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Crack propagation behavior in TiAl–Nb single and Bi-PST crystals
Intermetallics 15 (2007) 976e984 www.elsevier.com/locate/intermet
Crack propagation behavior in TiAleNb single and Bi-PST crystals S.W. Kim a,*, K.S. Kumar b, M.H. Oh c, D.M. Wee a a
Department of Materials Science and Engineering, KAIST, Daejeon 305-701, Republic of Korea b Division of Engineering, Brown University, Providence, RI 02912, USA c Department of Materials Science and Engineering, KIT, Gumi 730-701, Republic of Korea
Received 25 July 2006; received in revised form 1 December 2006; accepted 11 December 2006 Available online 6 March 2007
Abstract The fracture toughness of directional solidified Tie(45,47)Ale3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie(45,47)Ale3Nbe0.3Sie0.2C type I alloys and their contribution to crack growth resistance of TiAleNb alloys were studied using PST (polysynthetically twinned) crystals produced by directional solidification in FZ (floating zone) furnace. Lamellar orientations in the individual colonies are described using two angles defined with respect to the notch orientation: an in-plane kink angle and a through-thickness twist angle. Therefore, lamellar misorientation across an individual colony boundary is quantified as differences in these angles across the boundary. Crack growth resistance in colony boundary was identified by three-point bend test and crack advance was monitored by interrupted in situ test. From three-point bend test, it was found that the colony boundary could offer significant resistance to crack growth under large twist angle difference. Fracture toughness of type I specimens (in which crack propagates against lamellae boundaries) of the alloys decreased slightly with increasing Si and C contents and increased rapidly with decreasing Al content. The toughness for type I specimens was controlled by a2ea2 spacing in which the delamination-type separation occurred. Compared to 47Al alloys, a2ea2 spacing in 45Al alloys increased by decreasing Al content, therefore, fracture toughness increased rapidly. These results are discussed and the ability to improve toughness by changing Al content, Si and C addition in TiAleNb alloys produced by directional solidification is suggested. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; C. Crystal growth; B. Fatigue resistance and crack growth; F. Mechanical testing
1. Introduction Recently, there has been much interest in the research and development of titanium aluminides, particularly for lightweight high-temperature structural applications [1e4]. Furthermore, fully lamellar TiAl alloys composed of TiAl (g) and a small amount of Ti3Al (a2) exhibit superior fracture toughness and creep resistance, although compared to conventional hightemperature alloys, the room temperature ductility of TiAl alloys with the lamellar microstructure is poor. Fracture toughness of two-phase TiAl alloys has been known to increase with increasing colony size [5,6], on the other hand, the ductility decreases with increasing colony size [5]. Such inverse relation
* Corresponding author. Tel.: þ82 16 704 8564; fax: þ82 42 869 4275. E-mail address: [email protected] (S.W. Kim). 0966-9795/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.12.004
between tensile properties and fracture toughness is one of the key deficiencies of the current properties of two-phase TiAl alloys [7,8]. One approach for achieving balanced mechanical properties is by directional solidification (DS) techniques [3,5,6]. The mechanical properties of two-phase lamellar TiAl alloys are known to be dependent [9] on the lamellar orientation relative to the testing axis; thus it is essential to be able to control the orientation during directional solidification in order to obtain desirable properties. We have succeeded in obtaining TiAl alloys with aligned lamellar orientation using Tie46Ale1.5Moe0.2C, Tie46Ale1.5Moe1Si and Tie44.6Ale3Nbe0.6Sie0.2C alloys as seed materials [10e12]. Experimental investigations relating to crack propagation in lamellar TiAl-based alloys using both single colony (PST crystals) and polycolony specimens have demonstrated that intrinsic and extrinsic toughening mechanisms can be present in
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these alloys [13e21]. It has been reported that much of the fracture resistance in the lamellar structure arises from crack bridging by intact ligaments in the wake of the crack [21]. Recently, it was shown that colony boundaries can offer resistance to crack growth and that the magnitude of the resistance is dependent on the lamellar misorientation across the boundary [21]. In TiAl PST crystals, the fracture behavior was reported using type I (crack arrester), type II (crack divider) and type III (crack delamination) specimens [20]. The crack arrester type (type I) TiAl PST crystals showed excellent crack propagating resistance [21]. In this paper, the fracture toughness of Tie(45,47Al)e3Nb alloys with type I orientation and Bi-PST crystals containing Si and C was investigated. We previously reported that the fracture toughness of Tie47Ale3Nb alloys decreased with increasing Si and C contents [22] due to decreasing a-phase volume fraction. Therefore, it is focused to suggest the possibility of improving toughness by changing Al content, Si and C addition in TiAleNb alloys with type I orientation. Moreover, improving crack growth resistance of TiAleNb Bi-PST crystals is suggested. 2. Experimental procedure The alloy used in this research was produced from 99.99 wt.% Ti, 99.99 wt.% Al, 99.9 wt.% Nb, 99 wt.% TiC powder and 99.99 wt.% Si. Button ingots (Tie(45,47)Ale 3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie(45,47)Ale3Nbe 0.3Sie0.2C) were made by vacuum arc melting in an argon
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atmosphere and the resulting ingot was remelted at least five times or more to promote homogeneity. The feeder ingots for the DS process were obtained by re-melting the arc melting buttons into 14 mm 100 mm cylinders. The DS ingots were grown in an ASGAL FZ-SS35W optical floating zone furnace using Tie44.5Ale3Nbe0.6Sie0.2C alloy as a seed material at a growth rate of 5 mm/h for the first 10 mm, and subsequently at 10 mm/h. The microstructure of the alloys was examined optically as well as using scanning and transmission electron microscopy techniques (SEM and TEM e JSM 845 and Philips EM 420, respectively). The samples for TEM observation were prepared by twin-jet electropolishing using a solution of 925 ml methanol and 75 ml perchloric acid in the 20 C to 40 C range and a voltage of 12e15 V. Fracture surfaces were examined by SEM. Room temperature tensile tests were performed at a nominal strain rate of 2 104/s using a screw driven universal testing machine. Flat tensile specimens were electro-discharge machined with the tensile axis parallel to the growth direction, and each specimen had a 5 mm gauge length with a 2 mm 1 mm cross-section. All the tensile specimens were polished through 2000 grit emery paper and electropolished in a solution of 5% HClO4e35% n-butanole60% methanol by volume at 65 C. Fracture toughness at room temperature was obtained by using single edge notched bend (SENB) specimens. These specimens were electro-discharge machined from DS ingots with dimensions 20 mm 4 mm 3 mm (Fig. 1). For each composition, two ‘‘type I’’ specimens were prepared. Prior to performing the fracture toughness tests, the specimens were
Fig. 1. The schematic illustration of (a) type I specimen, (b) Bi-PST specimen and (c) the parameters a, b and q describing the lamellar orientation with respect to the notch and the colony boundary inclination [26].
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Fig. 2. Room temperature loadedisplacement curves from the three-point bend tests of (a) Tie47Ale3NbexSieyC (x ¼ 0, 0.2, 0.3, y ¼ 0, 0.1, 0.2) alloys and (b) Tie45Ale3NbexSieyC (x ¼ 0, 0.2, 0.3, y ¼ 0, 0.1, 0.2) alloys.
subjected to compressionecompression fatigue to generate a precrack at the notch tip. The compression load of fatigue test to generate crack was under 10 kN and the load ratio (S ) was 0.1. This means that the ratio of maximum and minimum load was 0.1. Moreover, the crack length, which was generated by compression fatigue test, was usually under 0.05 mm. The precracks were examined under an optical microscope to ensure that they were present on both sides of the specimens. The three-point bend tests were carried out using an Instron machine at room temperature in air using a crosshead speed of 0.01 mm/min and loadedisplacement data were acquired electronically during the test. Fracture toughness was obtained from loadedisplacement data. We defined PQ by 5% secant line instead of Pmax. 3. Results and discussion 3.1. Fracture toughness of type I specimens In our previous study, it was reported that the fracture toughness of Tie47Ale3Nb alloys decreased with increasing Si and C contents, besides, yield strength increased rapidly with increasing Si and C contents [22]. The fracture toughness of Tie47Ale3Nb alloys decreased due to decreasing a2 volume fraction which contributes to the toughness of type I specimens. In this study, the possibility of increasing toughness was investigated by changing Al content to 45 at.% to increase a2 volume fraction. Moreover, important factors that affect the toughness of type I specimens were discussed. As shown in Fig. 1, the type I specimens manufactured by DS
(directionally solidification) process were used for three-point bend test. Before the bend test, fatigue precracks were generated at the notch tip by compressionecompression fatigue test. Moreover, two specimens were prepared for bend test in same alloys, and the bend test results were very similar in those alloys. Loadedisplacement curves obtained from the bend test are shown in Fig. 2. As mentioned above, we previously did the bend test for Tie47Ale3Nb alloys [22], but the results shown in Fig. 2a were obtained newly for this study except for Tie 47Ale3Nbe0.3Sie0.2C alloy. In all the cases, the load dropped rapidly in alloys with Si and C, especially, dropping is more rapid in Tie45Ale3Nb alloys. Moreover, in the alloys with same Si and C contents, the maximum load is higher in Tie45Ale3Nb alloys than in Tie47Ale3Nb alloys. The fracture toughness (KQ) values for the alloys are presented in Table 1. Using standard methods and equations for three-point notched bend specimens [23], the fracture toughness KQ was determined in each case, then, using the plane-strain requirement it was validated as KIC. Only two KQ values (Tie47 Ale3Nbe0.3Sie0.2C, Tie45Ale3Nbe0.3Sie0.2C) satisfied the requirement. Fracture toughness (KQ) of Tie45Ale3Nb alloy was the highest of the alloys examined in this study. The reason for this result will be discussed later. In Table 1, the yield strengths of the alloys obtained from tensile tests are also shown. As mentioned in Ref. [22], the yield strength increased rapidly with increasing Si and C contents in Tie 47Ale3Nb alloys due to solid solution hardening and lamellar refinement. However, the yield strength of Tie45Ale3Nb alloys increased not so rapidly with increasing Si and C
Table 1 Summary of fracture toughness (KQ) and yield strength (sy.s.) vs microstructural factors Comp.
KQ sy.s. a vol.(%) a2ea2 (mm) g lamellae size (nm)
Tie47Ale3Nb
Tie45A1e3Nb
0Sie0C
0.2Sie0.1C
0.3Sie0.2C
0Sie0C
0.2Sie0.1C
0.3Sie0.2C
18.9 410 15.4 2.55 441.9
17.1 530 10.8 2.88 431.6
16.3 630 9.5 3.86 418.2
24.7 614 28.1 0.68 360.3
20.2 663 21.9 1.22 354.8
18.5 690 20.3 2.48 340.4
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Fig. 3. Schematic illustration of crack propagating process in type I alloy.
contents. From Fig. 2 and Table 1, the fracture toughness can be improved by decreasing Si and C contents and changing Al content from 47 to 45 at.%. 3.2. Controlling factors that affect the fracture toughness Fig. 3 shows the crack propagation mode in type I specimens [20]. When the sharp crack generated by fatigue precracking meets the lamellae boundaries (g/g and/or a2/g) and a2 lamellae, delamination occurred parallel to the boundaries or a2 lamellae. Therefore, the crack tip becomes blunt and propagation across g lamellae will be very hard. These blunting can contribute to the high fracture toughness of type I specimens and if such blunting occurred a lot of frequencies, fracture toughness will be maximized. Therefore, we can suggest the controlling factors of fracture toughness in type I alloys: a2 volume fraction and g lamellae thickness. Delamination will occur many times when a2 volume fraction is high and g lamellae are fine. In Table 1, a2 volume fraction and g lamellae thickness are summarized. The a2 volume fraction of alloys is calculated considering the shifting phase diagram effects of Nb, Si and C [22] and average g lamellae thickness is measured by TEM observation. a2 Volume fraction becomes smaller with high Si and C contents, and those of 45Al alloys are higher than those of 47Al alloys. Moreover, g lamellae thickness becomes slightly smaller with high Si and C additions and those of 45Al alloys are finer than those of 47Al alloys. Therefore, it is thought to be reasonable that fracture toughness of 45Al alloys is higher
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than 47Al alloys. However, among the alloys with same Al content, fracture toughness decreased with higher Si and C contents especially in 45Al alloys even though g lamellae thickness become finer. In the viewpoint of a2 volume fraction, fracture toughness increased with increasing a2 volume fractions compared to 45Al and 47Al alloys approximately. However, the fracture toughness of Tie45Ale3Nbe0.3Sie0.2C alloy increased slightly more than Tie47Ale3Nbe0.3Sie0.2C alloy even though a2 volume fraction is more than twice in latter alloy. Moreover, comparing the toughness of Tie47Ale3Nb and Tie45Ale3Nbe0.3Sie0.2C alloys, the fracture toughness of latter alloys slightly decreased even though the a2 volume fraction of Tie45Ale3Nbe0.3Sie0.2C alloy is higher than that of Tie47Ale3Nb alloy. Therefore, we concentrated on the a2 lamellae distribution rather than a2 volume fraction. a2 lamellae distributions of Tie47Ale3Nb, Tie47Ale3Nbe0.3Sie0.2C and Tie45Ale3Nbe0.3Sie0.2C alloys by TEM dark field images are shown in Fig. 4. The white lines in Fig. 4 indicate a2 lamellae. Comparing Tie47Ale3Nb and Tie45Ale3Nbe 0.3Sie0.2C alloys (Fig. 4a and c), a2 lamellae distributions are very similar even though a2 volume fraction is different for those alloys (notice that the magnification is not the same.). Moreover, a2 lamellae thickness of Tie45Ale3Nbe 0.3Sie0.2C alloy is approximately twice than that of Tie47 Ale3Nbe0.3Sie0.2C alloy so that a2 lamellae distributions of those alloys are not very different. The a2ea2 spacing of the alloys in this study is summarized in Table 1. a2ea2 spacing of Tie45Ale3Nb alloys of which a2 volume fraction is highest was most fine among those alloys, and that of Tie 47Ale3Nbe0.3Sie0.2C alloy was largest. The relationship between a2ea2 spacing and fracture toughness is shown in Fig. 5. In both 45Al and 47Al alloys, the two factors have linear relationship. Therefore, we can suggest that a2ea2 spacing is the most important factor that controls fracture toughness (KQ). However, the a2ea2 spacings in Fig. 5 are average numbers for alloys. We should notice that a2ea2 spacing in an alloy is distributed for a long range. Fig. 6 shows that the distributions of a2ea2 spacing of Tie47Ale3Nbe0.3Sie0.2C alloy.
Fig. 4. TEM dark field image of (a) Tie47Ale3Nb, (b) Tie47Ale3Nbe0.3Sie0.2C and (c) Tie45Ale3Nbe0.3Sie0.2C showing the distribution of a2 lamellae.
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Fig. 5. Relation between fracture toughness and a2ea2 spacing.
As we can see in Fig. 6, a2ea2 spacing is distributed for a long range of order. Therefore, the crack propagating resistance is not uniform during bend test and the maximum peak considered indicates the total resistance. 3.3. Modeling of crack propagating mode in type I specimens To understand the fracture toughness change with a2ea2 spacing, crack propagation sequence and paths should be examined. Hwu et al. examined the crack propagation behavior using metal/ceramic laminate systems [24,25]. They investigated the transition from single crack extension to multiple cracking near the crack tip over a range of metal/ceramic thickness ratios (tm/tc). The transition was controlled by brittle ceramic thickness; crack reinitiated with multiple cracking when the ceramic layer was thin, however, crack reinitiated with macroscopic crack propagation when the ceramic layer was thick. Macroscopic crack propagation showed higher
toughness because load cannot be concentrated at the crack tip. Considering this report and delamination in a2 phases, crack propagation behavior in type I specimens with small and high a2ea2 spacing is shown in Figs. 7 and 8. In the case of small a2ea2 spacing (Fig. 7), delamination will occur when the sharp crack meets the boundary (mainly a2/g), and high load will be needed for further propagation. Next, delamination in the next a2 phase will occur instead of fracture in the middle g. Therefore, it can be very hard to break middle g phase because load cannot be concentrated at the crack tip. In the case of high a2e a2 spacing (Fig. 8), delamination will occur when the sharp crack meets the boundary (mainly a2/g), and high load will be needed for further propagation (same with Fig. 7). However, middle g phases will break instead of delamination in next a2 phase because the next a2 phase is located out of the stress field. Therefore, low load will be needed to break middle g phases because load can be concentrated at the crack tip. Figs. 7 and 8 can explain the result of Fig. 2 and Table 1.
Fig. 6. The distribution of a2ea2 spacing of Tie47Ale3Nbe0.3Sie0.2C alloy.
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Fig. 7. Schematic illustration of crack propagating process in type I alloy with small a2ea2 spacing.
3.4. Crack initiation in Bi-PST crystals with different twist angle and compositions In this paper, crack propagation behavior in Bi-PST crystals was also investigated. To understand the crack propagation behavior, Bi-PST crystals with different twist angles (Fig. 1b and c) and compositions were prepared by DS process in FZ furnace. Table 2 illustrates the kink (a) and twist (b) angles of specimens in this paper. As shown in Table 2, three specimens of Tie45Al and one Tie45Ale3Nb specimen have similar twist angle differences (Db), however, that of Tie45Ale3Nbe 0.2Sie0.1C alloy was very different. Before the bend test, fatigue precracks were generated at the notch tip by compressionecompression fatigue test. All the precracks were examined by OM (optical microscope) and they were made parallel with lamellar boundaries. Loadedisplacement curves obtained from the bend test are shown in Fig. 9. As shown in Fig. 9, the maximum load of
Tie45Al was very similar (523e576 N). Considering they have similar Db and different Da, the toughness can be controlled by Db rather than Da as already reported by Ping et al. [26]. In case of Tie45Ale3Nb alloy, maximum load was higher than those of Tie45Al alloys in spite of similar Db. This was caused by different a2ea2 spacings in which delamination occur (it will be discussed later). The maximum load of Tie45Ale3Nbe0.2Sie0.1C alloy was very low compared to other compositions because of small Db. In the case of Tie45Ale3Nbe0.2Sie0.1C alloy, the kink angle of colony 1 was very high (45 ), therefore, more than two precracks were generated and propagated (Fig. 10). From Figs. 9 and 10, toughness of Bi-PST crystals was controlled by Db. Interrupted in situ test was performed to understand these results and the crack path was examined. Interrupted in situ test was performed by slightly loading using loading jig and observing using OM and SEM (BS images). Fig. 11 shows the cracks after 1st (Fig. 11a) and 2nd (Fig. 11b) loadings.
Fig. 8. Schematic illustration of crack propagating process in type I alloy with high a2ea2 spacing.
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982 Table 2 Kink and twist angles of Bi-PST crystals Specimen
a ( )
b ( )
Colony 1
Colony 2
Colony 1
Colony 2
45Al #1 45Al #2 45Al #3 45Ale3Nb 45Ale3Nbe0.2Sie0.1C
35 35 18 15 45
5 5 43 36 35
10 10 <5 <10 25
48 48 45 45 45
From Fig. 11, it was inferred that one of the precracks propagated after 1st loading (Fig. 11a) and stopped, then, the other crack propagated after 2nd loading (Fig. 11b). Fig. 12 shows the cracks after 3rd loading by OM (Fig. 12a) and SEM BS images (Fig. 12b). Main crack had already crossed boundary and propagation mode was changed from delamination (colony 1) to translamellar (colony 2). It was very noticeable that so many delaminations occurred in colony 2 (Fig. 12a) and these
Fig. 9. Room temperature loadedisplacement curves from the three-point bend tests of Tie45Al, Tie45Ale3Nb and Tie45Ale3Nbe0.2Sie0.1C alloys.
delaminations occurred mainly in a2/g lamellar boundaries and a2 phases (Fig. 12b). Yokoshima and Yamaguchi already reported that delamination cannot give any resistance for crack propagation and translamellar type propagation can give a lot
Fig. 10. Microstructures after three-point bend test of Tie45Ale3Nbe0.2Sie0.1C alloy.
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Fig. 11. Microstructures of Tie45Al alloy during interrupted in situ test after (a) first and (b) second loadings.
of resistance [20]. Therefore, the change from delamination to translamellar at the boundary can be explained as that twisting needs higher energy than translamellar. So, we can guess that crack will propagate by delamination after boundary when Db is small (Fig. 10). Moreover, if the delamination in colony 2 occurred more frequently (small a2ea2 spacing in colony 2), the maximum load after three-point bend test can be higher as in the case of Tie45Ale3Nb alloy. The average a2ea2 spacings of Tie45Al and Tie45Ale3Nb obtained by SEM BS images were 1.11 and 0.68 mm, respectively. Therefore, more delamination occurred in Tie45Ale3Nb alloy so that load was dispersed and maximum load was higher than Tie45Al alloys (Fig. 9). Therefore, crack resistance at the colony boundaries was controlled mainly by Db and secondly by a2ea2 spacing in colony 2.
4. Conclusions The crack propagation behavior of directional solidified Tie(45,47)Ale3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie (45,47)Ale3Nbe0.3Sie0.2C type I specimens and Tie45Al, Tie45Ale3Nb, Tie45Ale3Nbe0.2Sie0.1C Bi-PST crystals was investigated. It was found that fracture toughness of type I specimens increased with decreasing Al contents and decreased with increasing Si and C contents. The toughness of type I specimens was controlled by a2ea2 spacing. Moreover, crack propagation resistance of Bi-PST crystals was controlled by twist angle differences (Db) rather than kink angle differences (Da). The crack paths of Bi-PST crystals changed from delamination (colony 1) to translamellar (colony 2) with high Db, and did not change with low Db.
Fig. 12. Microstructures of Tie45Al alloy during interrupted in situ test after third loading by (a) OM and (b) SEM BS images.
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Acknowledgements This work was supported by Korea Research Foundation Grant (KRF-2004-042-D00111). References [1] Kim YW, Dimiduk DM. J Met 1991;43:40. [2] Huang SC, Shih DS. In: Kim YW, Boyer RR, editors. Microstructuree property correlation in TiAl-base alloys. Warrendale, PA: TMS; 1991. p. 105. [3] Yamaguchi M, Inui H. TiAl compounds for structural applications. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 127. [4] Kim YW. J Met 1994;46:30. [5] Chan KS, Kim YW. Acta Metall Mater 1995;43:439. [6] Chan KS, Kim YW. Metall Trans 1992;23A:1663. [7] London B, Larsen DK, Wheeler DA, Aimone PR. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 151. [8] Takeyama M, Kumagai T, Nakamura M, Kikuchi M. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 167.
[9] Inui H, Oh MH, Nakamura A, Yamaguchi M. Acta Mater 1992;40:3095. [10] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Acta Mater 2000;48:3221. [11] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Mater Sci Eng 2002;C329e331:19. [12] Kim JH, Kim SW, Oh MH, Inui H, Wee DM. Intermetallics 2005; 13:1038. [13] Chan KS. Metall Trans 1993;24A:569. [14] Chan KS, Kim YW. Metall Mater Trans 1994;25A:1217. [15] Appel F, Christoph U, Wagner R. Philos Mag A 1995;72:341. [16] Kim YW. Mater Sci Eng 1995;A192e193:519. [17] Chen M, Lin D, Chen D. Scr Mater 1997;36:497. [18] Chan KS, Shih DS. Metall Mater Trans 1997;28A:79. [19] Mitao S, Isawa T, Tsuyama S. Scr Metall Mater 1992;26:1405. [20] Yokoshima S, Yamaguchi M. Acta Mater 1996;44:873. [21] Chan KS, Onstott J, Kumar KS. Mater Trans A 2000;31:71. [22] Kim SW, Wang P, Oh MH, Wee DM, Kumar KS. Intermetallics 2004; 12:499. [23] Anderson TL. In: Fracture mechanics, fundamentals and applications. 2nd ed. New York: CRC Press; 1995. p. 603. [24] Hwu KL, Derby B. Acta Mater 1999;47:529. [25] Hwu KL, Derby B. Acta Mater 1999;47:545. [26] Wang P, Bhate N, Chan KS, Kumar KS. Acta Mater 2003;51:1573.
Crack propagation behavior in TiAleNb single and Bi-PST crystals S.W. Kim a,*, K.S. Kumar b, M.H. Oh c, D.M. Wee a a
Department of Materials Science and Engineering, KAIST, Daejeon 305-701, Republic of Korea b Division of Engineering, Brown University, Providence, RI 02912, USA c Department of Materials Science and Engineering, KIT, Gumi 730-701, Republic of Korea
Received 25 July 2006; received in revised form 1 December 2006; accepted 11 December 2006 Available online 6 March 2007
Abstract The fracture toughness of directional solidified Tie(45,47)Ale3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie(45,47)Ale3Nbe0.3Sie0.2C type I alloys and their contribution to crack growth resistance of TiAleNb alloys were studied using PST (polysynthetically twinned) crystals produced by directional solidification in FZ (floating zone) furnace. Lamellar orientations in the individual colonies are described using two angles defined with respect to the notch orientation: an in-plane kink angle and a through-thickness twist angle. Therefore, lamellar misorientation across an individual colony boundary is quantified as differences in these angles across the boundary. Crack growth resistance in colony boundary was identified by three-point bend test and crack advance was monitored by interrupted in situ test. From three-point bend test, it was found that the colony boundary could offer significant resistance to crack growth under large twist angle difference. Fracture toughness of type I specimens (in which crack propagates against lamellae boundaries) of the alloys decreased slightly with increasing Si and C contents and increased rapidly with decreasing Al content. The toughness for type I specimens was controlled by a2ea2 spacing in which the delamination-type separation occurred. Compared to 47Al alloys, a2ea2 spacing in 45Al alloys increased by decreasing Al content, therefore, fracture toughness increased rapidly. These results are discussed and the ability to improve toughness by changing Al content, Si and C addition in TiAleNb alloys produced by directional solidification is suggested. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; C. Crystal growth; B. Fatigue resistance and crack growth; F. Mechanical testing
1. Introduction Recently, there has been much interest in the research and development of titanium aluminides, particularly for lightweight high-temperature structural applications [1e4]. Furthermore, fully lamellar TiAl alloys composed of TiAl (g) and a small amount of Ti3Al (a2) exhibit superior fracture toughness and creep resistance, although compared to conventional hightemperature alloys, the room temperature ductility of TiAl alloys with the lamellar microstructure is poor. Fracture toughness of two-phase TiAl alloys has been known to increase with increasing colony size [5,6], on the other hand, the ductility decreases with increasing colony size [5]. Such inverse relation
* Corresponding author. Tel.: þ82 16 704 8564; fax: þ82 42 869 4275. E-mail address: [email protected] (S.W. Kim). 0966-9795/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.12.004
between tensile properties and fracture toughness is one of the key deficiencies of the current properties of two-phase TiAl alloys [7,8]. One approach for achieving balanced mechanical properties is by directional solidification (DS) techniques [3,5,6]. The mechanical properties of two-phase lamellar TiAl alloys are known to be dependent [9] on the lamellar orientation relative to the testing axis; thus it is essential to be able to control the orientation during directional solidification in order to obtain desirable properties. We have succeeded in obtaining TiAl alloys with aligned lamellar orientation using Tie46Ale1.5Moe0.2C, Tie46Ale1.5Moe1Si and Tie44.6Ale3Nbe0.6Sie0.2C alloys as seed materials [10e12]. Experimental investigations relating to crack propagation in lamellar TiAl-based alloys using both single colony (PST crystals) and polycolony specimens have demonstrated that intrinsic and extrinsic toughening mechanisms can be present in
S.W. Kim et al. / Intermetallics 15 (2007) 976e984
these alloys [13e21]. It has been reported that much of the fracture resistance in the lamellar structure arises from crack bridging by intact ligaments in the wake of the crack [21]. Recently, it was shown that colony boundaries can offer resistance to crack growth and that the magnitude of the resistance is dependent on the lamellar misorientation across the boundary [21]. In TiAl PST crystals, the fracture behavior was reported using type I (crack arrester), type II (crack divider) and type III (crack delamination) specimens [20]. The crack arrester type (type I) TiAl PST crystals showed excellent crack propagating resistance [21]. In this paper, the fracture toughness of Tie(45,47Al)e3Nb alloys with type I orientation and Bi-PST crystals containing Si and C was investigated. We previously reported that the fracture toughness of Tie47Ale3Nb alloys decreased with increasing Si and C contents [22] due to decreasing a-phase volume fraction. Therefore, it is focused to suggest the possibility of improving toughness by changing Al content, Si and C addition in TiAleNb alloys with type I orientation. Moreover, improving crack growth resistance of TiAleNb Bi-PST crystals is suggested. 2. Experimental procedure The alloy used in this research was produced from 99.99 wt.% Ti, 99.99 wt.% Al, 99.9 wt.% Nb, 99 wt.% TiC powder and 99.99 wt.% Si. Button ingots (Tie(45,47)Ale 3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie(45,47)Ale3Nbe 0.3Sie0.2C) were made by vacuum arc melting in an argon
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atmosphere and the resulting ingot was remelted at least five times or more to promote homogeneity. The feeder ingots for the DS process were obtained by re-melting the arc melting buttons into 14 mm 100 mm cylinders. The DS ingots were grown in an ASGAL FZ-SS35W optical floating zone furnace using Tie44.5Ale3Nbe0.6Sie0.2C alloy as a seed material at a growth rate of 5 mm/h for the first 10 mm, and subsequently at 10 mm/h. The microstructure of the alloys was examined optically as well as using scanning and transmission electron microscopy techniques (SEM and TEM e JSM 845 and Philips EM 420, respectively). The samples for TEM observation were prepared by twin-jet electropolishing using a solution of 925 ml methanol and 75 ml perchloric acid in the 20 C to 40 C range and a voltage of 12e15 V. Fracture surfaces were examined by SEM. Room temperature tensile tests were performed at a nominal strain rate of 2 104/s using a screw driven universal testing machine. Flat tensile specimens were electro-discharge machined with the tensile axis parallel to the growth direction, and each specimen had a 5 mm gauge length with a 2 mm 1 mm cross-section. All the tensile specimens were polished through 2000 grit emery paper and electropolished in a solution of 5% HClO4e35% n-butanole60% methanol by volume at 65 C. Fracture toughness at room temperature was obtained by using single edge notched bend (SENB) specimens. These specimens were electro-discharge machined from DS ingots with dimensions 20 mm 4 mm 3 mm (Fig. 1). For each composition, two ‘‘type I’’ specimens were prepared. Prior to performing the fracture toughness tests, the specimens were
Fig. 1. The schematic illustration of (a) type I specimen, (b) Bi-PST specimen and (c) the parameters a, b and q describing the lamellar orientation with respect to the notch and the colony boundary inclination [26].
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Fig. 2. Room temperature loadedisplacement curves from the three-point bend tests of (a) Tie47Ale3NbexSieyC (x ¼ 0, 0.2, 0.3, y ¼ 0, 0.1, 0.2) alloys and (b) Tie45Ale3NbexSieyC (x ¼ 0, 0.2, 0.3, y ¼ 0, 0.1, 0.2) alloys.
subjected to compressionecompression fatigue to generate a precrack at the notch tip. The compression load of fatigue test to generate crack was under 10 kN and the load ratio (S ) was 0.1. This means that the ratio of maximum and minimum load was 0.1. Moreover, the crack length, which was generated by compression fatigue test, was usually under 0.05 mm. The precracks were examined under an optical microscope to ensure that they were present on both sides of the specimens. The three-point bend tests were carried out using an Instron machine at room temperature in air using a crosshead speed of 0.01 mm/min and loadedisplacement data were acquired electronically during the test. Fracture toughness was obtained from loadedisplacement data. We defined PQ by 5% secant line instead of Pmax. 3. Results and discussion 3.1. Fracture toughness of type I specimens In our previous study, it was reported that the fracture toughness of Tie47Ale3Nb alloys decreased with increasing Si and C contents, besides, yield strength increased rapidly with increasing Si and C contents [22]. The fracture toughness of Tie47Ale3Nb alloys decreased due to decreasing a2 volume fraction which contributes to the toughness of type I specimens. In this study, the possibility of increasing toughness was investigated by changing Al content to 45 at.% to increase a2 volume fraction. Moreover, important factors that affect the toughness of type I specimens were discussed. As shown in Fig. 1, the type I specimens manufactured by DS
(directionally solidification) process were used for three-point bend test. Before the bend test, fatigue precracks were generated at the notch tip by compressionecompression fatigue test. Moreover, two specimens were prepared for bend test in same alloys, and the bend test results were very similar in those alloys. Loadedisplacement curves obtained from the bend test are shown in Fig. 2. As mentioned above, we previously did the bend test for Tie47Ale3Nb alloys [22], but the results shown in Fig. 2a were obtained newly for this study except for Tie 47Ale3Nbe0.3Sie0.2C alloy. In all the cases, the load dropped rapidly in alloys with Si and C, especially, dropping is more rapid in Tie45Ale3Nb alloys. Moreover, in the alloys with same Si and C contents, the maximum load is higher in Tie45Ale3Nb alloys than in Tie47Ale3Nb alloys. The fracture toughness (KQ) values for the alloys are presented in Table 1. Using standard methods and equations for three-point notched bend specimens [23], the fracture toughness KQ was determined in each case, then, using the plane-strain requirement it was validated as KIC. Only two KQ values (Tie47 Ale3Nbe0.3Sie0.2C, Tie45Ale3Nbe0.3Sie0.2C) satisfied the requirement. Fracture toughness (KQ) of Tie45Ale3Nb alloy was the highest of the alloys examined in this study. The reason for this result will be discussed later. In Table 1, the yield strengths of the alloys obtained from tensile tests are also shown. As mentioned in Ref. [22], the yield strength increased rapidly with increasing Si and C contents in Tie 47Ale3Nb alloys due to solid solution hardening and lamellar refinement. However, the yield strength of Tie45Ale3Nb alloys increased not so rapidly with increasing Si and C
Table 1 Summary of fracture toughness (KQ) and yield strength (sy.s.) vs microstructural factors Comp.
KQ sy.s. a vol.(%) a2ea2 (mm) g lamellae size (nm)
Tie47Ale3Nb
Tie45A1e3Nb
0Sie0C
0.2Sie0.1C
0.3Sie0.2C
0Sie0C
0.2Sie0.1C
0.3Sie0.2C
18.9 410 15.4 2.55 441.9
17.1 530 10.8 2.88 431.6
16.3 630 9.5 3.86 418.2
24.7 614 28.1 0.68 360.3
20.2 663 21.9 1.22 354.8
18.5 690 20.3 2.48 340.4
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Fig. 3. Schematic illustration of crack propagating process in type I alloy.
contents. From Fig. 2 and Table 1, the fracture toughness can be improved by decreasing Si and C contents and changing Al content from 47 to 45 at.%. 3.2. Controlling factors that affect the fracture toughness Fig. 3 shows the crack propagation mode in type I specimens [20]. When the sharp crack generated by fatigue precracking meets the lamellae boundaries (g/g and/or a2/g) and a2 lamellae, delamination occurred parallel to the boundaries or a2 lamellae. Therefore, the crack tip becomes blunt and propagation across g lamellae will be very hard. These blunting can contribute to the high fracture toughness of type I specimens and if such blunting occurred a lot of frequencies, fracture toughness will be maximized. Therefore, we can suggest the controlling factors of fracture toughness in type I alloys: a2 volume fraction and g lamellae thickness. Delamination will occur many times when a2 volume fraction is high and g lamellae are fine. In Table 1, a2 volume fraction and g lamellae thickness are summarized. The a2 volume fraction of alloys is calculated considering the shifting phase diagram effects of Nb, Si and C [22] and average g lamellae thickness is measured by TEM observation. a2 Volume fraction becomes smaller with high Si and C contents, and those of 45Al alloys are higher than those of 47Al alloys. Moreover, g lamellae thickness becomes slightly smaller with high Si and C additions and those of 45Al alloys are finer than those of 47Al alloys. Therefore, it is thought to be reasonable that fracture toughness of 45Al alloys is higher
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than 47Al alloys. However, among the alloys with same Al content, fracture toughness decreased with higher Si and C contents especially in 45Al alloys even though g lamellae thickness become finer. In the viewpoint of a2 volume fraction, fracture toughness increased with increasing a2 volume fractions compared to 45Al and 47Al alloys approximately. However, the fracture toughness of Tie45Ale3Nbe0.3Sie0.2C alloy increased slightly more than Tie47Ale3Nbe0.3Sie0.2C alloy even though a2 volume fraction is more than twice in latter alloy. Moreover, comparing the toughness of Tie47Ale3Nb and Tie45Ale3Nbe0.3Sie0.2C alloys, the fracture toughness of latter alloys slightly decreased even though the a2 volume fraction of Tie45Ale3Nbe0.3Sie0.2C alloy is higher than that of Tie47Ale3Nb alloy. Therefore, we concentrated on the a2 lamellae distribution rather than a2 volume fraction. a2 lamellae distributions of Tie47Ale3Nb, Tie47Ale3Nbe0.3Sie0.2C and Tie45Ale3Nbe0.3Sie0.2C alloys by TEM dark field images are shown in Fig. 4. The white lines in Fig. 4 indicate a2 lamellae. Comparing Tie47Ale3Nb and Tie45Ale3Nbe 0.3Sie0.2C alloys (Fig. 4a and c), a2 lamellae distributions are very similar even though a2 volume fraction is different for those alloys (notice that the magnification is not the same.). Moreover, a2 lamellae thickness of Tie45Ale3Nbe 0.3Sie0.2C alloy is approximately twice than that of Tie47 Ale3Nbe0.3Sie0.2C alloy so that a2 lamellae distributions of those alloys are not very different. The a2ea2 spacing of the alloys in this study is summarized in Table 1. a2ea2 spacing of Tie45Ale3Nb alloys of which a2 volume fraction is highest was most fine among those alloys, and that of Tie 47Ale3Nbe0.3Sie0.2C alloy was largest. The relationship between a2ea2 spacing and fracture toughness is shown in Fig. 5. In both 45Al and 47Al alloys, the two factors have linear relationship. Therefore, we can suggest that a2ea2 spacing is the most important factor that controls fracture toughness (KQ). However, the a2ea2 spacings in Fig. 5 are average numbers for alloys. We should notice that a2ea2 spacing in an alloy is distributed for a long range. Fig. 6 shows that the distributions of a2ea2 spacing of Tie47Ale3Nbe0.3Sie0.2C alloy.
Fig. 4. TEM dark field image of (a) Tie47Ale3Nb, (b) Tie47Ale3Nbe0.3Sie0.2C and (c) Tie45Ale3Nbe0.3Sie0.2C showing the distribution of a2 lamellae.
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Fig. 5. Relation between fracture toughness and a2ea2 spacing.
As we can see in Fig. 6, a2ea2 spacing is distributed for a long range of order. Therefore, the crack propagating resistance is not uniform during bend test and the maximum peak considered indicates the total resistance. 3.3. Modeling of crack propagating mode in type I specimens To understand the fracture toughness change with a2ea2 spacing, crack propagation sequence and paths should be examined. Hwu et al. examined the crack propagation behavior using metal/ceramic laminate systems [24,25]. They investigated the transition from single crack extension to multiple cracking near the crack tip over a range of metal/ceramic thickness ratios (tm/tc). The transition was controlled by brittle ceramic thickness; crack reinitiated with multiple cracking when the ceramic layer was thin, however, crack reinitiated with macroscopic crack propagation when the ceramic layer was thick. Macroscopic crack propagation showed higher
toughness because load cannot be concentrated at the crack tip. Considering this report and delamination in a2 phases, crack propagation behavior in type I specimens with small and high a2ea2 spacing is shown in Figs. 7 and 8. In the case of small a2ea2 spacing (Fig. 7), delamination will occur when the sharp crack meets the boundary (mainly a2/g), and high load will be needed for further propagation. Next, delamination in the next a2 phase will occur instead of fracture in the middle g. Therefore, it can be very hard to break middle g phase because load cannot be concentrated at the crack tip. In the case of high a2e a2 spacing (Fig. 8), delamination will occur when the sharp crack meets the boundary (mainly a2/g), and high load will be needed for further propagation (same with Fig. 7). However, middle g phases will break instead of delamination in next a2 phase because the next a2 phase is located out of the stress field. Therefore, low load will be needed to break middle g phases because load can be concentrated at the crack tip. Figs. 7 and 8 can explain the result of Fig. 2 and Table 1.
Fig. 6. The distribution of a2ea2 spacing of Tie47Ale3Nbe0.3Sie0.2C alloy.
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Fig. 7. Schematic illustration of crack propagating process in type I alloy with small a2ea2 spacing.
3.4. Crack initiation in Bi-PST crystals with different twist angle and compositions In this paper, crack propagation behavior in Bi-PST crystals was also investigated. To understand the crack propagation behavior, Bi-PST crystals with different twist angles (Fig. 1b and c) and compositions were prepared by DS process in FZ furnace. Table 2 illustrates the kink (a) and twist (b) angles of specimens in this paper. As shown in Table 2, three specimens of Tie45Al and one Tie45Ale3Nb specimen have similar twist angle differences (Db), however, that of Tie45Ale3Nbe 0.2Sie0.1C alloy was very different. Before the bend test, fatigue precracks were generated at the notch tip by compressionecompression fatigue test. All the precracks were examined by OM (optical microscope) and they were made parallel with lamellar boundaries. Loadedisplacement curves obtained from the bend test are shown in Fig. 9. As shown in Fig. 9, the maximum load of
Tie45Al was very similar (523e576 N). Considering they have similar Db and different Da, the toughness can be controlled by Db rather than Da as already reported by Ping et al. [26]. In case of Tie45Ale3Nb alloy, maximum load was higher than those of Tie45Al alloys in spite of similar Db. This was caused by different a2ea2 spacings in which delamination occur (it will be discussed later). The maximum load of Tie45Ale3Nbe0.2Sie0.1C alloy was very low compared to other compositions because of small Db. In the case of Tie45Ale3Nbe0.2Sie0.1C alloy, the kink angle of colony 1 was very high (45 ), therefore, more than two precracks were generated and propagated (Fig. 10). From Figs. 9 and 10, toughness of Bi-PST crystals was controlled by Db. Interrupted in situ test was performed to understand these results and the crack path was examined. Interrupted in situ test was performed by slightly loading using loading jig and observing using OM and SEM (BS images). Fig. 11 shows the cracks after 1st (Fig. 11a) and 2nd (Fig. 11b) loadings.
Fig. 8. Schematic illustration of crack propagating process in type I alloy with high a2ea2 spacing.
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982 Table 2 Kink and twist angles of Bi-PST crystals Specimen
a ( )
b ( )
Colony 1
Colony 2
Colony 1
Colony 2
45Al #1 45Al #2 45Al #3 45Ale3Nb 45Ale3Nbe0.2Sie0.1C
35 35 18 15 45
5 5 43 36 35
10 10 <5 <10 25
48 48 45 45 45
From Fig. 11, it was inferred that one of the precracks propagated after 1st loading (Fig. 11a) and stopped, then, the other crack propagated after 2nd loading (Fig. 11b). Fig. 12 shows the cracks after 3rd loading by OM (Fig. 12a) and SEM BS images (Fig. 12b). Main crack had already crossed boundary and propagation mode was changed from delamination (colony 1) to translamellar (colony 2). It was very noticeable that so many delaminations occurred in colony 2 (Fig. 12a) and these
Fig. 9. Room temperature loadedisplacement curves from the three-point bend tests of Tie45Al, Tie45Ale3Nb and Tie45Ale3Nbe0.2Sie0.1C alloys.
delaminations occurred mainly in a2/g lamellar boundaries and a2 phases (Fig. 12b). Yokoshima and Yamaguchi already reported that delamination cannot give any resistance for crack propagation and translamellar type propagation can give a lot
Fig. 10. Microstructures after three-point bend test of Tie45Ale3Nbe0.2Sie0.1C alloy.
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Fig. 11. Microstructures of Tie45Al alloy during interrupted in situ test after (a) first and (b) second loadings.
of resistance [20]. Therefore, the change from delamination to translamellar at the boundary can be explained as that twisting needs higher energy than translamellar. So, we can guess that crack will propagate by delamination after boundary when Db is small (Fig. 10). Moreover, if the delamination in colony 2 occurred more frequently (small a2ea2 spacing in colony 2), the maximum load after three-point bend test can be higher as in the case of Tie45Ale3Nb alloy. The average a2ea2 spacings of Tie45Al and Tie45Ale3Nb obtained by SEM BS images were 1.11 and 0.68 mm, respectively. Therefore, more delamination occurred in Tie45Ale3Nb alloy so that load was dispersed and maximum load was higher than Tie45Al alloys (Fig. 9). Therefore, crack resistance at the colony boundaries was controlled mainly by Db and secondly by a2ea2 spacing in colony 2.
4. Conclusions The crack propagation behavior of directional solidified Tie(45,47)Ale3Nb, Tie(45,47)Ale3Nbe0.2Sie0.1C, Tie (45,47)Ale3Nbe0.3Sie0.2C type I specimens and Tie45Al, Tie45Ale3Nb, Tie45Ale3Nbe0.2Sie0.1C Bi-PST crystals was investigated. It was found that fracture toughness of type I specimens increased with decreasing Al contents and decreased with increasing Si and C contents. The toughness of type I specimens was controlled by a2ea2 spacing. Moreover, crack propagation resistance of Bi-PST crystals was controlled by twist angle differences (Db) rather than kink angle differences (Da). The crack paths of Bi-PST crystals changed from delamination (colony 1) to translamellar (colony 2) with high Db, and did not change with low Db.
Fig. 12. Microstructures of Tie45Al alloy during interrupted in situ test after third loading by (a) OM and (b) SEM BS images.
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Acknowledgements This work was supported by Korea Research Foundation Grant (KRF-2004-042-D00111). References [1] Kim YW, Dimiduk DM. J Met 1991;43:40. [2] Huang SC, Shih DS. In: Kim YW, Boyer RR, editors. Microstructuree property correlation in TiAl-base alloys. Warrendale, PA: TMS; 1991. p. 105. [3] Yamaguchi M, Inui H. TiAl compounds for structural applications. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 127. [4] Kim YW. J Met 1994;46:30. [5] Chan KS, Kim YW. Acta Metall Mater 1995;43:439. [6] Chan KS, Kim YW. Metall Trans 1992;23A:1663. [7] London B, Larsen DK, Wheeler DA, Aimone PR. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 151. [8] Takeyama M, Kumagai T, Nakamura M, Kikuchi M. In: Daloria R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 167.
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