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Effect of purity and second phase on ductility of TiAl
Scripta
METALLURGICA
Vol. 22, pp. 1 7 2 5 - 1 7 3 0 , 1988 P r i n t e d in the U.S.A.
P e r g a m o n Press plc All rights r e s e r v e d
Z~'F~OT OF PURITY AND SECOND PHASE ON DUCTILITY OF TiAI
T. Kawabata, M. Tadano* and O. Izumi Institute for Materials Research (Formerly The Research Institute for Iron, Steel and Other Metals) Tohoku University, Sendai, 980, Japan. * Graduate Student, Tohoku University ( R e c e i v e d May 19, 1988) ( R e v i s e d A u g u s t 29, 1988)
Introduction TiAI has high strength [I, 2], high creep resistance [I, 3], good oxidation resistance [1] at high temperatures and low density so that it is potentially useful for a heat resistant structural alloy. TiAI is a candidate for materials of rotary structural parts in jet or gas turbine engines [4]. However, it has not been used yet industrially [4] because it is very brittle below about 900 K [I, 2]. The phase diagram of Ti and A1 binary alloys contains Ti3AI , TiAI and TiAlq intermetallic compounds [5, 6]. The hardness of TiAI alloys is lower than those of TiqAl, TiqAI+TiAI and TiAIB+TiAI alloys [6, 7]. The tensile properties of TiAI were studied ~irst i K 1 9 5 6 [I] by McAn~rew and Kessler who found that fracture occurred in a brittle mode with no appreciable elongation at room temperature. The temperature dependence of yield stress and ductility in TiAI formed by powder metallurgy has been studied by Lipsitt et al in a temperature range from 293 to 1273 K [2]. The ductility was nearly zero below 973 K but increased rapidly above 973 K. The ductility improvement due to the addition of third elements has been studied by Blackburn and Smith [8], Martin et al [9], Tsujimoto [10] and Hashimoto et al [11]. The ductility of Ti-48 at% A1 has been improved up to 5.1 Z by the addition of 0.5 and 2.5 at% V at 533 K [8] and to nearly 3 Z by the addition of I at% Mn at room temperature [11]. A fracture mechanism in TiAI has been proposed by Kawabata and Izumi [12] to explain (111) slip plane fracture and cleavage fracture on [110), [IO1), (010) and (001) planes. There have been very few studies on ductility in TiAI based on structural observations. In the present study we report that the ductility of a binary Ti-48 at% A1 alloy is improved by the use of high purity base metals while a lamellar structure of the matrix, TiAI, and the second phase, Ti3AI , in the alloy also contributes additionally to the increase of ductility. Experimental Procedure A Ti-48 at% A1 alloy was melted from high purity Ti and A1 using a nonconsumable tungsten electrode arc-melting furnace in an argon atmosphere. The size of button ingots was 50 mm in diameter and 15 mm in thickness. Ti with 99.91 mass % purity in a plate-like shape was used. The Ti contained the impurities Fe;O.030, Cr;O.OO6, Ni;O.O08, 0;0.038, C;0.O06, N;O.O03 and H;0.O04 mass %. T h ~ p u r i t y of A1 was 99.99 mass Z. After annealing at 1573 K for 259.2 ks in vacuum of Ixi0 -5 Pa, the ingots were cut with a precision resinoid cutter into slices 1.5 mm in thickness. The plates were polished by emery papers from No. 800 to 2400 into 1.0 mm in thickness. Tensile specimens with a gage portion of 2 mm in width, 14 mm in length and I mm in thickness were formed with a spark erosion machine. The surfaces cut by the spark erosion machine were polished by emery papers from No. 250 to 1500. The specimens were electrolytically polished by the method shown in Ref. [13]. The tensile tests were performed with an Instron type testing machine at a strain rate of 6xi0 -4 s -1 in air. A supplementary test of a Ti-52 at % A1 alloy made from 99.7 mass % purity Ti and 99.99 mass % purity A1 was also performed in tension. A specimen for a bending test was cut with a precision resinoid cutter into 0.8 mm in thickness and polished by emery papers from No. 240 to 2400 finally into 0.4 mm in thickness. The specimen, which was not electropolished, was bent slowly by hand. The microstructures of undeformed and deformed specimens were observed with a Nomarski type differential interference optical microscope. In order to observe interaction between the second phase and slip bands the specimen
1725 0 0 5 6 - 9 7 4 8 / 8 8 $5.00 + .00 C o p y r i g h t (c) 1988 P e r g a m o n P r e s s
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was etched by the Kroll solution after the tensile t e s t . The fracture surface was observed with a scanning electron microscope (JSM-T20). Results and Discussion Mechanical properties Figure I shows the tensile stress-strain curve of an annealed specimen. The proof stress at 0.2 % plastic strain is 296.0 MPa, the fracture stress 411.5 MPa and the fracture strain 2.7 %. The curve is smooth and the yield stress is lower than that in Ref. [8]. The unannealed specimen showed that the proof stress is 466.0 MPa, the fracture stress 490.7 MPa and the fracture strain 0.3 %. In the ordinary purity Ti-52 at% A1 alloy, very brittle fracture occurred at a stress level far lower than the elastic limit. Even if the A1 rich alloy was made from high purity Ti, the ductility of the alloy might be restricted within the elastic limit, which means that the composition is important for ductility. However, a single phase structure of Ti-48 at % A1 cannot be made by quenching so that we cannot compare the ductility between the single phase and the two phase structures of the alloy because of the limitation imposed by the form of the phase diagram. The ductility in bending of Ti rich TiA1 has been improved by a Mn addition resulting in a fracture strain of' 3 % [10]. At 533 K, the tensile elongation of the Ti-48 at% AI, Ti-48 at% A1-0.5 at% V and Ti-48 at% A1-2.5 at% V alloys has been reported as 2.1, 5.1 and 5.1%, respectively [8]. However, the tensile elongation 2.7 % in the present study is the highest data at room temperature for the binary alloy. Figure 2 shows a bent specimen, for which the radius of curvature is about 9 mm and the tensile strain at the surface is about 2.2 %. One of the important reasons that ductility was improved, in the present study, is the use of high purity base metals. The decrease of the impurity contents is expected to lower the flow stress resulting in the suppression of cleavage and grain boundary fracture. However,' as shown in the above references and also in the present study, it seems that the Ti-48 at % A1 alloys generally have a larger elongation than alloys with different compositions, indicating that the reason for the increased ductility is present in the microstructure itself. Microstructure The microstructures of unannealed and annealed specimens are shown in Fig. 3 (a) and (b), respectively. A very fine lamellar structure is observable in Fig. 3 (a). The lamellae consist of the matrix, TiA1 (y), and the second phase, TisA1 (a2). After annealing at 1573 K for 259.2 ks, the structure changes into nearly equiaxed g~ains and the TisA1 phase becomes coarse plates, which are also seen on the grain boundaries. The grain size o~f the annealed specimen measured by the intercept method was about 40 wm, the average periodic distance between the second phase plates 3 ~m and the thickness of the plates 0.7~3 Dm. The habit Rlane and the orientation relationship between TiA1 and Ti.A1 nhases have been renorted as (111~ //(0001) ^ and (111)v//(OOO1)a9 and <1TO>v//<1120>~9 [I~], respectively. However, the TidAl plates in Fig. ~(b) are not precisely p'arallel an~ the interfaces are not strictly smooth, s~ggesting that the coherency between phases is not perfect in that interface steps and dislocations are present. It is characteristic that the grain size of the annealed specimen is smaller than that of the unannealed specimen. Figure 4 shows optical micrographs near the fracture portion of the tensile specimen observed at (a) low and (b) high magnifications. The surface becomes rough and deformation occurs uniformly over the gage length. Figure 5 (a) shows the feature of grain boundary cracking, suggesting that such a crack could be the initiation site of fracture, as illustrated by the arrow and the letter "c". The tensile axis is illustrated by the arrow and the letter "TA". Other grain boundary cracks were observed on surfaces which were inclined about 45 degrees to the tensile axis, suggesting that the cracks were caused by grain boundary sliding and accommodating multiple slips near the grain boundary. That is, the heavy deformation near the grain boundary would cause dislocation reactions [12] resulting in fracture along the grain boundaries. The existence of the isolated cracks means that the material has crack tolerance and is rather tough. Ductility is related to how deformation occurs in the lamellar structure and whether cracking occurs at intersecting points of the slip bands and plates. Both the flow stress and ductility in TidAl are known to be larger than those in single phase TiA1 alloys at room temperature [15]~ although the ductility of Ti3A1 is not as large as that in the Ti-48 at % A1. In the present study, it is shown that the composite structure, such as that of the Ti-48 at % A1 alloy, results in a larger ductility than that expected from the rule of mixing. Three cases on the relationship between the second phase plate and slip bands were observed. Figure 5(b) represents the slip traces parallel to the plates of the second phase resulting in the straight and rather heavy slip markings, which are shown by the arrow and the letter "p". The arrow and
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the letter "s" in Fig. 5 (b) show a thick plate obstructing slips which shortens the length of slip bands and plays the same role as a grain boundary. On the contrary, as shown by the letters "t1" and "t2" in Fig. 5 (c), some slip bands seem to cross thin plates of the second phase. The deformation occurs continuously through the plates and no crack occurs at the crossing points. Therefore, the presence of the second phase plates, with higher strength and ductility than the matrix, and the goodness of conjunction of the interfaces do not bring brittleness but result in strengthening. This is the reason why the appropriate distribution of the second phase plates brings toughening. The orientation relationship between TidAl and TiA1 is also satisfied in the present alloy. We consider the case (111)v//(0001)~2 and [TTOSv//K1120]a2. When slip bands cross the second phase, the (TIT) plane in T'iA1, for example, corresponds nearly to a pyramidal plane (TI01) in TidAl. The Burgers vectors [STO) and [101] in TiA1 are very near to the Burgers vectors [1120] an8 [1121S in TiBA1. Therefore, the continuity of deformation would be kept at interfaces between the TiA1 ~nd Ti3A1 phases. The unannealed and also annealed TiA1 with A1 less than about 50 at ~ contained many twins and second phase thin plates [16, 17). Many dislocations were observed on the interfaces of the twins and second phase plates [16, 17], which were easily movable [16, 18]. Therefore, interface slip might be one of the deformation mechanisms in TiA1 with a lamellar structure such as in the present study. In Fig. 5 (c), the slips along the interface between TiA1 and TiBA1 is actually observed at the portion shown by the arrow and the letter "tS". However, in-the unannealed specimen, the fracture strain was 0.3 ~, although most slip bands occurred along plates in the matrix of the lamellar structure. Therefore, it was concluded that the interface slips do not contribute much to fracture strain. Figure 5 (c) exhibits wavy slip traces in a large grain, shown by the arrow and the letter "w", which is evidence for cross-slip. Cross-slip in TiA1 has not been observed by optical microscopy. Double slip events were observed in some grains, which would cause dislocation reactions K12S at the intersection resulting in strengthening and also would relate to fracturing [12]. The work hardening rate in TiA1 single crystals is higher than that in ordinary metals and alloys [19S. Therefore, in TiA1, the dislocation reactions such as those proposed in Ref. [12] should be considered, additionally to the ordinary dislocation reactions for polycrystals. The use of high purity base metals in the present study decreased flow stress in grains and toughened grain boundaries by lowering the grain boundary segregation of harmful elements. The decrease of flow stress also causes toughening because the fracture caused by a given fracture mechanism [12] is suppressed. The macroscopic structure such as the lamellar spacing and the thickness of the second phase plates would not be changed much by the purity of alloy. Fracture strain Fractography showed that the fracture mode in the present study was mainly of the transgranular type, namely, cleavage fracture with rather flat planes and slip plane fracture with saw-toothed zigzag planes due to dislocation reactions. Thus, it seems that the dislocation pile-ups caused transgranular fracture. We consider the fracture stresses in the structures with and without the second phase plates, based on a fracture mechanism due to stress concentration near the front of the dislocation pile-ups. We assume that when the stress concentration reaches a given stress level near the pile-up front, fracture will initiate. As a first approximation, we consider simple cubic grains with and without second phase plates having a side length which corresponds to the grain size being d, the lamellar spacing s and the thickness of the second phase plates w, in which the pile-up lengths are (s - w) and d and the applied stresses o S and 42, in the structures with and without the second phase plates, respectively. Subsequently, the subscripts, 1 and 2, show the structures with and without the second phase plates, respectively. Under %he condition that fracture occurs at the same level of the stress concentration, the following relations are obtained [20). N 1 = (1 - ~ ) ( s
- W)~l/~b
N2 = (1 - ~ ) ( s
- W)Ol(d/(s
= Nl(d/(s
- w)) 1/2
(1) - w))l/2/~b
(2) (2')
When each dislocation sweeps an area (s - w)d and d 2 in the structures with and without the
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second phase plates, respectively, the ratio of shear strains, Yl and Y2' is given as the ratio of the number of pile-up dislocations. YIIY2 = NIIN2" (3) The ratio of the stresses is also given as
o11o2
=
(dl(s
-
w))I12.
(~)
The structural parameters in the present study are s=3 ~m, w=0.7 ~m, d=40 ~m and the applied yield stress, O 0 o~ i=296.0 MPa, while the material parameters are ~=O.247, W=71.4 GPa and b=0.2828 nm, resulting in N1=25 , N2=I06 and y1/Y2=1/4.2. This analysis shows that the presence of the second phase plates, which play the same role as grain boundaries, lowers the fracture strain. This conclusion can partly explain the tendency of the experimental results, that the ductility in the annealed specimen with a larger lamellar spacing is larger than that in the unannealed specimen with a smaller lamellar spacing while the yield stress varies in the opposite manner. Moreover, the above result seems to imply that the matrix of the Ti-48 at % A1 alloy has far larger ductility than that in the composite structure containing the second phase plates. However, TiA1 with an A1 composition of less than 50 at % already contains the second phase [16]. Therefore, we cannot investigate the ductility of only the matrix. Also, the large fracture strain in the Ti-48 at % A1 alloy should be related to the characteristics of the interfaces between Ti3A1 and TiA1 and twins, the second phase and, perhaps, the matrix. Conclusion The mechanical properties of high purity Ti-48 at% AI were studied in tensile tests. The fracture elongation was improved up to about 2.7 %. The improvement of ductility is caused by the lowering of flow stress because of both the use of high purity base metals and the characteristics of the second phase Ti3A1. Acknowledgment This study was supported partly by Keikinzoku Shogakukai (The Light Metal Educational Foundation, Incorporated) and Grant-in-Aid for Scientific Research of the Ministry of Education, Science and Culture, Japan (C62550512). Reference I. 2. 3. 4. 5. 6. 7. 8. 9.
J.B. MacAndrew and H.D. Kessler, J. Metals, 8, 1348 (1956). H.A. Lipsitt, D. Shechtman and R.E. Schafrik, Metall. Trans. 6A, 1991 (1975). P.L. Martin, M.G. Mendiratta and H.A. Lipsitt, Metall. Trans. 14A, 2170 (1983). H.A. Lipsitt, in High Temperature Ordered Intermetallic Alloys, Edited by C.C. Koch, C.T. Liu and N.S. Stoloff, Materials Research Society Symposia Proc. 39, p. 351, (1985). H.R. Ogden, D.J. Maykuth, W.L. Finlay and R.I. Jaffee, J. Metals, 3, 1150 (1951). E.S. Bumps, H.D. Kessler and M. Hansen, J. Metals, 4, 609 (1952). H.R. Ogden, D.J. Maykuth, W.L. Finlay and R.I. Jaffee, J. Metals, 5, 267 (1953). M.J. Blackburn and M.P. Smith, United States Patent No. 4,294,615, Oct. 13, 1981. P.L. Martin, H.A. Lipsitt, N.T. Nuhfer and J.C. Williams, Titanium '80, Science and Technology, Edited by H. Kimura and O. Izumi, Metall. Society of AIME, New York, p. 1245
(1980). 10. T. Tsujimoto, Titanium and Zirconium, 33, 159 (1985) (In Japanese). 11. T. Hashimoto, H. Dohi and Tsujimcto, Preprint of symposium of Japan Institute of Metals, in Plastic Deformation of Ordered Alloys and Intermetallic Compounds (In Japanese), p. 17, July, 1986. 12. T. Kawabata, T. Takezono, T. Kanai and O. Izumi, Acta Metall. 36, 963 (1988). 13. M.J. Blackburn and J.C. Williams, Trans. Met. Soc. AIME, 239, 287 (1967). 14. M.J. Blackburn, in The Science, Technology and Application of Titanium, Edited by R.'Jaffee and N. Promisel, Pergamon Press, New York, p. 633, (1970). 15. H.A. Lipsitt, D. Shechtman and R.E. Schafrik, Metall. Trans. 11A, 1369 (1980). 16. D. Shechtman, M.J. Blackburn and H.A. Lipsitt, Metall. Trans. 5, 1373 (1974). 17. Ye.I. Teytel and E.S. Yakovleva, Fiz. metal, metalloved. 40, No. I, 129 (1975) (in English edition). 18. T. Kawabata and 0. Izumi, Unpublished work. 19. T. Kawabata, T. Kanai and O. Izumi, Acta Metall. 33, 1355 (1985). 20. J.P. HIrth and J. Lothe, in Theory of Dislocations, McGraw-Hill, New York, p. 694, (1968).
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LU n." I.tn
STRAIN
FIG. I
The tensile stress-strain curve of high purity Ti-48 at% AI.
FIG. 3
FIG. 2
A specimen bent by hand. The size of the specimen is 0.4 mm in thickness, 4.2 mm in width and 29.6 mm in length. The radius of the curvature is 9 mm and the tensile strain, Es, at the surface is 2.2 %.
Optical microstructures of (a) an unannealed and (b) an annealed specimen which is heated at 1573 K for 259.2 ks.
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FIG. 4
The gage portion of the tensile specimen deformed at room temperature at strain rate 1.2×10 -~ s-I in air. (a) Low and (b) high magnifications.
FIG. 5
Optical micrographs showing (a) grain boundary cracking, (b) slip bands parallel to the plates of the second phase and (c) wavy slip bands. The arrow and the letter "TA" in (a) are the tensile axis and "c" the grain boundary cracking. The arrow and the letters "s" and "p" in (b) show the second phase plates blocking and parallel to slip bands, respectively. The arrows and the letters "w" and "t1" and "t2" in (c) i l l u s t r a t e the w a v y slip bands and the slip bands r u n n i n g across the second phases, respectively.
METALLURGICA
Vol. 22, pp. 1 7 2 5 - 1 7 3 0 , 1988 P r i n t e d in the U.S.A.
P e r g a m o n Press plc All rights r e s e r v e d
Z~'F~OT OF PURITY AND SECOND PHASE ON DUCTILITY OF TiAI
T. Kawabata, M. Tadano* and O. Izumi Institute for Materials Research (Formerly The Research Institute for Iron, Steel and Other Metals) Tohoku University, Sendai, 980, Japan. * Graduate Student, Tohoku University ( R e c e i v e d May 19, 1988) ( R e v i s e d A u g u s t 29, 1988)
Introduction TiAI has high strength [I, 2], high creep resistance [I, 3], good oxidation resistance [1] at high temperatures and low density so that it is potentially useful for a heat resistant structural alloy. TiAI is a candidate for materials of rotary structural parts in jet or gas turbine engines [4]. However, it has not been used yet industrially [4] because it is very brittle below about 900 K [I, 2]. The phase diagram of Ti and A1 binary alloys contains Ti3AI , TiAI and TiAlq intermetallic compounds [5, 6]. The hardness of TiAI alloys is lower than those of TiqAl, TiqAI+TiAI and TiAIB+TiAI alloys [6, 7]. The tensile properties of TiAI were studied ~irst i K 1 9 5 6 [I] by McAn~rew and Kessler who found that fracture occurred in a brittle mode with no appreciable elongation at room temperature. The temperature dependence of yield stress and ductility in TiAI formed by powder metallurgy has been studied by Lipsitt et al in a temperature range from 293 to 1273 K [2]. The ductility was nearly zero below 973 K but increased rapidly above 973 K. The ductility improvement due to the addition of third elements has been studied by Blackburn and Smith [8], Martin et al [9], Tsujimoto [10] and Hashimoto et al [11]. The ductility of Ti-48 at% A1 has been improved up to 5.1 Z by the addition of 0.5 and 2.5 at% V at 533 K [8] and to nearly 3 Z by the addition of I at% Mn at room temperature [11]. A fracture mechanism in TiAI has been proposed by Kawabata and Izumi [12] to explain (111) slip plane fracture and cleavage fracture on [110), [IO1), (010) and (001) planes. There have been very few studies on ductility in TiAI based on structural observations. In the present study we report that the ductility of a binary Ti-48 at% A1 alloy is improved by the use of high purity base metals while a lamellar structure of the matrix, TiAI, and the second phase, Ti3AI , in the alloy also contributes additionally to the increase of ductility. Experimental Procedure A Ti-48 at% A1 alloy was melted from high purity Ti and A1 using a nonconsumable tungsten electrode arc-melting furnace in an argon atmosphere. The size of button ingots was 50 mm in diameter and 15 mm in thickness. Ti with 99.91 mass % purity in a plate-like shape was used. The Ti contained the impurities Fe;O.030, Cr;O.OO6, Ni;O.O08, 0;0.038, C;0.O06, N;O.O03 and H;0.O04 mass %. T h ~ p u r i t y of A1 was 99.99 mass Z. After annealing at 1573 K for 259.2 ks in vacuum of Ixi0 -5 Pa, the ingots were cut with a precision resinoid cutter into slices 1.5 mm in thickness. The plates were polished by emery papers from No. 800 to 2400 into 1.0 mm in thickness. Tensile specimens with a gage portion of 2 mm in width, 14 mm in length and I mm in thickness were formed with a spark erosion machine. The surfaces cut by the spark erosion machine were polished by emery papers from No. 250 to 1500. The specimens were electrolytically polished by the method shown in Ref. [13]. The tensile tests were performed with an Instron type testing machine at a strain rate of 6xi0 -4 s -1 in air. A supplementary test of a Ti-52 at % A1 alloy made from 99.7 mass % purity Ti and 99.99 mass % purity A1 was also performed in tension. A specimen for a bending test was cut with a precision resinoid cutter into 0.8 mm in thickness and polished by emery papers from No. 240 to 2400 finally into 0.4 mm in thickness. The specimen, which was not electropolished, was bent slowly by hand. The microstructures of undeformed and deformed specimens were observed with a Nomarski type differential interference optical microscope. In order to observe interaction between the second phase and slip bands the specimen
1725 0 0 5 6 - 9 7 4 8 / 8 8 $5.00 + .00 C o p y r i g h t (c) 1988 P e r g a m o n P r e s s
plc
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OF TiAI
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was etched by the Kroll solution after the tensile t e s t . The fracture surface was observed with a scanning electron microscope (JSM-T20). Results and Discussion Mechanical properties Figure I shows the tensile stress-strain curve of an annealed specimen. The proof stress at 0.2 % plastic strain is 296.0 MPa, the fracture stress 411.5 MPa and the fracture strain 2.7 %. The curve is smooth and the yield stress is lower than that in Ref. [8]. The unannealed specimen showed that the proof stress is 466.0 MPa, the fracture stress 490.7 MPa and the fracture strain 0.3 %. In the ordinary purity Ti-52 at% A1 alloy, very brittle fracture occurred at a stress level far lower than the elastic limit. Even if the A1 rich alloy was made from high purity Ti, the ductility of the alloy might be restricted within the elastic limit, which means that the composition is important for ductility. However, a single phase structure of Ti-48 at % A1 cannot be made by quenching so that we cannot compare the ductility between the single phase and the two phase structures of the alloy because of the limitation imposed by the form of the phase diagram. The ductility in bending of Ti rich TiA1 has been improved by a Mn addition resulting in a fracture strain of' 3 % [10]. At 533 K, the tensile elongation of the Ti-48 at% AI, Ti-48 at% A1-0.5 at% V and Ti-48 at% A1-2.5 at% V alloys has been reported as 2.1, 5.1 and 5.1%, respectively [8]. However, the tensile elongation 2.7 % in the present study is the highest data at room temperature for the binary alloy. Figure 2 shows a bent specimen, for which the radius of curvature is about 9 mm and the tensile strain at the surface is about 2.2 %. One of the important reasons that ductility was improved, in the present study, is the use of high purity base metals. The decrease of the impurity contents is expected to lower the flow stress resulting in the suppression of cleavage and grain boundary fracture. However,' as shown in the above references and also in the present study, it seems that the Ti-48 at % A1 alloys generally have a larger elongation than alloys with different compositions, indicating that the reason for the increased ductility is present in the microstructure itself. Microstructure The microstructures of unannealed and annealed specimens are shown in Fig. 3 (a) and (b), respectively. A very fine lamellar structure is observable in Fig. 3 (a). The lamellae consist of the matrix, TiA1 (y), and the second phase, TisA1 (a2). After annealing at 1573 K for 259.2 ks, the structure changes into nearly equiaxed g~ains and the TisA1 phase becomes coarse plates, which are also seen on the grain boundaries. The grain size o~f the annealed specimen measured by the intercept method was about 40 wm, the average periodic distance between the second phase plates 3 ~m and the thickness of the plates 0.7~3 Dm. The habit Rlane and the orientation relationship between TiA1 and Ti.A1 nhases have been renorted as (111~ //(0001) ^ and (111)v//(OOO1)a9 and <1TO>v//<1120>~9 [I~], respectively. However, the TidAl plates in Fig. ~(b) are not precisely p'arallel an~ the interfaces are not strictly smooth, s~ggesting that the coherency between phases is not perfect in that interface steps and dislocations are present. It is characteristic that the grain size of the annealed specimen is smaller than that of the unannealed specimen. Figure 4 shows optical micrographs near the fracture portion of the tensile specimen observed at (a) low and (b) high magnifications. The surface becomes rough and deformation occurs uniformly over the gage length. Figure 5 (a) shows the feature of grain boundary cracking, suggesting that such a crack could be the initiation site of fracture, as illustrated by the arrow and the letter "c". The tensile axis is illustrated by the arrow and the letter "TA". Other grain boundary cracks were observed on surfaces which were inclined about 45 degrees to the tensile axis, suggesting that the cracks were caused by grain boundary sliding and accommodating multiple slips near the grain boundary. That is, the heavy deformation near the grain boundary would cause dislocation reactions [12] resulting in fracture along the grain boundaries. The existence of the isolated cracks means that the material has crack tolerance and is rather tough. Ductility is related to how deformation occurs in the lamellar structure and whether cracking occurs at intersecting points of the slip bands and plates. Both the flow stress and ductility in TidAl are known to be larger than those in single phase TiA1 alloys at room temperature [15]~ although the ductility of Ti3A1 is not as large as that in the Ti-48 at % A1. In the present study, it is shown that the composite structure, such as that of the Ti-48 at % A1 alloy, results in a larger ductility than that expected from the rule of mixing. Three cases on the relationship between the second phase plate and slip bands were observed. Figure 5(b) represents the slip traces parallel to the plates of the second phase resulting in the straight and rather heavy slip markings, which are shown by the arrow and the letter "p". The arrow and
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the letter "s" in Fig. 5 (b) show a thick plate obstructing slips which shortens the length of slip bands and plays the same role as a grain boundary. On the contrary, as shown by the letters "t1" and "t2" in Fig. 5 (c), some slip bands seem to cross thin plates of the second phase. The deformation occurs continuously through the plates and no crack occurs at the crossing points. Therefore, the presence of the second phase plates, with higher strength and ductility than the matrix, and the goodness of conjunction of the interfaces do not bring brittleness but result in strengthening. This is the reason why the appropriate distribution of the second phase plates brings toughening. The orientation relationship between TidAl and TiA1 is also satisfied in the present alloy. We consider the case (111)v//(0001)~2 and [TTOSv//K1120]a2. When slip bands cross the second phase, the (TIT) plane in T'iA1, for example, corresponds nearly to a pyramidal plane (TI01) in TidAl. The Burgers vectors [STO) and [101] in TiA1 are very near to the Burgers vectors [1120] an8 [1121S in TiBA1. Therefore, the continuity of deformation would be kept at interfaces between the TiA1 ~nd Ti3A1 phases. The unannealed and also annealed TiA1 with A1 less than about 50 at ~ contained many twins and second phase thin plates [16, 17). Many dislocations were observed on the interfaces of the twins and second phase plates [16, 17], which were easily movable [16, 18]. Therefore, interface slip might be one of the deformation mechanisms in TiA1 with a lamellar structure such as in the present study. In Fig. 5 (c), the slips along the interface between TiA1 and TiBA1 is actually observed at the portion shown by the arrow and the letter "tS". However, in-the unannealed specimen, the fracture strain was 0.3 ~, although most slip bands occurred along plates in the matrix of the lamellar structure. Therefore, it was concluded that the interface slips do not contribute much to fracture strain. Figure 5 (c) exhibits wavy slip traces in a large grain, shown by the arrow and the letter "w", which is evidence for cross-slip. Cross-slip in TiA1 has not been observed by optical microscopy. Double slip events were observed in some grains, which would cause dislocation reactions K12S at the intersection resulting in strengthening and also would relate to fracturing [12]. The work hardening rate in TiA1 single crystals is higher than that in ordinary metals and alloys [19S. Therefore, in TiA1, the dislocation reactions such as those proposed in Ref. [12] should be considered, additionally to the ordinary dislocation reactions for polycrystals. The use of high purity base metals in the present study decreased flow stress in grains and toughened grain boundaries by lowering the grain boundary segregation of harmful elements. The decrease of flow stress also causes toughening because the fracture caused by a given fracture mechanism [12] is suppressed. The macroscopic structure such as the lamellar spacing and the thickness of the second phase plates would not be changed much by the purity of alloy. Fracture strain Fractography showed that the fracture mode in the present study was mainly of the transgranular type, namely, cleavage fracture with rather flat planes and slip plane fracture with saw-toothed zigzag planes due to dislocation reactions. Thus, it seems that the dislocation pile-ups caused transgranular fracture. We consider the fracture stresses in the structures with and without the second phase plates, based on a fracture mechanism due to stress concentration near the front of the dislocation pile-ups. We assume that when the stress concentration reaches a given stress level near the pile-up front, fracture will initiate. As a first approximation, we consider simple cubic grains with and without second phase plates having a side length which corresponds to the grain size being d, the lamellar spacing s and the thickness of the second phase plates w, in which the pile-up lengths are (s - w) and d and the applied stresses o S and 42, in the structures with and without the second phase plates, respectively. Subsequently, the subscripts, 1 and 2, show the structures with and without the second phase plates, respectively. Under %he condition that fracture occurs at the same level of the stress concentration, the following relations are obtained [20). N 1 = (1 - ~ ) ( s
- W)~l/~b
N2 = (1 - ~ ) ( s
- W)Ol(d/(s
= Nl(d/(s
- w)) 1/2
(1) - w))l/2/~b
(2) (2')
When each dislocation sweeps an area (s - w)d and d 2 in the structures with and without the
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second phase plates, respectively, the ratio of shear strains, Yl and Y2' is given as the ratio of the number of pile-up dislocations. YIIY2 = NIIN2" (3) The ratio of the stresses is also given as
o11o2
=
(dl(s
-
w))I12.
(~)
The structural parameters in the present study are s=3 ~m, w=0.7 ~m, d=40 ~m and the applied yield stress, O 0 o~ i=296.0 MPa, while the material parameters are ~=O.247, W=71.4 GPa and b=0.2828 nm, resulting in N1=25 , N2=I06 and y1/Y2=1/4.2. This analysis shows that the presence of the second phase plates, which play the same role as grain boundaries, lowers the fracture strain. This conclusion can partly explain the tendency of the experimental results, that the ductility in the annealed specimen with a larger lamellar spacing is larger than that in the unannealed specimen with a smaller lamellar spacing while the yield stress varies in the opposite manner. Moreover, the above result seems to imply that the matrix of the Ti-48 at % A1 alloy has far larger ductility than that in the composite structure containing the second phase plates. However, TiA1 with an A1 composition of less than 50 at % already contains the second phase [16]. Therefore, we cannot investigate the ductility of only the matrix. Also, the large fracture strain in the Ti-48 at % A1 alloy should be related to the characteristics of the interfaces between Ti3A1 and TiA1 and twins, the second phase and, perhaps, the matrix. Conclusion The mechanical properties of high purity Ti-48 at% AI were studied in tensile tests. The fracture elongation was improved up to about 2.7 %. The improvement of ductility is caused by the lowering of flow stress because of both the use of high purity base metals and the characteristics of the second phase Ti3A1. Acknowledgment This study was supported partly by Keikinzoku Shogakukai (The Light Metal Educational Foundation, Incorporated) and Grant-in-Aid for Scientific Research of the Ministry of Education, Science and Culture, Japan (C62550512). Reference I. 2. 3. 4. 5. 6. 7. 8. 9.
J.B. MacAndrew and H.D. Kessler, J. Metals, 8, 1348 (1956). H.A. Lipsitt, D. Shechtman and R.E. Schafrik, Metall. Trans. 6A, 1991 (1975). P.L. Martin, M.G. Mendiratta and H.A. Lipsitt, Metall. Trans. 14A, 2170 (1983). H.A. Lipsitt, in High Temperature Ordered Intermetallic Alloys, Edited by C.C. Koch, C.T. Liu and N.S. Stoloff, Materials Research Society Symposia Proc. 39, p. 351, (1985). H.R. Ogden, D.J. Maykuth, W.L. Finlay and R.I. Jaffee, J. Metals, 3, 1150 (1951). E.S. Bumps, H.D. Kessler and M. Hansen, J. Metals, 4, 609 (1952). H.R. Ogden, D.J. Maykuth, W.L. Finlay and R.I. Jaffee, J. Metals, 5, 267 (1953). M.J. Blackburn and M.P. Smith, United States Patent No. 4,294,615, Oct. 13, 1981. P.L. Martin, H.A. Lipsitt, N.T. Nuhfer and J.C. Williams, Titanium '80, Science and Technology, Edited by H. Kimura and O. Izumi, Metall. Society of AIME, New York, p. 1245
(1980). 10. T. Tsujimoto, Titanium and Zirconium, 33, 159 (1985) (In Japanese). 11. T. Hashimoto, H. Dohi and Tsujimcto, Preprint of symposium of Japan Institute of Metals, in Plastic Deformation of Ordered Alloys and Intermetallic Compounds (In Japanese), p. 17, July, 1986. 12. T. Kawabata, T. Takezono, T. Kanai and O. Izumi, Acta Metall. 36, 963 (1988). 13. M.J. Blackburn and J.C. Williams, Trans. Met. Soc. AIME, 239, 287 (1967). 14. M.J. Blackburn, in The Science, Technology and Application of Titanium, Edited by R.'Jaffee and N. Promisel, Pergamon Press, New York, p. 633, (1970). 15. H.A. Lipsitt, D. Shechtman and R.E. Schafrik, Metall. Trans. 11A, 1369 (1980). 16. D. Shechtman, M.J. Blackburn and H.A. Lipsitt, Metall. Trans. 5, 1373 (1974). 17. Ye.I. Teytel and E.S. Yakovleva, Fiz. metal, metalloved. 40, No. I, 129 (1975) (in English edition). 18. T. Kawabata and 0. Izumi, Unpublished work. 19. T. Kawabata, T. Kanai and O. Izumi, Acta Metall. 33, 1355 (1985). 20. J.P. HIrth and J. Lothe, in Theory of Dislocations, McGraw-Hill, New York, p. 694, (1968).
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LU n." I.tn
STRAIN
FIG. I
The tensile stress-strain curve of high purity Ti-48 at% AI.
FIG. 3
FIG. 2
A specimen bent by hand. The size of the specimen is 0.4 mm in thickness, 4.2 mm in width and 29.6 mm in length. The radius of the curvature is 9 mm and the tensile strain, Es, at the surface is 2.2 %.
Optical microstructures of (a) an unannealed and (b) an annealed specimen which is heated at 1573 K for 259.2 ks.
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FIG. 4
The gage portion of the tensile specimen deformed at room temperature at strain rate 1.2×10 -~ s-I in air. (a) Low and (b) high magnifications.
FIG. 5
Optical micrographs showing (a) grain boundary cracking, (b) slip bands parallel to the plates of the second phase and (c) wavy slip bands. The arrow and the letter "TA" in (a) are the tensile axis and "c" the grain boundary cracking. The arrow and the letters "s" and "p" in (b) show the second phase plates blocking and parallel to slip bands, respectively. The arrows and the letters "w" and "t1" and "t2" in (c) i l l u s t r a t e the w a v y slip bands and the slip bands r u n n i n g across the second phases, respectively.