at room temperature and cP =0 and 45° under a stress-controlled condition. Depending on the applied stress, a great difference
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Fig. 6. Schematic pictures of the X-dependence of deformation substructure in cyclically deformed TiAI PST crystals at r/J (a) X = O. (b) X = 15°. (c) X = 30°. Stress (MPa)
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4 Strain (%)
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Fig. 7. Hysteresis loop of Ti-49·1 at% Al PST crystals cyclically deformed at several stress-controlled amplitudes: (a) r/J N = 20 cycles; (b) r/J = 45°. N = 20 cycles.
in shape of the loop was found between specimens at cf> = 0 and 45°. After the first cycle the plastic strain remained at zero stress level in a hysteresis loop, and then increased monotonically with increasing applied stress. However, an anomalous change in the loop at cf> = 0 can be seen in the course of further deformation, e.g. at N = 20 cycles, while there is no anomalous dependence of applied stress on the hysteresis loop at cf> = 45° even at further cycles. The plastic strain energy, which is determined by the area of the hysteresis loop, is closely related to the energy required for twinning andlor the to-and-fro motion of dislocations during one cycle. Figure 8 shows the variation in plastic strain energy with applied stress at N = 20 cycles. At cf> = 45° the plastic strain energy increases as the applied 0 the energy increases stress rises, while at cf>
=
= O.
sharply with increasing stress and then shows a sudden drop around 440 MPa followed by a rapid increase. A peak of plastic strain energy is therefore observed around 430 MPa. In both cases this energy decreases with increasing number of cycles, since dislocations are piled up, tangled and interact with each other resulting in adequate cyclic hardening. The anomalous change of plastic strain energy was observed in the course of cyclic deformation although the value of the energy decreased with increasing number of cycles. The relation between the plastic strain energy and applied stress at cf> = 0 in Fig. 8 can be divided into three regions. In region I where stress amplitudes below 430 MPa are applied, ordinary dislocations are operative in V-type domains and twinning occurs in T-type domains at an early stage of deformation. The increase in applied stress requires
Plastic anisotropy and fatigue of TiAI PST crystals
120
140
160
180
200
Stress amplitude (MPa)
Fig. 8. Plastic strain energy-stress amplitude curves of Ti-49·1 at'Y., Al PST crystals cyclicalIy deformed at stress-controlIed amplitude with c/J = 0 and 45°.
the motion of a large number of dislocations, resulting in an increase in plastic strain energy. During the to-and-fro motion of dislocations, a vein-like structure is formed in V-type domains and cyclic hardening occurs preferentially in these domains with increasing applied stress and number of cycles. However, twinning in T-type domains does not effectively contribute to the cyclic hardening, although extrusions develop on the specimen surface during cyclic loading at stress near the peak in Fig. 8. In the arphase no significant feature is observed on the specimen surface or in the substructure. In region II where the plastic strain energy decreases sharply, twinning begins to occur in V-type domains accompanied by the motion of ordinary dislocations, while twinning is predominantly operative in T-type domains similar to region I, although no significant surface extrusions are observed. The to-and-fro motion of ordinary dislocations, which form a vein-like structure with highly tangled dislocations at a further stage, is interrupted by deformation twins in T-type domains, and these domains are hardened. The hardening of V-type domains-which is accelerated by twinning~orresponds to a sudden drop of plastic strain energy in this region. Prism slip is activated in the a 2-phase since sufficient stress and strain are applied. In region III, twins are observed in both T- and V-type domains since there is enough stress applied to activate them. In Tvtype domains heavy extrusions develop on the specimen surface and slips aiding the development of these extrusions
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are accommodated near twin boundaries. In V-type domains, ordinary dislocations which produce numerous loops and form veins are operative, accompanied by twinning. Although cyclic hardening becomes remarkable with an increase in applied stress that accelerates the formation of veins, dislocations can glide against the residual stress from lattice defects under high applied stress, resulting in an increase in plastic strain energy. Slip traces on the prism plane in the a2-phase also became remarkable. In contrast, plastic strain energy increases monotonically at q, = 45° as the applied stress rises. Twinning and slips occur on (1 1 1) planes parallel to the lamellar planes in T-type and Dtype domains. There is no significant stress dependence in slip markings and deformation substructure, although the frequency of slip traces and the density of twins and dislocations increase with increasing stress. Even at high applied stress no slip markings were observed in the arphase, which does not directly play a key role in cyclic deformation. In general, fatigue life decreases with increasing applied stress: in fact, the fatigue life of TiAI PST crystals cyclically deformed at q, = 45° decreases monotonically as the applied stress rises. However, the fatigue life and fracture behaviour of TiAl PST crystals at q, = 0 are closely related to the deformation mode of T-type and V-type domains. At the lower stress level in region I, TiAI PST crystals showed a long fatigue life and specimens did not fracture even after I X 105 cycles. Anomalous damage appeared for fatigue lives between regions I and II, however. At the peak in the plastic strain energy-applied stress curves, fatigue life was suddenly shortened. In region II on the border with region I, fatigue life was long and specimens did not fail at I X 105 cycles. Fatigue life then decreased with increasing applied stress as is true for other materials. At the peak of the curve, surface extrusions in T-type domains developed and the stress concentration near the extrusions was responsible for the initiation of microcracking, resulting in specimen failure. In region II plastic deformation is uniformly distributed and divided in T-type and V-type domains. Formation of the surface extrusions is suppressed and fatigue life is again improved. In region III the increase in applied stress hardens V-type domains and highly localized stress initiates cracking near a2/Y boundaries and in the a2-phase. In the previous section it was described that twinning is directional and occurs in only one
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direction; in aT-type domain twins can only be activated in tension or in compression depending on the loading axis. Even if twins were activated in the positive half-cycle during cyclic loading, twinning would not occur in the negative halfcycle. When cyclic load is applied in tension/tension mode and in tension/compression, the frequency of activation of twins can be changed between specimens with the same loading axis. At the same applied stress, the number of cycles to fracture of a specimen loaded in the tension/tension mode is larger than that of a specimen loaded in tension/ compression mode. For example, at a = 500 MPa a specimen of Ti-49·1 at% AI PST crystal deformed at cP = 0 under cyclic loading in the tension/tension mode fractured at N = 746 cycles, while in the tension/compression mode it fractured at N = 37 cycles. The number of deformation twins increased more in the tension/compression mode than in the tension/tension mode. There was also a significant difference in slip markings on the specimen surface between the two cyclic loading modes as shown in Fig. 9. In the tension/compression mode the surface extrusions developed in a peculiar domain and fatigue life was shortened. This result strongly suggests that deformation twins are detrimental to microcrack initiation and fatigue life. 6 EFFECT OF ADDITIONAL ELEMENTS ON CYCLIC BEHAVIOUR AND FATIGUE LIFE In monotonic deformation, the addition of a third element-such as Nb, V, Mn, Cr or Mo-is known to be effective in improving strength and ductility.P-" Addition of V, which may decrease the stacking fault energy, easily activates deformation twins in the ,,-phase, resulting in improved ductility at room temperature.P:" The addition of Nb, in contrast, increases the stacking fault energy
but may suppress the anisotropy of plastic behaviour due to the activation of pyramidal slip in the €X2-phase. 30 The effect on cyclic behaviour and fatigue life of adding V and Nb, which respectively may induce a decrease and an increase in the stacking fault energy of 'Y"phase TiAI, has also been investigated." Figure 10 shows the variation in stress amplitude with cumulative plastic strain for binary and ternary TiAI PST crystals cyclically deformed at cP = 45° at de = ±D·5 and ±0·3%. Adding a third element has a significant effect on the fatigue life. Nb is effective but V is harmful to fatigue life at cP = 45°. A similar tendency was observed at cP = 0 but was not significant in the case of Nb addition. In V addition, twinning is more activated and T-type domains are adequately hardened at an early stage of cyclic loading. With further loading, plastic deformation is concentrated in D-type domains. Highly dense dislocations in these domains produced during their to-and-fro motion cannot easily pass through domain boundaries and enter the neighbouring T-type domains because the T-type domains are hardened by the high density of deformation twins. In addition, frequent twinning may produce the surface steps and extrusions which act as initiation sites for microcracking. Once a microcrack is initiated at cP = 45°, it easily develops into a macrocrack and fracture occurs. Frequent twinning induced by V addition aids in starting a microcrack, resulting in short fatigue life. In contrast, since Nb addition suppresses twinning, plastic strain during cyclic loading does not concentrate in 250
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Fig. 10. Variation in stress amplitude with cumulative strain of binary and ternary TiAI PST crystals cyclically deformed at q, 45°.
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Plastic anisotropy and fatigue of TiAI PST crystals
specific domains. The highly tangled dislocation structure is gradually formed, accompanied by moderate cyclic hardening. In D-type domains, dense locally tangled dislocations are formed since Nb addition decreases the separation distance between Schockley partials dissociated from an ordinary dislocation and the dislocation can easily cross-slip. Nb addition improves the fatigue life at cP =45°. At cP = 0, addition of Nb improves the fatigue life but it does not work effectively as it does at cP = 45°. The fatigue life is shortened by V addition. In ternary TiAI PST crystals containing Nb the average domain size is similar to that of binary ones, but the average lamellar spacing is much larger. The propagation of cracking, which is effectively interrupted by lamellar boundaries, is important in fatigue life at c/J = O. 7 TEMPERATURE DEPENDENCE OF CYCLIC DEFORMATION AND FATIGUELIFE The cyclic behaviour and fatigue life of TiAI PST crystals show a strong temperature dependence. Figure 11 shows the relation between stress amplitude (S) and number of cycles to failure (N) of TiAI PST crystals cyclically deformed at different temperatures. The S-N curves shifted to higher stress level with decreasing temperature, although there was no significant difference between specimens at 25°C and 400°C with cP = O. At cP = 45° the S-N relation showed a straight line at all temperatures and the fatigue life became shorter as the stress amplitude increased. At cP = 0 a similar tendency in the S-N relation was found.
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In monotonic deformation, the mobility of ordinary and superlattice dislocations decreases and yield stress increases with decreasing temperature. Since the cross-slip event and/or climb motion of moving dislocation are thermally assisted, dislocations show the tendency to glide on the same slip plane at low temperature. As previously described, the vein-like structure composed of highly dense dislocations and low density regions in V-type domains is responsible for the cyclic hardening. Some segments of moving dislocations that perform the climb motion and/or the cross-slip event produce dislocation, loops and debris during cyclic loading. During the to-and-fro motion of dislocations, highly tangled dislocations form a vein-like structure whose formation is assisted by the formation of loops and debris. Since the crossslip of ordinary dislocations hardly occurs at -196°C during their to-and-fro motion in V-type domains, formation of the vein-like structure is suppressed, resulting in weak cyclic hardening. Accommodation slip near deformation twins in 'f-type domains, which causes development of surface extrusions, is also suppressed at low temperature. In contrast, cross-slip and/or climb motion occurred easily at room temperature during cyclic loading. The mobility of dislocations increased rapidly with increasing temperature. Deformation is controlled by the motion of ordinary dislocations and twinning is not dominant even at high applied stress in V-type domains. However, the surface extrusions were still significant in T-type domains since accommodation slip near deformation twins occurred easily. At high temperatures, for example at 700°C, the mobility and climb motion of ordinary dislocations increased rapidly and the formation of extrusions was suppressed. The rearrangement of dislocations during cyclic loading softens the crystal, resulting in short fatigue life.
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The initiation and propagation of cracking and the fracture of cyclically deformed TiAI PST crystals depend strongly on strain amplitude, stress amplitude, orientation of loading axis and deformation temperature. At c/> = 45°, twinning and a long movement of dislocations that were concentrated in localized regions of the y-phase occurred predominantly on the (l 1 I) plane in y-domains parallel to the lamellar boundaries and aided the development of heavy extrusions. On a macroscopic scale fracture occurs parallel to the lamellar
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boundaries, independent of deformation temperature, microstructure and the applied stress. Highly concentrated stress at extrusions-which are closely related to twinning-has an important function in crack initiation. At low temperatures twinning stress increases and formation of the vein-like structure that causes strong cyclic hardening is hindered due to the suppression of cross-slip of moving dislocations during cyclic loading. The formation of extrusions, which act as a trigger to initiate a microcrack, is suppressed. The fractograph in Fig. 12(a) shows a flat surface in a cleavage mode but was not accompanied by a zigzag motion, suggesting that y/y domain boundaries do not play an important role in suppressing the crack propagation. Refinement of lamellae was effective in improving the fatigue life at all tested temperatures. Fine lamellae and small ordered domains lead to homogeneous deformation which assists in suppressing the formation of extrusions and initiation of a microcrack, and prolongs fatigue life. However, there is not much difference in the fracture morphology between specimens with fine and coarse lamellae. At cP = 0, slip markings on the {I I I} plane in the y-phase crossing the lamellar boundaries were observed but they were unable to enter the a2phase at low strain amplitude. At high strain amplitude, microcracks were often seen on the (000 I) plane in a2-plates but they did not have a direct connection with specimen failure. According to our recent result on the orientation dependence of fracture of Ti3AI single crystals, not only the basal planes but also the {I 0 I 2} pyramidal planes were susceptible to brittle fracture." In TiAI PST crystals, a crack propagating on the (000 1) plane in an arplate suddenly changed course across the y-phase and proceeded again in a neighbouring arplate, resulting in a zigzag crack path between neighbouring a2-plates. The
brittle nature of fracture on {I 0 12} planes in the arphase may act as a trigger to change the course of a crack path and be responsible for this zigzag motion. Figure 12(b) shows a scanning electron micrograph of the fracture surface of a specimen cyclically deformed at cP =O. There were many terraces and projections on y-domains and a2-plates, suggesting that the lamellar boundaries may act as a barrier to suppress the further propagation of cracks, and they do this effectively in fine lamellae. In T-type domains, twins accompanied by dislocations were observed but multiplication of new twins does not occur easily because twins do not perform the reverse motion during cyclic loading. When V-type domains are sufficiently hardened by the formation of vein-like structure during the to-and-fro motion of dislocations, T-type domains begin to deform preferentially with further cyclic deformation: this deformation concentrates near the twins, which act as nucleation and propagation sites for slip and/or twinning and aid the development of surface steps. The large residual stress field associated with inhomogeneous distribution of dislocations and twins acts as a trigger to nucleate a microcrack. The extrusions also act as a nucleation site for microcracks. At room temperatures and below, an anomalous peak appears in the plastic strain energy-applied stress curves. At small applied stress below the peak, twinning occurs in T-type domains but V-type domains are deformed by the motion of dislocations. In contrast, at large applied stress above the peak, twinning is operative even in V-type domains. Interaction between twins and highly dense dislocations forms a concentrated residual stress field which acts as a trigger for microcrack initiation. Therefore, frequent microcrack nucleation at high applied stress leads to a short fatigue life. At high temperatures the anomalous peak disappears in the plastic strain energy-applied stress curves because twinning is suppressed in V-type domains even at higher applied stress. Fracture morphology, therefore, does not show a strong dependence on applied stress.
9 CONCLUDING REMARKS (a)
(b)
Fig. 12. Scanning electron fractographs of Ti-49·) at°/r, Al PST crystals cyclically deformed at room temperature: (a) t/J =45°, ae = to·)'X" N = 3300 cycles; (b) t/J =0, ae =to·3%. N =2500 cycles.
Large anisotropy of strength, ductility, cyclic hardening and fatigue life is observed in TiAI PST crystals. Variation in deformation mode, such as twinning and the to-and-fro motion of ordinary
Plastic anisotropy and fatigue of TiAI PST crystals
dislocations, is responsible for this. The extrusions connected with twinning are important in the initiation of microcracks. In monotonic deformation, addition of V which may decrease the stacking fault energy activates twinning in Ti-rich TiAI alloys at room temperature, resulting in improved ductility at room temperature. However, twinning is harmful to the fatigue life because it accelerates the formation of extrusions on the specimen surface and leads to highly concentrated residual stress near lamellar and domain boundaries. From the viewpoint of twinning, addition of Nb, which may increase the stacking fault energy, suppresses the twinning but does not effectively improve the fatigue life at cP = O. Deformation of arphase does not directly contribute to the cyclic deformation behaviour. but the accommodation of slip in a2- and 'Y'"phases is required to relax the highly concentrated residual stress at the a21y lamellar boundaries and suppress microcrack initiation. Since a2/y lamellar boundaries act as an effective barrier to stop microcrack propagation, the fatigue life must be controlled by the initiation and propagation of the microcrack. Polycrystalline alloys with lamellar structure show better fatigue life than those with duplex structure. To obtain a good combination of strength, ductility and fatigue life, refinement of lamellae is the most effective method. Even at high temperatures the lamellar boundaries still act as an effective barrier for microcrack propagation and good fatigue properties are maintained. The choice of alloying elements and thermo-mechanical treatment should be considered to obtain fine lamellae.
ACKNOWLEDGEMENT Y. Umakohi would like to thank the Ministry of Education. Science. Sports and Culture of Japan for a Grant-in-Aid for Scientific Research on the Priority Area 'Intermetallic Compounds as New High Temperature Structure Materials'.
REFERENCES I. Lipsitt, H. A., Shechtman. D. & Schfrik, R. E., Metall. Trans.• 8 (1956) 1348. 2. Yamaguchi. M. & Inui, H., in Structural lntermetallics, eds R. Darolia et al.. TMS, Warrendale. 1993. p. 127. 3. Kim. Y.-W. & Dimiduk, D. M.• J. Metals, 46 (1994) 30.
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4. Umakoshi, Y., Nakano, T. & Yasuda. H. Y., in Mechanical Properties and Phase Transformation of Multi-Phase Intermetallic Alloys. eds A. F. Giami et al.. TMS, Warrendale, 1995, p. lSI. 5. Fujiwara, T., Nakamura. A., Hosomi, M.• Nishitani, S. R., Shirai, T. & Yamaguchi, M.• Phil. Mag.• A61 (1990) 591. 6. Inui, H., Oh, M. H., Nakamura, A. & Yamaguchi, M., Acta Metall. Mater., 40 (1992) 3095. 7. Umakoshi, Y. & Nakano, T., Acta Metall. Mater., 41 (1993) 1155. 8. Umakoshi. Y. & Nakano, T., Mater. Sci. Eng. , AIS2 (1992) 81. 9. Umakoshi, Y. & Nakano, T.• ISIJ Int., 32 (1992) 1339. 10. Sastry, S. M. L. & Lipsitt, H. A.• Metall. Trans, 8A (1977)299. II. Dowling, W. E. Jr, Donlon, W. T. & Allison. J. E., in High- Temperature Ordered lntermetallic Alloys IV, eds L. A. Johnson et al.. MRS Symp. Proc., 213 (1991) 561. 12. Soboyejo, W.O., Deffeyes, J. E. & Aswath, P. B., Mater. Sci. Eng., A138(1991) 95. 13. James. A. W. & Bowen. P., Mater. Sci. Eng.• AlS3 (1992) 486. 14. Davidson, D. L. & Campbell. J. B., Metall. Trans.,24A (1993) 1555. 15. Kumpfert, J., Kim, Y.-W. & Dimiduk, D. M.• Mater. Sci. Eng.• A1921193 (1995) 465. 16. Sloloff, N. S.• Fuchs, G. E.• Kuruvilla, A. K. & Choe, S. 1., in High-Temperature Ordered Intermetallic Alloys Il, ed. N. S. Stoloff, MRS Symp. Proc., 81 (1987) 247. 17. Stoloff, N. S.• Smith. T. R. & Castagna, A.. in HtghTemperature Ordered Intermetallic Alloys V, ed. 1. Baker, MRS Symp. Proc., 288 (1993) 59. 18. Yauda, H . Y.• Nakano, T. & Umakoshi, Y., Phil. Mag., A71 (1995) 127. 19. Umakoshi, Y., Yasuda, H. Y. & Nakano, T., Mater. Sci. Eng., A194 (1995) 43. 20. Yasuda. H. Y., Nakano, T. & Umakoshi, Y., Phil. Mag.• (1996) in press. 21. Yasuda. H. Y.• Nakano, T. & Umakoshi, Y., Phil. Mag., (1996) in press. 22. Nakano, T. & Umakoshi, Y., Intermetailics, 2 (1994) 185. 23. Schechtman. D., Blackburn. M. J. & Lipsitt, H. A.• Metall. Trans.,5 (1974) 1373. 24. Court, S. A.• Vasudevan, V. K. & Fraser. H. L.. Phil. Mag.• A61 (1990) 141. 25. Inui, H., Nakamura, A., Oh, M. H. & Yamaguchi, M.• Phil. Mag., A66 (1992) 557. 26. Hung, S. C; in Structural Intermetallics, ed. R. Darolia et al.• TMS. Warrendale, 1993, p. 299. 27. Hashimoto. K. & Kimura, M" in Structural Intermetallics, ed. R. Darolia et al... TMS, Warrendale, 1993, p.309. 28. Whang. S. H.• Hahn, Y. D., u, Z. X. & u Z. c, in JIMIS-6 : Intermetallic Compounds-Structure and Mechanical Properties, ed, O. Izumi, lap. Inst. Metals, Sendai, 1991. p. 763. 29. Yao, K.-F.• Inui, H., Kishida, K. & Yamaguchi, M.• Acta Metall. Mater.• 43 (1995) 1075. 30. Nakano, T., Yokoyama, A. & Umakoshi, Y., Scripta Metall. Mater., 27 (1992) 1253. 31. Umakoshi, Y., Yasuda, H. Y.• Nakano. T. & Nakazawa, J.• in Gamma Titanium Aluminides, eds Y.-M. Kim, R. Wagner & M. Yamaguchi. TMS. Warrendale, 1995, p.697. 32. Umakoshi, Y., Nakano, T. & Ogawa, B., Scripta Metall. A!ater., (1996) in press.
lmermetullics 4 (1996) S65-S75 " 1996 Elsevier Science Limited Printedin Great Britain. All rightsreserved 0966-97951961$15,00
ELSEVIER
Plastic anisotropy and fatigue of TiAl PST crystals: a review Y. Umakoshi, H. Y. Yasuda & T. Nakano Department of Materials Science and Engineering, Faculty of Engineering, Osaka University, 2-/ Yamada-Oka, Suita, Osaka 565. Japan (Received I November 1995: accepted 27 December 1995)
The cyclic deformation and fracture behaviour of TiAI PST crystals are reviewed. Anisotropy in cyclic hardening. fracture and fatigue life of the crystals deformed at 4> 0 and 45°. where 4> is the angle between the loading axis and lamellar planes. are discussed focusing on the different deformation substructures in '}'-domains. The effects of deformation temperature. addition of third elements and lamel1ar structure on fatigue life are also described. Copyright © 1996 Elsevier Science Ltd
=
Key words: A. titanium aluminides, based on TiAI, B. anisotropy. fatigue resistance and crack growth.
1 INTRODUCTION
lamellar spacing, domain size and the structure of lamellar and domain boundaries: the slip mode of the a2-phase was responsible for the
The discovery of a combination of high strength and good ductility in two-phase Ti-rich TiAI alloys has accelerated attempts to develop the alloys as potential high-temperature structural materials for use in airframes, car bodies, automobile engine valves and turbo rotors.':' Upon solidification of Ti-rich TiAI, a lamellar structure composed of y-matrix and a small amount of thin arplates forms during the a-to y-phase transformation. Since mechanical properties such as strength, ductility, fracture toughness and fatigue are known to be very sensitive to the orientation and microstructure of lamellae, knowledge of the effect of lamellar structure on plastic behaviour of TiAI is necessary to improve these properties. The successful achievement of polysynthetically twinned (PST) crystals, in which a single set of lamellae is unidirectionally aligned with no grain boundaries, has provided clear information on the effect of lamellar structure."? In monotonic deformation, the y-phase in equilibrium with the a2-phase in two-phase TiAI showed good ductility with low yield stress, but lamellar boundaries containing a2-phase-which were often accompanied by a loss of ductility-were required for strengthening. The anisotropy in strength, ductility and fracture behaviour was closely related to the angle between the loading axis and lamellar planes,
anisotropy.r" Since industrial materials are often subjected to repeatedly applied stress and strain, knowledge of the plastic behaviour under cyclic loading conditions is required for practical use in addition to monotonic deformation. The fatigue life, crack growth and fracture behaviour of polycrystalline TiAI alloys have been examined and good fatigue properties have been demonstrated at room and high temperatures.P:" The fatigue life of metals and alloys is controlled by the initiation of a microcrack and its propagation accompanied by an intergranular and/or transgranular crack path. Crack initiation in intermetallic compounds is known to play an important role in fatigue life because of their low ductility.P''" However, the lamellar boundaries of TiAI alloys often act as an effective barrier to the propagation of a crack. Similar to that in monotonic deformation, remarkable progress has been made in understanding the effect of lamellar structure on the plastic behaviour of TiAI under cyclic loading using TiAI PST crystals.v"?' This article reviews recent advances in our understanding of anisotropy and the mechanism of plastic and fracture behaviour of TiAI alloys under cyclic loading, concentrating on recent results on TiAI PST crystals. S65
Y. Umakoshi et al.
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2 LOADING AXES AND SCHMID FACTOR OF y-DOMAINS The plastic behaviour of TiAI PST crystals is well known to exhibit a strong dependence on orientation c/J, where c/J is the angle between the loading axis and lamellar planes. Deformation in easy mode occurs only on (I I I) planes in the ')'-matrix parallel to the lamellar boundaries, and good ductility with low yield stress is shown at c/J = 45°.6.22 For hard orientations at c/J =0 and/or 90°, lamellar boundaries act as an effective barrier to the motion of dislocations, resulting in high yield stress but a loss of ductility. The angle c/J was chosen to be 0 and 45° to examine the effect of hard and easy mode on the cyclic deformation and fatigue properties of TiAI PST crystals. At c/J = 90° specimens were broken within a few cycles under push/pull loading because no slip systems other than pyramidal slip were activated in the a2-plates and the a2-plates hardly deformed. Figure 1 shows the stress axes of the arphase and each domain in the y-matrix of PST crystals c/J = 0 and 45°. The loading axis of the a2-phase occupies one position for each specimen, but since there are six types of ')'-domain in consideration of the matrix and twinning relation, the corresponding stress axes in each domain of the ')'-matrix are plotted as AM-CT in Fig. I(a), where the suffixes M and T represent the matrix and twinning relationship to each other. To examine the effect of deformation modes of twinning and slip in each domain, the loading axis is rotated on a (l I l) lamellar plane at c/J = 0 as shown in Fig. I(c). When the rotation angle (X) between the stress axis and the [I I 2] plane in AM domains changes from 0 to 30°, the combination of equivalent
domains on the basis of stress factors for the dominant slip and twinning systems is alternated. The two domains in the following parentheses are equivalent: (AM' AT). (BM• BT) and (CM, CT) at X =0; (AM' AT), (BM, CT) and (BT, CM) at X = 15°; (AM' BT), (BM, ~) and (CM, AT) at X = 30°. In particular, high stress factors for twinning exist in domains of (B M • ~), (CM, CT) at X =0, (B M , CT) at X = 15° and (BM , CT) at X = 30°. Therefore, as X varies at c/J = 0, the number of domains in which twinning would be activated can be changed to examine the effect of deformation twins on cyclic deformation and fatigue properties.i-" When x-value is not indicated in this paper, X =0 is chosen. 3 CYCLIC HARDENING AND FATIGUE LIFE UNDER A TOTAL STRAIN CONTROLLED CONDITION In monotonic deformation, the strong anisotropy of the plastic behaviour of TiAl PST crystals was found to depend on the angle c/J. In soft deformation mode at c/J = 45°. shear deformation occurs predominantly on (11 I) planes in the ')'-matrix parallel to the lamellar planes, giving low yield stress and adequate elongation. In contrast, slip and defonnation twins occur on {Ill} planes crossing the lamellar planes in hard deformation at c/J = 0 or 90°, resulting in high yield stress but poor ductility. Figure 2 shows cyclic hardening curves of TiAl PST crystals containing 49·1 at% Al and 50·8 at% AI with c/J = 0 and 45° at different total strain amplitudes. 19 The volume fraction of arphase, the mean lamellar spacing and the mean domain size of Ti-49·1 at% Al and Ti-50·8 at% Al PST crystals are 8 and 2 vol%, 0·75 and 1·13 J..l.m, and Loading axis
~(H21
I
XI
-:
1010
011
b
Cr a
b
/~A
A~
001
101
(a)
100
0001
(b)
2110
(c)
Fig. 1. Stereographic projections of loading axes of TiAI PST crystals: (a) 0 and 6 show loading axes for each domain in the y-phase at c/J = 0 with X = 0 and at c/J = 45°. respectively; (b) loading axes in the arphase at c/J = 0 (0) and 45° (6); (c) schematic drawing of loading axis for changing x-value.
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Plastic anisotropy and fatigue of TiAI PST crystals
o ..
x~e =:to. 7%
~£ =:to.5%
600
~£=:t:0.4%
~500
~e=:t0.3%
6
....
~
0
10
~a
_ •••••• ~e=:t0.2%
~e=:t:0.3%
a.
."
~ 300
.....x
.
III III
• TI-SO.8al%AI, .-0 ....5·
"
~o
I as
~ 400
.-0 ....5·
~ o~~. TI-SO.8al%Al,
E
.~
~9.1at%A1,
~.l""Al,
A
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.~
ell
c75 200
=-:-.-.--100
-_ ..--_ ....
~£=:t:0.3%
4>=45° Ti-49.1 at%AJ --- ••• Ti-50.8at%AI 3
10 1 102 10 Number of cycles
Fig. 2. Cyclic hardening curves for Ti-49 ·1 at% Al and Ti-50·g at'}'o Al PST crystals deformed at ¢ = 0 with X == 0 and at ¢ = 45° at various stra in amplitudes (.1e). Deformation was performed at room temperature in air.
23 and 64 J,Lm. respectively. If compared at the same total strain. the stress amplitudes are about three times greater at f/> = 0 than at f/> = 45° on the basis of different interactions of moving dislocations at the lamellar boundaries. The cyclic hardening behaviour depended strongly on the orientation and the strain amplitude. At f/> = 0 the stress amplitude increased with the number of cycles and cyclic hardening became noticeable as the strain amplitude rose. In the range of strain amplitudes less than ±O·3% the stress amplitude was saturated after an initial sharp increase giving no failure of specimens until I X 104 cycles. while an increase in the strain amplitude to more than ±O·4% produced greater cyclic hardening and a short fatigue life. In contrast. at f/> = 45° the stress amplitude (which increased with increasing strain amplitude) was saturated in an early stage after slight cyclic hardening and no noticeable hardening was observed even when the specimen broke. Anisotropy in cyclic hardening between specimens deformed at f/> = 0 and 45° occurred similarly in both Ti-49·1 at% Al and Ti-50·g at% Al crystals. but there was a difference in fatigue life. Figure 3 shows the fatigue lives of Ti-49·1 at% Al and Ti 50·8 at% AI PST crystals with cP = 0 and 45° at different strain amplitudes." A decrease in strain amplitude and refinement of lamellae prolonged fatigue life. At large applied total strains. specimens with f/> = 45° showed a more
10-' 10,
102 103 Cycles to failure
10·
Fig. 3. Fatigue life of Ti-49 ·1 at%AI and Ti-50·8 at%AI PST crystals cyclically deformed with ¢ = 0 and 45° at different total strain amplitudes.
longer fatigue life than those with cP =O. while the reverse was true at low strain amplitude. In monotonic deformation a great difference was noticed between the yield stresses of specimens with cP = 0 and 45°. At the same total strain amplitude the specimen with cP = 0 showed a narrow hysteresis loop extending to the stress axis between applied stress and total strain. while at f/> = 45° there was a wide and orthorhombic-shaped loop expanding to the strain axis. If the yield stress increases. the elastic region in the hysteresis loop is extended. resulting in reduction of the plastic strain at the same total strain. The plastic strain given in specimens at cP = 45° for each cycle was larger than that at cP =O. The larger plastic strain at cP = 45° may be responsible for the shorter fatigue life at cP = 45° than at cP = 0 in the range of low strain amplitudes. The x-value also affects the cyclic hardening behaviour and fatigue life because activated slip and twinning systems in each domain vary with the change of this value. Details are given in the next section. 4 ANISOTROPY IN DEFORMATION SUBSTRUCTURE OF CYCLICALLY DEFORMED TiAI PST CRYSTALS
The strong anisotropy of the cyclic behaviour of TiAI PST crystals at cP = 0 and 45° is closely related to the deformation substructure which is driven by the difference in the roles of oiy lamellar and vv domain boundaries; the boundaries act as a barrier to the motion of dislocations and the propagation of deformation twins on {I II} planes parallel or across the lamellar planes.
568
Y. Umakoshi et al.
At cP = 45° dislocations andlor twins were mainly observed on (I I I) planes parallel to the lamellar planes. The density of dislocations and twins depended strongly on the type of ;,-domain. Focusing on the difference in deformation substructure, the domains in Fig. I(a) can be classified into T(twin)-type and D(dislocation)-type domains. Deformation twins and dislocations were predominantly observed in T-type and D-type domains, respectively, as shown in Fig. 4. Considering the analogous crystallographic relationship and the deformation substructure among these six types of domain, (AM and AT) and (BM, BT, CM and CT) belong to T-type and D-type domains, respectively. In monotonic deformation of Ti-rich TiAI alloys, ~< I I OJ-type ordinary dislocations and deformation twins (which are equivalently activated) are preferentially more operative than <101]type superlattice' dislocations.P:" During cyclic deformation twinning occurs in the initial stage, but it does not directly contribute to further plastic behaviour since it cannot perform the reversible motion in tension and compression. In D-type domains, a high density of dislocations, loops and debris-which were uniformly distributed on the (I I I) planes-were observed as shown in Fig. 4(b). Although some segments of [0 j I] and ~[I I 2] superlattice dislocations were observed, a large number of dislocations and almost all loops and debris were composed of ![I j 0] ordinary dislocations. At cP = 45° dislocations and twins, predominantly activated on the (II I) planes in the ;,-phase parallel to the lamellar planes, were homogeneously distributed on the planes and their motion may be interrupted at the "11"1 domain boundaries, which did not act as effectively as did the lamellar boundaries at cP = O. ~[I j 0] ordinary dislocations are dominant in D-type domains but not in neighbouring different T-type domains
(a)
(b)
Fig. 4. Deformation substructure of Ti-49·1 at% Al PST crystals cyclically deformed at room temperature to 1000 cycles !!t rb = 45°_and Je = ±Q·3%: (a) Ar
because of the low stress factor for (I 1 1)1[1 j 0] slip. Shear displacement induced by (l 1 I)~[l j 0] slip in a domain is parallel to that induced by (l I 1)
Plastic anisotropy and fatigue of TiAI PST crystals
l
l
BM C,
l
B~,
l
l
I
I
l
l
I
CT BM AT A I AI,4 CT AI,4
Fig. 5. Deformation substructure of Ti--49·! at% AI PST crystal deformed to 1000 cycles at f/J = 0 and .1£ = iO ·]'Yf•.
the motion of dislocations and the propagation of twins, ~< I IOJ-type ordinary dislocations performed the to-and-fro motion in the same domain during cyclic loading. During this motion, they tangled with each other and formed a vein-like structure with high residual stress, resulting in strong cyclic hardening. The increase in number of V-type domains was responsible for the sharp cyclic hardening and short fatigue life. Since twinning is directional and occurs in only one direction, twins can be activated only under tension (positive loading) or compression (negative loading) depending on the type of domain. In the positive half-cycle during cyclic loading, twins can be activated in (BM • Br• eM and CT ) domains, while they can be activated in (AM and AT) domains in the negative half-cycle. In Tvtype (AM and Ar) domains both deformation twins and UI 10] ordinary dislocations could be operative. but the motion was interrupted at the yly domain boundaries since their Burgers vector was not parallel to the lamellar planes. In addition, the Schmid factor for the activation of ~II 10] dislocations in T-type domains was smaller than that for UI 10] dislocations in V-type (BM , Br• eM and c,.) domains. Therefore. twins were predominantly activated in the negative half-cycle accompanied by a small
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amount of hard mode dislocations, and high stress was required to deform the T-type domains. Since deformation twins were induced at an early stage of fatigue and their to-and-fro motion did not occur easily, the density of twins did not increase remarkably during cyclic loading, resulting in a weak contribution to cyclic'hardening. The deformation substructure in each domain varies with the X-value. The combination of V-type and T-type domains can be changed with different X-values. Figure 6 shows schematic pictures of the x-dependence of deformation substructure in cyclically deformed TiAI PST crystals at cP = O. As the number of V-type domains increases, the stress amplitude of PST crystals at the initial stage of fatigue decreases and the cyclic hardening becomes pronounced. The fatigue life is related to the cyclic hardening behaviour, with rapid cyclic hardening shortening the fatigue life. During cyclic loading, plastic deformation is concentrated in V-type domains by the to-and-fro motion of dislocations. When V-type domains harden sufficiently, deformation in T-type domains is accelerated under high applied stress and the deformation concentrates near the twins, which act as nucleation and propagation sites for slip andlor twinning and aid the development of surface steps. The heavy surface steps may play an important role in the initiation of a microcrack. Highly dense dislocations pile up near the lamellar and domain boundaries forming high residual stress which may also have an important function as a trigger for crack initiation. Thus, the increase in number of V-type domains accelerates cyclic hardening and is detrimental to fatigue life.
5 PLASTIC BEHAVIOUR OF TiAI PST CRYSTALS FATIGUED UNDER CONSTANT APPLIED STRESS Since the six types of y-domain and u2-phase in TiAI PST crystals show a significant difference in deformation mode and strength, their plastic behaviour depends strongly on the applied stress under cyclic loading. The effect of the deformation mode of y-domains and urphase on fatigue behaviour can be clearly recognised by a constant loading fatigue test. Figure 7 shows the hysteresis loops between applied stress and total strain of TiAI PST crystals cyclically deformed at room temperature and cP =0 and 45° under a stress-controlled condition. Depending on the applied stress, a great difference
Y. Umakoshi et al.
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(c)
(b)
~
deformation twin
I~I
tangle of ordinary dislocations
Fig. 6. Schematic pictures of the X-dependence of deformation substructure in cyclically deformed TiAI PST crystals at r/J (a) X = O. (b) X = 15°. (c) X = 30°. Stress (MPa)
490 MPa
= O.
Stress (MPa)
200
400
190 MPa
4 Strain (%)
-2 Strain (%)
-400
(a)
-200
(b)
Fig. 7. Hysteresis loop of Ti-49·1 at% Al PST crystals cyclically deformed at several stress-controlled amplitudes: (a) r/J N = 20 cycles; (b) r/J = 45°. N = 20 cycles.
in shape of the loop was found between specimens at cf> = 0 and 45°. After the first cycle the plastic strain remained at zero stress level in a hysteresis loop, and then increased monotonically with increasing applied stress. However, an anomalous change in the loop at cf> = 0 can be seen in the course of further deformation, e.g. at N = 20 cycles, while there is no anomalous dependence of applied stress on the hysteresis loop at cf> = 45° even at further cycles. The plastic strain energy, which is determined by the area of the hysteresis loop, is closely related to the energy required for twinning andlor the to-and-fro motion of dislocations during one cycle. Figure 8 shows the variation in plastic strain energy with applied stress at N = 20 cycles. At cf> = 45° the plastic strain energy increases as the applied 0 the energy increases stress rises, while at cf>
=
= O.
sharply with increasing stress and then shows a sudden drop around 440 MPa followed by a rapid increase. A peak of plastic strain energy is therefore observed around 430 MPa. In both cases this energy decreases with increasing number of cycles, since dislocations are piled up, tangled and interact with each other resulting in adequate cyclic hardening. The anomalous change of plastic strain energy was observed in the course of cyclic deformation although the value of the energy decreased with increasing number of cycles. The relation between the plastic strain energy and applied stress at cf> = 0 in Fig. 8 can be divided into three regions. In region I where stress amplitudes below 430 MPa are applied, ordinary dislocations are operative in V-type domains and twinning occurs in T-type domains at an early stage of deformation. The increase in applied stress requires
Plastic anisotropy and fatigue of TiAI PST crystals
120
140
160
180
200
Stress amplitude (MPa)
Fig. 8. Plastic strain energy-stress amplitude curves of Ti-49·1 at'Y., Al PST crystals cyclicalIy deformed at stress-controlIed amplitude with c/J = 0 and 45°.
the motion of a large number of dislocations, resulting in an increase in plastic strain energy. During the to-and-fro motion of dislocations, a vein-like structure is formed in V-type domains and cyclic hardening occurs preferentially in these domains with increasing applied stress and number of cycles. However, twinning in T-type domains does not effectively contribute to the cyclic hardening, although extrusions develop on the specimen surface during cyclic loading at stress near the peak in Fig. 8. In the arphase no significant feature is observed on the specimen surface or in the substructure. In region II where the plastic strain energy decreases sharply, twinning begins to occur in V-type domains accompanied by the motion of ordinary dislocations, while twinning is predominantly operative in T-type domains similar to region I, although no significant surface extrusions are observed. The to-and-fro motion of ordinary dislocations, which form a vein-like structure with highly tangled dislocations at a further stage, is interrupted by deformation twins in T-type domains, and these domains are hardened. The hardening of V-type domains-which is accelerated by twinning~orresponds to a sudden drop of plastic strain energy in this region. Prism slip is activated in the a 2-phase since sufficient stress and strain are applied. In region III, twins are observed in both T- and V-type domains since there is enough stress applied to activate them. In Tvtype domains heavy extrusions develop on the specimen surface and slips aiding the development of these extrusions
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are accommodated near twin boundaries. In V-type domains, ordinary dislocations which produce numerous loops and form veins are operative, accompanied by twinning. Although cyclic hardening becomes remarkable with an increase in applied stress that accelerates the formation of veins, dislocations can glide against the residual stress from lattice defects under high applied stress, resulting in an increase in plastic strain energy. Slip traces on the prism plane in the a2-phase also became remarkable. In contrast, plastic strain energy increases monotonically at q, = 45° as the applied stress rises. Twinning and slips occur on (1 1 1) planes parallel to the lamellar planes in T-type and Dtype domains. There is no significant stress dependence in slip markings and deformation substructure, although the frequency of slip traces and the density of twins and dislocations increase with increasing stress. Even at high applied stress no slip markings were observed in the arphase, which does not directly play a key role in cyclic deformation. In general, fatigue life decreases with increasing applied stress: in fact, the fatigue life of TiAI PST crystals cyclically deformed at q, = 45° decreases monotonically as the applied stress rises. However, the fatigue life and fracture behaviour of TiAl PST crystals at q, = 0 are closely related to the deformation mode of T-type and V-type domains. At the lower stress level in region I, TiAI PST crystals showed a long fatigue life and specimens did not fracture even after I X 105 cycles. Anomalous damage appeared for fatigue lives between regions I and II, however. At the peak in the plastic strain energy-applied stress curves, fatigue life was suddenly shortened. In region II on the border with region I, fatigue life was long and specimens did not fail at I X 105 cycles. Fatigue life then decreased with increasing applied stress as is true for other materials. At the peak of the curve, surface extrusions in T-type domains developed and the stress concentration near the extrusions was responsible for the initiation of microcracking, resulting in specimen failure. In region II plastic deformation is uniformly distributed and divided in T-type and V-type domains. Formation of the surface extrusions is suppressed and fatigue life is again improved. In region III the increase in applied stress hardens V-type domains and highly localized stress initiates cracking near a2/Y boundaries and in the a2-phase. In the previous section it was described that twinning is directional and occurs in only one
Y. Umakoshi et al.
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direction; in aT-type domain twins can only be activated in tension or in compression depending on the loading axis. Even if twins were activated in the positive half-cycle during cyclic loading, twinning would not occur in the negative halfcycle. When cyclic load is applied in tension/tension mode and in tension/compression, the frequency of activation of twins can be changed between specimens with the same loading axis. At the same applied stress, the number of cycles to fracture of a specimen loaded in the tension/tension mode is larger than that of a specimen loaded in tension/ compression mode. For example, at a = 500 MPa a specimen of Ti-49·1 at% AI PST crystal deformed at cP = 0 under cyclic loading in the tension/tension mode fractured at N = 746 cycles, while in the tension/compression mode it fractured at N = 37 cycles. The number of deformation twins increased more in the tension/compression mode than in the tension/tension mode. There was also a significant difference in slip markings on the specimen surface between the two cyclic loading modes as shown in Fig. 9. In the tension/compression mode the surface extrusions developed in a peculiar domain and fatigue life was shortened. This result strongly suggests that deformation twins are detrimental to microcrack initiation and fatigue life. 6 EFFECT OF ADDITIONAL ELEMENTS ON CYCLIC BEHAVIOUR AND FATIGUE LIFE In monotonic deformation, the addition of a third element-such as Nb, V, Mn, Cr or Mo-is known to be effective in improving strength and ductility.P-" Addition of V, which may decrease the stacking fault energy, easily activates deformation twins in the ,,-phase, resulting in improved ductility at room temperature.P:" The addition of Nb, in contrast, increases the stacking fault energy
but may suppress the anisotropy of plastic behaviour due to the activation of pyramidal slip in the €X2-phase. 30 The effect on cyclic behaviour and fatigue life of adding V and Nb, which respectively may induce a decrease and an increase in the stacking fault energy of 'Y"phase TiAI, has also been investigated." Figure 10 shows the variation in stress amplitude with cumulative plastic strain for binary and ternary TiAI PST crystals cyclically deformed at cP = 45° at de = ±D·5 and ±0·3%. Adding a third element has a significant effect on the fatigue life. Nb is effective but V is harmful to fatigue life at cP = 45°. A similar tendency was observed at cP = 0 but was not significant in the case of Nb addition. In V addition, twinning is more activated and T-type domains are adequately hardened at an early stage of cyclic loading. With further loading, plastic deformation is concentrated in D-type domains. Highly dense dislocations in these domains produced during their to-and-fro motion cannot easily pass through domain boundaries and enter the neighbouring T-type domains because the T-type domains are hardened by the high density of deformation twins. In addition, frequent twinning may produce the surface steps and extrusions which act as initiation sites for microcracking. Once a microcrack is initiated at cP = 45°, it easily develops into a macrocrack and fracture occurs. Frequent twinning induced by V addition aids in starting a microcrack, resulting in short fatigue life. In contrast, since Nb addition suppresses twinning, plastic strain during cyclic loading does not concentrate in 250
,,
TI-4B.2at%AI-1.2at%V
"iii 200 0-
6
,, I
I
I
CD
"0 ::J
:t::
a. ~
TI-49.1at%A1
150
CIl CIl
CD
--
b
fJ)
100 ~e=:t0.5%
1000%
(a)
~e=:t0.3%
50 L...L.
--J
(b)
Fig. 9. Effect of deformation mode on slip markings of Ti--49·I at% AI PST crystals cyclically deformed at q, =0 and ,ju = t420 MPa to )05 cycles: (a) tension/tension mode; (b) tension/compression mode.
Cumulative plastic strain (%)
Fig. 10. Variation in stress amplitude with cumulative strain of binary and ternary TiAI PST crystals cyclically deformed at q, 45°.
=
Plastic anisotropy and fatigue of TiAI PST crystals
specific domains. The highly tangled dislocation structure is gradually formed, accompanied by moderate cyclic hardening. In D-type domains, dense locally tangled dislocations are formed since Nb addition decreases the separation distance between Schockley partials dissociated from an ordinary dislocation and the dislocation can easily cross-slip. Nb addition improves the fatigue life at cP =45°. At cP = 0, addition of Nb improves the fatigue life but it does not work effectively as it does at cP = 45°. The fatigue life is shortened by V addition. In ternary TiAI PST crystals containing Nb the average domain size is similar to that of binary ones, but the average lamellar spacing is much larger. The propagation of cracking, which is effectively interrupted by lamellar boundaries, is important in fatigue life at c/J = O. 7 TEMPERATURE DEPENDENCE OF CYCLIC DEFORMATION AND FATIGUELIFE The cyclic behaviour and fatigue life of TiAI PST crystals show a strong temperature dependence. Figure 11 shows the relation between stress amplitude (S) and number of cycles to failure (N) of TiAI PST crystals cyclically deformed at different temperatures. The S-N curves shifted to higher stress level with decreasing temperature, although there was no significant difference between specimens at 25°C and 400°C with cP = O. At cP = 45° the S-N relation showed a straight line at all temperatures and the fatigue life became shorter as the stress amplitude increased. At cP = 0 a similar tendency in the S-N relation was found.
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In monotonic deformation, the mobility of ordinary and superlattice dislocations decreases and yield stress increases with decreasing temperature. Since the cross-slip event and/or climb motion of moving dislocation are thermally assisted, dislocations show the tendency to glide on the same slip plane at low temperature. As previously described, the vein-like structure composed of highly dense dislocations and low density regions in V-type domains is responsible for the cyclic hardening. Some segments of moving dislocations that perform the climb motion and/or the cross-slip event produce dislocation, loops and debris during cyclic loading. During the to-and-fro motion of dislocations, highly tangled dislocations form a vein-like structure whose formation is assisted by the formation of loops and debris. Since the crossslip of ordinary dislocations hardly occurs at -196°C during their to-and-fro motion in V-type domains, formation of the vein-like structure is suppressed, resulting in weak cyclic hardening. Accommodation slip near deformation twins in 'f-type domains, which causes development of surface extrusions, is also suppressed at low temperature. In contrast, cross-slip and/or climb motion occurred easily at room temperature during cyclic loading. The mobility of dislocations increased rapidly with increasing temperature. Deformation is controlled by the motion of ordinary dislocations and twinning is not dominant even at high applied stress in V-type domains. However, the surface extrusions were still significant in T-type domains since accommodation slip near deformation twins occurred easily. At high temperatures, for example at 700°C, the mobility and climb motion of ordinary dislocations increased rapidly and the formation of extrusions was suppressed. The rearrangement of dislocations during cyclic loading softens the crystal, resulting in short fatigue life.
600
'lsoo
8 CRACK PROPAGATION AND FRACTURE
~ ~ 400
.E
c.
~300
_ _ _ --=-e_ +=45' e _-::-::::-::-
_ t-
-196 'C
CIl CIl
~ 200 100
~
---..----a-a_ 4OO'C
1~
1~
1~
1~
Cycles to failure
1~
1~
Fig. 11. Variation in stress amplitude with the number of CYcles to failure of Ti-49·1 at% Al PST crystals cyclically deformed at various temperatures.
The initiation and propagation of cracking and the fracture of cyclically deformed TiAI PST crystals depend strongly on strain amplitude, stress amplitude, orientation of loading axis and deformation temperature. At c/> = 45°, twinning and a long movement of dislocations that were concentrated in localized regions of the y-phase occurred predominantly on the (l 1 I) plane in y-domains parallel to the lamellar boundaries and aided the development of heavy extrusions. On a macroscopic scale fracture occurs parallel to the lamellar
574
Y. Umakoshi et al.
boundaries, independent of deformation temperature, microstructure and the applied stress. Highly concentrated stress at extrusions-which are closely related to twinning-has an important function in crack initiation. At low temperatures twinning stress increases and formation of the vein-like structure that causes strong cyclic hardening is hindered due to the suppression of cross-slip of moving dislocations during cyclic loading. The formation of extrusions, which act as a trigger to initiate a microcrack, is suppressed. The fractograph in Fig. 12(a) shows a flat surface in a cleavage mode but was not accompanied by a zigzag motion, suggesting that y/y domain boundaries do not play an important role in suppressing the crack propagation. Refinement of lamellae was effective in improving the fatigue life at all tested temperatures. Fine lamellae and small ordered domains lead to homogeneous deformation which assists in suppressing the formation of extrusions and initiation of a microcrack, and prolongs fatigue life. However, there is not much difference in the fracture morphology between specimens with fine and coarse lamellae. At cP = 0, slip markings on the {I I I} plane in the y-phase crossing the lamellar boundaries were observed but they were unable to enter the a2phase at low strain amplitude. At high strain amplitude, microcracks were often seen on the (000 I) plane in a2-plates but they did not have a direct connection with specimen failure. According to our recent result on the orientation dependence of fracture of Ti3AI single crystals, not only the basal planes but also the {I 0 I 2} pyramidal planes were susceptible to brittle fracture." In TiAI PST crystals, a crack propagating on the (000 1) plane in an arplate suddenly changed course across the y-phase and proceeded again in a neighbouring arplate, resulting in a zigzag crack path between neighbouring a2-plates. The
brittle nature of fracture on {I 0 12} planes in the arphase may act as a trigger to change the course of a crack path and be responsible for this zigzag motion. Figure 12(b) shows a scanning electron micrograph of the fracture surface of a specimen cyclically deformed at cP =O. There were many terraces and projections on y-domains and a2-plates, suggesting that the lamellar boundaries may act as a barrier to suppress the further propagation of cracks, and they do this effectively in fine lamellae. In T-type domains, twins accompanied by dislocations were observed but multiplication of new twins does not occur easily because twins do not perform the reverse motion during cyclic loading. When V-type domains are sufficiently hardened by the formation of vein-like structure during the to-and-fro motion of dislocations, T-type domains begin to deform preferentially with further cyclic deformation: this deformation concentrates near the twins, which act as nucleation and propagation sites for slip and/or twinning and aid the development of surface steps. The large residual stress field associated with inhomogeneous distribution of dislocations and twins acts as a trigger to nucleate a microcrack. The extrusions also act as a nucleation site for microcracks. At room temperatures and below, an anomalous peak appears in the plastic strain energy-applied stress curves. At small applied stress below the peak, twinning occurs in T-type domains but V-type domains are deformed by the motion of dislocations. In contrast, at large applied stress above the peak, twinning is operative even in V-type domains. Interaction between twins and highly dense dislocations forms a concentrated residual stress field which acts as a trigger for microcrack initiation. Therefore, frequent microcrack nucleation at high applied stress leads to a short fatigue life. At high temperatures the anomalous peak disappears in the plastic strain energy-applied stress curves because twinning is suppressed in V-type domains even at higher applied stress. Fracture morphology, therefore, does not show a strong dependence on applied stress.
9 CONCLUDING REMARKS (a)
(b)
Fig. 12. Scanning electron fractographs of Ti-49·) at°/r, Al PST crystals cyclically deformed at room temperature: (a) t/J =45°, ae = to·)'X" N = 3300 cycles; (b) t/J =0, ae =to·3%. N =2500 cycles.
Large anisotropy of strength, ductility, cyclic hardening and fatigue life is observed in TiAI PST crystals. Variation in deformation mode, such as twinning and the to-and-fro motion of ordinary
Plastic anisotropy and fatigue of TiAI PST crystals
dislocations, is responsible for this. The extrusions connected with twinning are important in the initiation of microcracks. In monotonic deformation, addition of V which may decrease the stacking fault energy activates twinning in Ti-rich TiAI alloys at room temperature, resulting in improved ductility at room temperature. However, twinning is harmful to the fatigue life because it accelerates the formation of extrusions on the specimen surface and leads to highly concentrated residual stress near lamellar and domain boundaries. From the viewpoint of twinning, addition of Nb, which may increase the stacking fault energy, suppresses the twinning but does not effectively improve the fatigue life at cP = O. Deformation of arphase does not directly contribute to the cyclic deformation behaviour. but the accommodation of slip in a2- and 'Y'"phases is required to relax the highly concentrated residual stress at the a21y lamellar boundaries and suppress microcrack initiation. Since a2/y lamellar boundaries act as an effective barrier to stop microcrack propagation, the fatigue life must be controlled by the initiation and propagation of the microcrack. Polycrystalline alloys with lamellar structure show better fatigue life than those with duplex structure. To obtain a good combination of strength, ductility and fatigue life, refinement of lamellae is the most effective method. Even at high temperatures the lamellar boundaries still act as an effective barrier for microcrack propagation and good fatigue properties are maintained. The choice of alloying elements and thermo-mechanical treatment should be considered to obtain fine lamellae.
ACKNOWLEDGEMENT Y. Umakohi would like to thank the Ministry of Education. Science. Sports and Culture of Japan for a Grant-in-Aid for Scientific Research on the Priority Area 'Intermetallic Compounds as New High Temperature Structure Materials'.
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4. Umakoshi, Y., Nakano, T. & Yasuda. H. Y., in Mechanical Properties and Phase Transformation of Multi-Phase Intermetallic Alloys. eds A. F. Giami et al.. TMS, Warrendale, 1995, p. lSI. 5. Fujiwara, T., Nakamura. A., Hosomi, M.• Nishitani, S. R., Shirai, T. & Yamaguchi, M.• Phil. Mag.• A61 (1990) 591. 6. Inui, H., Oh, M. H., Nakamura, A. & Yamaguchi, M., Acta Metall. Mater., 40 (1992) 3095. 7. Umakoshi, Y. & Nakano, T., Acta Metall. Mater., 41 (1993) 1155. 8. Umakoshi. Y. & Nakano, T., Mater. Sci. Eng. , AIS2 (1992) 81. 9. Umakoshi, Y. & Nakano, T.• ISIJ Int., 32 (1992) 1339. 10. Sastry, S. M. L. & Lipsitt, H. A.• Metall. Trans, 8A (1977)299. II. Dowling, W. E. Jr, Donlon, W. T. & Allison. J. E., in High- Temperature Ordered lntermetallic Alloys IV, eds L. A. Johnson et al.. MRS Symp. Proc., 213 (1991) 561. 12. Soboyejo, W.O., Deffeyes, J. E. & Aswath, P. B., Mater. Sci. Eng., A138(1991) 95. 13. James. A. W. & Bowen. P., Mater. Sci. Eng.• AlS3 (1992) 486. 14. Davidson, D. L. & Campbell. J. B., Metall. Trans.,24A (1993) 1555. 15. Kumpfert, J., Kim, Y.-W. & Dimiduk, D. M.• Mater. Sci. Eng.• A1921193 (1995) 465. 16. Sloloff, N. S.• Fuchs, G. E.• Kuruvilla, A. K. & Choe, S. 1., in High-Temperature Ordered Intermetallic Alloys Il, ed. N. S. Stoloff, MRS Symp. Proc., 81 (1987) 247. 17. Stoloff, N. S.• Smith. T. R. & Castagna, A.. in HtghTemperature Ordered Intermetallic Alloys V, ed. 1. Baker, MRS Symp. Proc., 288 (1993) 59. 18. Yauda, H . Y.• Nakano, T. & Umakoshi, Y., Phil. Mag., A71 (1995) 127. 19. Umakoshi, Y., Yasuda, H. Y. & Nakano, T., Mater. Sci. Eng., A194 (1995) 43. 20. Yasuda. H. Y., Nakano, T. & Umakoshi, Y., Phil. Mag.• (1996) in press. 21. Yasuda. H. Y.• Nakano, T. & Umakoshi, Y., Phil. Mag., (1996) in press. 22. Nakano, T. & Umakoshi, Y., Intermetailics, 2 (1994) 185. 23. Schechtman. D., Blackburn. M. J. & Lipsitt, H. A.• Metall. Trans.,5 (1974) 1373. 24. Court, S. A.• Vasudevan, V. K. & Fraser. H. L.. Phil. Mag.• A61 (1990) 141. 25. Inui, H., Nakamura, A., Oh, M. H. & Yamaguchi, M.• Phil. Mag., A66 (1992) 557. 26. Hung, S. C; in Structural Intermetallics, ed. R. Darolia et al.• TMS. Warrendale, 1993, p. 299. 27. Hashimoto. K. & Kimura, M" in Structural Intermetallics, ed. R. Darolia et al... TMS, Warrendale, 1993, p.309. 28. Whang. S. H.• Hahn, Y. D., u, Z. X. & u Z. c, in JIMIS-6 : Intermetallic Compounds-Structure and Mechanical Properties, ed, O. Izumi, lap. Inst. Metals, Sendai, 1991. p. 763. 29. Yao, K.-F.• Inui, H., Kishida, K. & Yamaguchi, M.• Acta Metall. Mater.• 43 (1995) 1075. 30. Nakano, T., Yokoyama, A. & Umakoshi, Y., Scripta Metall. Mater., 27 (1992) 1253. 31. Umakoshi, Y., Yasuda, H. Y.• Nakano. T. & Nakazawa, J.• in Gamma Titanium Aluminides, eds Y.-M. Kim, R. Wagner & M. Yamaguchi. TMS. Warrendale, 1995, p.697. 32. Umakoshi, Y., Nakano, T. & Ogawa, B., Scripta Metall. A!ater., (1996) in press.