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lamellar competition
Intermetallics 15 (2007) 327e332 www.elsevier.com/locate/intermet
On the massive phase transformation regime in TiAl alloys: The alloying effect on massive/lamellar competition D. Hu*, A.J. Huang, X. Wu Interdisciplinary Research Centre in Materials, The University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 11 May 2006; received in revised form 10 July 2006; accepted 13 July 2006 Available online 8 September 2006
Abstract The influence of alloy composition on the solid-state transformations which occur during continuous cooling of TiAl-based alloys has been studied with the aim of defining the alloy compositions which most strongly influence the massive transformation. It has been found that heavy elements such as Nb and Ta decrease the cooling rates required to form massive gamma by suppressing the formation of feathery and lamellar structures to even lower cooling rates. It is proposed that the low diffusivity of such heavy elements retards these diffusion-controlled phase transformations (feathery and lamellar) so that the massive transformation, which does not require diffusion, but which occurs at a lower temperature can take place. The significance of the selection of alloying elements, which allow the formation of massive gamma, is briefly discussed in terms of the refinement of the microstructures of castings by heat treating massively transformed samples. Ó 2006 Elsevier Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; B. Phase transformations; B. Alloy design
1. Introduction The study of continuous cooling phase transformations has been an active field in the long quest into high performance engineering TiAl-based alloys [1e11], because different microstructures in these alloys can be obtained through solid-state phase transformations during continuous cooling. Those microstructures can be the final microstructures in components, such as the lamellar microstructure, or a precursor to the final microstructures, such as the massive gamma microstructure. The massive transformation in TiAl-based alloys was first reported in 1992 and has attracted much attention since then [12e15]. It is a diffusionless transformation in which a highly faulted gamma is generated from the parent high-temperature alpha phase by appropriate rapid cooling. Although massive gamma has not been used as a microstructure in service, it has been used as a precursor for the
* Corresponding author. Tel.: þ44 1214143445. E-mail address: [email protected] (D. Hu). 0966-9795/$ - see front matter Ó 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.07.007
formation of microstructures having shown improved properties. Thus heat treatments of massive gamma have been developed to refine the microstructures in cast TiAl alloys which are normally very difficult to refine without thermo-mechanical processing [16e23]. By heat treating massive gamma at temperatures above/close to the alpha transus for a very short time, fine alpha grains can be obtained. Those small alpha grains will, on subsequent cooling, transform into small lamellar colonies [16,17,20]. If the massive gamma is heat treated in the a þ g two-phase field, alpha plates will precipitate on all {111} planes of the massive gamma to form fine microstructures [18,19,21,22,24]. Microstructural refinement based on massive gamma can lead to improvement in properties, especially room temperature tensile ductility [22] and this has a potential benefit in improving the limited ductility e a characteristic of cast TiAl-based alloys. There are two potential problems in developing the massive transformation as a useful process route for TiAl-based alloys; firstly, the lack of understanding of the factors controlling this transformation and secondly the high cooling rates commonly required for massive transformation. A high cooling rate can lead to
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fracture in brittle alloys and is difficult to achieve in large samples. Thus the sample size is limited. Recent work aimed at reducing quench cracking has focused on quenching samples to below the massive transformation temperature but above the ductile-to-brittle temperature by using a salt bath [22]. It is known [9,10] that alloys with different compositions have different responses to continuous cooling but there is at present no way of predicting the range of cooling rates required to produce massive gamma in the given TiAl-based alloys. Thus, in order to utilize the massive transformation as a method of microstructural refinement in cast TiAl alloys it is necessary to define the massive transformation regime in each alloy. Not all the alloys developed so far are suitable for microstructural refinement through massive transformation since in some alloys the massive transformation regime is either too narrow or is shifted to very high cooling rate side which is not easy to reach in practice. This paper is aimed at understanding the alloying factors which control the massive transformation regime in TiAl alloys in order to develop alloys in which massive transformation can occur on medium or slow cooling. 2. Experimental Various TiAl-based alloys of compositions listed in Table 1 were prepared via cold hearth plasma arc melting and cast into either ingots or buttons. The ingots are 100 mm in diameter and about 25 kg each. The buttons are semi-spherical, weighing about 1 kg each. The oxygen level of the alloys is in given in Table 1 as well. In order to investigate the influence of cooling rate on the transformation behaviour of different alloys several approaches have been adopted. Firstly, standard Jominy end quenching test pieces with a diameter of 25 mm and a net length of 97 mm were machined from ingots. The cooling rate profile along the axis in TiAl Jominy end quenching specimen was established in the previous work, as between 4 and 350 C s1, varying continuously with the distance from the quenched end [25]. Jominy end quenching specimens were sectioned along the axis after quenching. By examining the microstructures along the axis and measuring the distance from the quenching end, the microstructureecooling rate relationship can be established. One example is reported Table 1 TiAl alloys used in this study (all in at%) prepared by cold heath plasma arc melting Alloy
Form
Weight (kg)
Oxygen (wt ppm)
Tie48Ale2Cre2Nb Tie47Ale2Nbe1 Mne1We0.2Si Tie46Ale5Nbe1W Tie46Ale8Nb Tie46Ale10Nb Tie46Ale4Nbe4Ta
Ingot Ingot
25 25
450 610
Ingot Ingot Button Button
25 25 1 1
540 520 560 500
elsewhere [26]. Secondly, since the buttons are not large enough for Jominy end quenching specimens, wedge-shaped samples of 20 mm thick and 45 mm long with the widths at two ends being 5 mm and 20 mm respectively were taken from alloy buttons. When the wedge samples were air cooled, different positions give different slow cooling rates. Cooling rates beyond the range available from Jominy samples and from air-cooled wedge samples were achieved by controlled furnace cooling and water quenching of 10 10 10 mm cubic samples. The cooling rate used in this study is referred to as the ‘nominal cooling rate’ since it was determined from the cooling curves from 1150 C to avoid reaction between samples and the thermocouples and the cooling rate at 900 C is used. The solution treatment temperature was 20e30 C above the alpha transus of the alloys and holding time was 1 h. The microstructures were examined using optical microscopy. Grain sizes and volume fractions of the phases observed were determined using linear intercept method [27]. The alpha grain size in the solution treated samples was in the range of 550e 850 mm. An early study showed that when the alpha grain size was decreased to about 300 mm, its influence on lamellar/massive competition could be observed [26]. Thus, in the range in this study, the influence of alpha grain size is regarded as insignificant. 3. Results and discussion 3.1. Microstructureecooling rate relationship from Jominy samples All the alloys used in this investigation have a single alpha phase field followed by an a þ g two-phase field. Upon continuous cooling from the alpha single phase field the alpha phase may decompose into different microstructures depending on the cooling rate. For the highest cooling rate retained a2 is formed, followed by massively transformed gamma, feathery lamellar microstructures and at the slowest rates a lamellar microstructure is formed, in general. Each type of microstructure is formed over a defined cooling rate range but the cooling rate ranges for different microstructures often overlap with others. Thus in a given alloy a totally massive microstructure will be formed for a specific cooling rate range, but if the cooling rate is increased some regions of the sample will contain retained a2 and if the cooling rate is decreased some regions will form feathery/lamellar microstructures. Not all microstructure regimes exist in all the TiAl-based alloys for the range of cooling rates investigated here. Fig. 1, for example, shows the typical microstructures in Tie46Ale8Nb after Jominy end quenching. Close to the quenched end of the sample, which corresponds to a nominal cooling rate of 180 C s1, the microstructure is fully massive g as shown in Fig. 1a and for this alloy higher cooling rates are required to give rise to retained a2. Away from the quenched end a mixture of massive g, feathery lamellar and lamellar microstructures is formed, as shown in Fig. 1b, corresponding to a nominal cooling rate of 25 C s1 at which three regimes
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overlap. Further away from the quenched end (at a nominal cooling rate of 10 C s1) a mixture of feathery, Widmansta¨tten lamellar and lamellar microstructures is formed, as shown in Fig. 1c. The variation of microstructure with cooling rate obtained for four of the alloys is plotted in Fig. 2. The two vertical dotted lines represent the cooling rate range from Jominy end quenching. The four alloys cover both low-alloyed TiAl
Fig. 2. The microstructureecooling rate relationships of Tie46Ale8Nb (468), Tie46Ale5Nbe1W (4651), Tie47Ale2Nbe1Mne1We0.2Si (47211) and Tie48Ale2Cre2Nb (4822). The alpha grain size of the alloys prior to quenching is, as indicated, in the range of 550e850 mm. The two vertical dashed lines are the cooling rate limits of Jominy end quenching.
Fig. 1. Typical optical microstructures in continuously cooled Tie46Ale8Nb from 1360 C (alpha phase field). (a) Fully massive g obtained at a cooling rate of 180 C s1, (b) a mixture of massive g þ feathery þ lamellar microstructures obtained at a cooling rate of 25 C s1 and (c) a mixture of feathery þ Widmansta¨tten þ lamellar microstructures obtained at a cooling rate of 10 C s1 (there is still a tiny fraction of massive g at this cooling rate). Letters M, F, W and L stand for massive, feathery, Widmansta¨tten and lamellar microstructures, respectively.
(typified by Tie48Ale2Cre2Nb) and high-alloyed alloys (Tie46Ale8Nb). It can be seen from this figure that three out of four alloys show all the possible microstructure regimes for different cooling rates but that Tie46Ale5Nbe1W does not appear to have a feathery microstructure regime. It is not clear at this stage why the feathery regime is suppressed in this alloy. The most heavily alloyed in Fig. 2 is Tie46Ale8Nb and is shown on the top of this chart and the extent of alloying decreases towards the bottom of the chart. It is obvious from this chart that the range of cooling rates over which the massive microstructure is obtained has a strong dependency on the alloy composition but that the cooling rate required to form massive is not very sensitive to alloy composition at high cooling rates in the range of this study (or this technique is not sensitive enough to reveal the difference clearly). Tie46Ale8Nb has the widest fully massive g regime whilst Tie48Ale2Cre 2Nb has the narrowest. Comparison of the extent of the massive regimes in Tie46Ale8Nb and Tie46Ale5Nbe1W illustrates the effect of small changes in alloy content, since the Al concentration is the same and the grain size is similar. The results obtained from the Jominy samples show that although the maximum cooling rate at which massive gamma is obtained is similar in the four alloys, the range of cooling rates over which a fully massive g microstructure is obtained varies significantly in the four alloys. The reduction in the massive regime is mainly caused by upward moving of the lamellar/feathery regimes’ upper boundary, i.e. the highest cooling rate at which the formation of the feathery or lamellar microstructure occurs. When the upper boundary of the feathery/lamellar regimes is suppressed towards low cooling rates, the fully massive g regime is extended. In view of the fact that
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the massive transformation does not involve diffusion and that the generation of feathery and lamellar microstructures requires diffusion it is likely that alloying additions which extend the massive regime towards lower cooling rates are those that diffuse slowly. 3.2. Microstructureecooling rate relationship from alloys with heavy elements In order to assess the role of some heavy alloying additions on the behaviour of the transformations during cooling wedge-shaped samples have been used since the buttons were not large enough for Jominy samples. The two final alloys listed in Table 1 were used for this investigationd Tie46Ale10Nb and Tie46Ale4Nbe4Ta. After solution treatment at 1360 C, the samples were air cooled. The microstructures at the position where the nominal cooling rate was 9 C s1 are shown in Fig. 3 together with that from Tie46Ale8Nb cooled in the same way for comparison. At that cooling rate, there is only a little massive g (arrowed in Fig. 3a) in Tie46Ale8Nb, as expected from the Jominy specimens. The amount of massive gamma in Tie46Ale 10Nb was significantly increased, and was about 40 vol% as can be seen in Fig. 3b. This effect of Nb on formation of massive g was observed earlier in Tie48Al based alloys by Takeyama et al. [28]. By substituting 4%Nb with 4%Ta, the amount of massive g was increased to about 70 vol% (Fig. 3c). These observations show that either increasing the concentration of Nb from 8 to 10 or substituting Nb with Ta will extend the range of cooling rates over which massive gamma is formed towards the low cooling rate side effectively. 4. General discussion Continuous cooling transformation (CCT) curves in TiAl alloys can be divided into the three regions shown in Fig. 4, according to the relative importance of diffusion during the various transformations. On the left hand side (i.e. at short times), in region (I) the diffusionless massive transformation occurs at high cooling rate. The massive transformation involves atoms transferring across the interface between the parent a-phase and the massive g-phase. On the right hand side (i.e. at long times) is the diffusion-related region (III) in which all three phase transformations, the formation of the feathery structure, the Widmansta¨tten lamellar structure and the lamellar structure, are diffusion related since they are all a þ g two-phase microstructures and the two phases have different compositions. Region (II) at intermediate times corresponds to cooling rates where diffusionless and diffusion-controlled transformations both occur. The optical microstructures of Tie46Ale8Nb after Jominy end quenching shown in Fig. 1 correspond to the three regions in Fig. 4. The lamellar transformation in TiAl-based alloys has long been considered to occur via the Blackburn mechanism in which lamellae form by gliding of Shockley partials in the parent a-phase [29]. However, the formation of the g lamellae
Fig. 3. Optical microstructures at the same position of wedge samples of (a) Tie46Ale8Nb, (b) Tie46Ale10Nb and (c) Tie46Ale4Nbe4Ta after air cooling from 1360 C. The nominal cooling rate was 9 C s1.
also involves diffusion as shown by several studies. Thus atom probe analysis has been used to show that the extremely thin a and g lamellae, which are formed when cooling Tie 48Al at 103e104 C s1, have significant compositional differences even when the lamellae were only a few nanometres thick [30e32]. This indicates that diffusion is involved as soon as the g-nuclei are formed. Furthermore, analysis of the interfacial structures of the a2/g lamellae showed that glide and climb of dislocations must be involved and that glide of Shockley partials alone cannot generate these interfaces; diffusion is required as an additional process [33]. Although there are no such detailed studies on the feathery microstructure, it has been shown by Dimiduk and Vasudevan that the feathery
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5. Summary
Fig. 4. Schematic illustration of CCT curves with different regions. M stands for massive, F for feathery, L for lamellar and W for Widmansta¨tten. This diagram clearly indicates that cooling at intermediate rates will generate mixtures of lamellar-type microstructures and massively transformed regions.
microstructure is an a2 þ g two-phase microstructure, with significant compositional differences between the phases as is the lamellar microstructure [9]. The formation of these diffusion-related phases will be affected by the diffusivity of the solutes. The fact that Nb and Ta both slow down the formation of feathery and lamellar microstructures implies that it may be related to their low diffusivity. There are not many published studies on diffusion in TiAl alloys, but it has been shown that Nb is a slow diffuser in both TiAl and Ti3Al with a diffusion coefficient about an order lower than that of Ti [34,35]. Unfortunately, no data is available for diffusion in high temperature a, but judging from their diffusion behaviours in a2 and g phases, it is reasonable to believe Nb diffuses more slowly than Ti in the a-phase too. It could be expected that Ta is an even slower diffuser than Nb in these phases from its relative position in the Periodic Table of Elements to Nb’s. The driving force for diffusion during these transformations is the partitioning of Nb and Ta between the g and a2. The partition coefficient for Nb is about 1.2, with the Nb content being higher in the g-phase [36] so that a significant amount of diffusion of Nb into the gamma phase is required during the lamellar/feathery/Widmansta¨tten transformations. Ta has a partition coefficient of 0.84 [36] and Ta is rich in the alpha phase. Since the mass transportation of alloying elements is delayed due to the low diffusivity, the formation of feathery/lamellar microstructures is suppressed. In Fig. 4 it can be seen that the feathery/lamellar transformation starts at a higher temperature than massive transformation. The suppression of feathery/lamellar transformation gives way to the formation of massive transformation. This should be the way how the massive regime is extended by elements like Nb and Ta. Alloying elements Nb and Ta have another advantage over the others. They are weak beta stabilising elements and will not change the phase constituents significantly at the concentrations considered here [37]. Thus, it is easy to maintain a considerably wide a-phase field which is essential for engineering alloys designed to be grain refined through massive transformation.
Continuous cooling phase transformation behaviour of some TiAl alloys has been studied using Jominy end quenching and air cooling of wedge-shaped samples. These experiments provide a direct correlation between the continuously changing microstructures and cooling rate over a wide range of cooling rates. The results show that the massive regime varies significantly with alloy composition and heavy elements like Nb and Ta can effectively extend the massive regime to the low cooling rate side. The mechanism proposed here is that alloying elements like Nb and Ta have low diffusivity in TiAl alloys, which suppresses the formation of diffusionrelated feathery/lamellar transformation, allowing massive transformation to occur. It is also shown that Ta has a greater effect than Nb in extending the massive regime to low cooling rate side in TiAl alloys, which leads to the possibility to develop alloys able to undergo massive transformation during air cooling in large-sized samples. Acknowledgement The authors would like to thank Prof MH Loretto for his stimulating discussions throughout this work and for comments on the manuscript. DH would like to thank the EPSRC for financial support through various contracts for many years. Support from Nb Products Company is also gratefully acknowledged. References [1] McQuay PA, Dimiduk DM, Semiatin SL. Scripta Metall Mater 1991;25:1689. [2] McQuay PA, Dimiduk DM, Lipsitt HA, Semiatin SL. In: Froes FH, Caplan I, editors. Titanium’92 science and technology. Warrendale, PA: TMS; 1993. p. 1041. [3] Wang P, Vasudevan VK. In: Baker I, Darolia R, Whittenberger JD, Yoo M, editors. High-temperature ordered intermetallic alloys V. Pittsburgh, PA: MRS; 1993. p. 229. [4] Ramanath G, Vasudevan VK. In: Baker I, Darolia R, Whittenberger JD, Yoo M, editors. High-temperature ordered intermetallic alloys V. Pittsburgh, PA: MRS; 1993. p. 223. [5] Jones SA, Kaufman MJ. Acta Metall Mater 1993;41:387. [6] Takeyama M, Kumagai T, Nakamura M, Kikuchi M. In: Darolia R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 167. [7] Denquin A, Naka S. In: Kim Y-W, Wagner R, Yamaguchi M, editors. Gamma titanium aluminides. Warrendale, PA: TMS; 1995. p. 141. [8] Kumagai T, Abe E, Nakamura M. Metall Trans 1998;29A:19. [9] Dimiduk DM, Vasudevan VK. In: Kim Y-W, Dumiduk DM, Loretto MH, editors. Gamma titanium aluminides 1999. Warrendale, PA: TMS; 1999. p. 239. [10] Prasad U, Chaturvedi MC. Metall Trans 2003;34A:2053. [11] Hu D, Blenkinsop PA, Loretto MH. In: Gornin IV, Ushkov SS, editors. Titanium’99. St Petersburg: CRISM; 2000. p. 290. [12] Wang P, Viswanathan GB, Vasudevan VK. Metall Mater Trans 1992;23A:691. [13] Wang P, Vasudevan VK. Scripta Metall Mater 1992;27:89. [14] Denquin A, Naka S. Acta Metall Mater 1996;44:353. [15] Dey SR, Bouzy E, Hazotte A. Intermetallics 2006;14:444. [16] Zhang WJ, Francesconi L, Evangelista E. Mater Lett 1996;27:135. [17] Zhao L, Au P, Beddoes JC, Wallace W. US patent US5653828, 1995.
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[18] Kikuchi M, Nakamura H, Yamabe Y. Japanese patent JP6116691, 1994. [19] Takeyama M, Kumagai T, Aritomi T, Nakamura M. Japanese patent JP6279964, 1994. [20] Wang JN, Xie K. Intermetallics 2000;8:545. [21] Hu D. In: Lutjering G, Albrecht J, editors. Ti-2003 Science and technology. Weinheim: Wiley-VCH; 2004. p. 2369. [22] Wu X, Hu D. Scripta Mater 2005;52:731. [23] Clemens H, Bartels A, Bystrzanowski S, Chladil H, Leitner H, Dehm G, et al. Intermetallics 2006;14:1380. [24] Kumagai T, Abe E, Takeyama M, Nakamura M. Scripta Mater 1997;36:523. [25] Hu D, Huang A, Gregoire A, Li XY, Wu X, Loretto MH. In: Nie JF, Barnett M, editors. Proceedings of the third international conference on advanced materials processing. Melbourne: Institute of Materials Engineering Australia Ltd; 2004. p. 172. [26] Hu D, Huang A, Novovic D, Wu X. Intermetallics 2006;14:818. [27] DeHoff RT, Rhines FN. Quantitative Microscopy. New York: McGrawHill Inc; 1968. p. 45.
[28] Takeyama M, Ohmura Y, Kikuchi M, Matsuo T. Intermetallics 1998;6:643. [29] Blackburn MJ. In: Jaffee RI, Promisel NE, editors. The science, technology, and applications of titanium. Oxford: Pergamon Press; 1970. p. 633. [30] Hono K, Abe E, Kumagai T, Harada H. Scripta Mater 1996;35:495. [31] Menand A, Zapolsky-Tatarenko H, Ne´rac-Partaix A. Mater Sci Eng 1998;A250:55. [32] Tanimura M, Inoue Y. Scripta Mater 1998;39:907. [33] Pond RC, Shang P, Cheng TT, Aindow M. Acta Mater 2000;48:1047. [34] Herzig Chr, Przeorski T, Friesel M, Hisker F, Divinski S. Intermetallics 2001;9:461. [35] Mishin Y, Herzig Chr. Acta Mater 2000;48:589. [36] Forwood CT, Gibson MA, Miller PR, Rossouw CJ, Morton AJ. In: Nathal MV, Darolia R, Liu CT, Martin PL, Miracle DB, Wagner R, et al., editors. Structural intermetallics 1997. Warrendale, PA: TMS; 1997. p. 545. [37] Kainuma R, Fijita Y, Mitsui H, Ohnuma I, Ishida K. Intermetallics 2000;8:855.
On the massive phase transformation regime in TiAl alloys: The alloying effect on massive/lamellar competition D. Hu*, A.J. Huang, X. Wu Interdisciplinary Research Centre in Materials, The University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 11 May 2006; received in revised form 10 July 2006; accepted 13 July 2006 Available online 8 September 2006
Abstract The influence of alloy composition on the solid-state transformations which occur during continuous cooling of TiAl-based alloys has been studied with the aim of defining the alloy compositions which most strongly influence the massive transformation. It has been found that heavy elements such as Nb and Ta decrease the cooling rates required to form massive gamma by suppressing the formation of feathery and lamellar structures to even lower cooling rates. It is proposed that the low diffusivity of such heavy elements retards these diffusion-controlled phase transformations (feathery and lamellar) so that the massive transformation, which does not require diffusion, but which occurs at a lower temperature can take place. The significance of the selection of alloying elements, which allow the formation of massive gamma, is briefly discussed in terms of the refinement of the microstructures of castings by heat treating massively transformed samples. Ó 2006 Elsevier Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; B. Phase transformations; B. Alloy design
1. Introduction The study of continuous cooling phase transformations has been an active field in the long quest into high performance engineering TiAl-based alloys [1e11], because different microstructures in these alloys can be obtained through solid-state phase transformations during continuous cooling. Those microstructures can be the final microstructures in components, such as the lamellar microstructure, or a precursor to the final microstructures, such as the massive gamma microstructure. The massive transformation in TiAl-based alloys was first reported in 1992 and has attracted much attention since then [12e15]. It is a diffusionless transformation in which a highly faulted gamma is generated from the parent high-temperature alpha phase by appropriate rapid cooling. Although massive gamma has not been used as a microstructure in service, it has been used as a precursor for the
* Corresponding author. Tel.: þ44 1214143445. E-mail address: [email protected] (D. Hu). 0966-9795/$ - see front matter Ó 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.07.007
formation of microstructures having shown improved properties. Thus heat treatments of massive gamma have been developed to refine the microstructures in cast TiAl alloys which are normally very difficult to refine without thermo-mechanical processing [16e23]. By heat treating massive gamma at temperatures above/close to the alpha transus for a very short time, fine alpha grains can be obtained. Those small alpha grains will, on subsequent cooling, transform into small lamellar colonies [16,17,20]. If the massive gamma is heat treated in the a þ g two-phase field, alpha plates will precipitate on all {111} planes of the massive gamma to form fine microstructures [18,19,21,22,24]. Microstructural refinement based on massive gamma can lead to improvement in properties, especially room temperature tensile ductility [22] and this has a potential benefit in improving the limited ductility e a characteristic of cast TiAl-based alloys. There are two potential problems in developing the massive transformation as a useful process route for TiAl-based alloys; firstly, the lack of understanding of the factors controlling this transformation and secondly the high cooling rates commonly required for massive transformation. A high cooling rate can lead to
D. Hu et al. / Intermetallics 15 (2007) 327e332
328
fracture in brittle alloys and is difficult to achieve in large samples. Thus the sample size is limited. Recent work aimed at reducing quench cracking has focused on quenching samples to below the massive transformation temperature but above the ductile-to-brittle temperature by using a salt bath [22]. It is known [9,10] that alloys with different compositions have different responses to continuous cooling but there is at present no way of predicting the range of cooling rates required to produce massive gamma in the given TiAl-based alloys. Thus, in order to utilize the massive transformation as a method of microstructural refinement in cast TiAl alloys it is necessary to define the massive transformation regime in each alloy. Not all the alloys developed so far are suitable for microstructural refinement through massive transformation since in some alloys the massive transformation regime is either too narrow or is shifted to very high cooling rate side which is not easy to reach in practice. This paper is aimed at understanding the alloying factors which control the massive transformation regime in TiAl alloys in order to develop alloys in which massive transformation can occur on medium or slow cooling. 2. Experimental Various TiAl-based alloys of compositions listed in Table 1 were prepared via cold hearth plasma arc melting and cast into either ingots or buttons. The ingots are 100 mm in diameter and about 25 kg each. The buttons are semi-spherical, weighing about 1 kg each. The oxygen level of the alloys is in given in Table 1 as well. In order to investigate the influence of cooling rate on the transformation behaviour of different alloys several approaches have been adopted. Firstly, standard Jominy end quenching test pieces with a diameter of 25 mm and a net length of 97 mm were machined from ingots. The cooling rate profile along the axis in TiAl Jominy end quenching specimen was established in the previous work, as between 4 and 350 C s1, varying continuously with the distance from the quenched end [25]. Jominy end quenching specimens were sectioned along the axis after quenching. By examining the microstructures along the axis and measuring the distance from the quenching end, the microstructureecooling rate relationship can be established. One example is reported Table 1 TiAl alloys used in this study (all in at%) prepared by cold heath plasma arc melting Alloy
Form
Weight (kg)
Oxygen (wt ppm)
Tie48Ale2Cre2Nb Tie47Ale2Nbe1 Mne1We0.2Si Tie46Ale5Nbe1W Tie46Ale8Nb Tie46Ale10Nb Tie46Ale4Nbe4Ta
Ingot Ingot
25 25
450 610
Ingot Ingot Button Button
25 25 1 1
540 520 560 500
elsewhere [26]. Secondly, since the buttons are not large enough for Jominy end quenching specimens, wedge-shaped samples of 20 mm thick and 45 mm long with the widths at two ends being 5 mm and 20 mm respectively were taken from alloy buttons. When the wedge samples were air cooled, different positions give different slow cooling rates. Cooling rates beyond the range available from Jominy samples and from air-cooled wedge samples were achieved by controlled furnace cooling and water quenching of 10 10 10 mm cubic samples. The cooling rate used in this study is referred to as the ‘nominal cooling rate’ since it was determined from the cooling curves from 1150 C to avoid reaction between samples and the thermocouples and the cooling rate at 900 C is used. The solution treatment temperature was 20e30 C above the alpha transus of the alloys and holding time was 1 h. The microstructures were examined using optical microscopy. Grain sizes and volume fractions of the phases observed were determined using linear intercept method [27]. The alpha grain size in the solution treated samples was in the range of 550e 850 mm. An early study showed that when the alpha grain size was decreased to about 300 mm, its influence on lamellar/massive competition could be observed [26]. Thus, in the range in this study, the influence of alpha grain size is regarded as insignificant. 3. Results and discussion 3.1. Microstructureecooling rate relationship from Jominy samples All the alloys used in this investigation have a single alpha phase field followed by an a þ g two-phase field. Upon continuous cooling from the alpha single phase field the alpha phase may decompose into different microstructures depending on the cooling rate. For the highest cooling rate retained a2 is formed, followed by massively transformed gamma, feathery lamellar microstructures and at the slowest rates a lamellar microstructure is formed, in general. Each type of microstructure is formed over a defined cooling rate range but the cooling rate ranges for different microstructures often overlap with others. Thus in a given alloy a totally massive microstructure will be formed for a specific cooling rate range, but if the cooling rate is increased some regions of the sample will contain retained a2 and if the cooling rate is decreased some regions will form feathery/lamellar microstructures. Not all microstructure regimes exist in all the TiAl-based alloys for the range of cooling rates investigated here. Fig. 1, for example, shows the typical microstructures in Tie46Ale8Nb after Jominy end quenching. Close to the quenched end of the sample, which corresponds to a nominal cooling rate of 180 C s1, the microstructure is fully massive g as shown in Fig. 1a and for this alloy higher cooling rates are required to give rise to retained a2. Away from the quenched end a mixture of massive g, feathery lamellar and lamellar microstructures is formed, as shown in Fig. 1b, corresponding to a nominal cooling rate of 25 C s1 at which three regimes
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overlap. Further away from the quenched end (at a nominal cooling rate of 10 C s1) a mixture of feathery, Widmansta¨tten lamellar and lamellar microstructures is formed, as shown in Fig. 1c. The variation of microstructure with cooling rate obtained for four of the alloys is plotted in Fig. 2. The two vertical dotted lines represent the cooling rate range from Jominy end quenching. The four alloys cover both low-alloyed TiAl
Fig. 2. The microstructureecooling rate relationships of Tie46Ale8Nb (468), Tie46Ale5Nbe1W (4651), Tie47Ale2Nbe1Mne1We0.2Si (47211) and Tie48Ale2Cre2Nb (4822). The alpha grain size of the alloys prior to quenching is, as indicated, in the range of 550e850 mm. The two vertical dashed lines are the cooling rate limits of Jominy end quenching.
Fig. 1. Typical optical microstructures in continuously cooled Tie46Ale8Nb from 1360 C (alpha phase field). (a) Fully massive g obtained at a cooling rate of 180 C s1, (b) a mixture of massive g þ feathery þ lamellar microstructures obtained at a cooling rate of 25 C s1 and (c) a mixture of feathery þ Widmansta¨tten þ lamellar microstructures obtained at a cooling rate of 10 C s1 (there is still a tiny fraction of massive g at this cooling rate). Letters M, F, W and L stand for massive, feathery, Widmansta¨tten and lamellar microstructures, respectively.
(typified by Tie48Ale2Cre2Nb) and high-alloyed alloys (Tie46Ale8Nb). It can be seen from this figure that three out of four alloys show all the possible microstructure regimes for different cooling rates but that Tie46Ale5Nbe1W does not appear to have a feathery microstructure regime. It is not clear at this stage why the feathery regime is suppressed in this alloy. The most heavily alloyed in Fig. 2 is Tie46Ale8Nb and is shown on the top of this chart and the extent of alloying decreases towards the bottom of the chart. It is obvious from this chart that the range of cooling rates over which the massive microstructure is obtained has a strong dependency on the alloy composition but that the cooling rate required to form massive is not very sensitive to alloy composition at high cooling rates in the range of this study (or this technique is not sensitive enough to reveal the difference clearly). Tie46Ale8Nb has the widest fully massive g regime whilst Tie48Ale2Cre 2Nb has the narrowest. Comparison of the extent of the massive regimes in Tie46Ale8Nb and Tie46Ale5Nbe1W illustrates the effect of small changes in alloy content, since the Al concentration is the same and the grain size is similar. The results obtained from the Jominy samples show that although the maximum cooling rate at which massive gamma is obtained is similar in the four alloys, the range of cooling rates over which a fully massive g microstructure is obtained varies significantly in the four alloys. The reduction in the massive regime is mainly caused by upward moving of the lamellar/feathery regimes’ upper boundary, i.e. the highest cooling rate at which the formation of the feathery or lamellar microstructure occurs. When the upper boundary of the feathery/lamellar regimes is suppressed towards low cooling rates, the fully massive g regime is extended. In view of the fact that
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the massive transformation does not involve diffusion and that the generation of feathery and lamellar microstructures requires diffusion it is likely that alloying additions which extend the massive regime towards lower cooling rates are those that diffuse slowly. 3.2. Microstructureecooling rate relationship from alloys with heavy elements In order to assess the role of some heavy alloying additions on the behaviour of the transformations during cooling wedge-shaped samples have been used since the buttons were not large enough for Jominy samples. The two final alloys listed in Table 1 were used for this investigationd Tie46Ale10Nb and Tie46Ale4Nbe4Ta. After solution treatment at 1360 C, the samples were air cooled. The microstructures at the position where the nominal cooling rate was 9 C s1 are shown in Fig. 3 together with that from Tie46Ale8Nb cooled in the same way for comparison. At that cooling rate, there is only a little massive g (arrowed in Fig. 3a) in Tie46Ale8Nb, as expected from the Jominy specimens. The amount of massive gamma in Tie46Ale 10Nb was significantly increased, and was about 40 vol% as can be seen in Fig. 3b. This effect of Nb on formation of massive g was observed earlier in Tie48Al based alloys by Takeyama et al. [28]. By substituting 4%Nb with 4%Ta, the amount of massive g was increased to about 70 vol% (Fig. 3c). These observations show that either increasing the concentration of Nb from 8 to 10 or substituting Nb with Ta will extend the range of cooling rates over which massive gamma is formed towards the low cooling rate side effectively. 4. General discussion Continuous cooling transformation (CCT) curves in TiAl alloys can be divided into the three regions shown in Fig. 4, according to the relative importance of diffusion during the various transformations. On the left hand side (i.e. at short times), in region (I) the diffusionless massive transformation occurs at high cooling rate. The massive transformation involves atoms transferring across the interface between the parent a-phase and the massive g-phase. On the right hand side (i.e. at long times) is the diffusion-related region (III) in which all three phase transformations, the formation of the feathery structure, the Widmansta¨tten lamellar structure and the lamellar structure, are diffusion related since they are all a þ g two-phase microstructures and the two phases have different compositions. Region (II) at intermediate times corresponds to cooling rates where diffusionless and diffusion-controlled transformations both occur. The optical microstructures of Tie46Ale8Nb after Jominy end quenching shown in Fig. 1 correspond to the three regions in Fig. 4. The lamellar transformation in TiAl-based alloys has long been considered to occur via the Blackburn mechanism in which lamellae form by gliding of Shockley partials in the parent a-phase [29]. However, the formation of the g lamellae
Fig. 3. Optical microstructures at the same position of wedge samples of (a) Tie46Ale8Nb, (b) Tie46Ale10Nb and (c) Tie46Ale4Nbe4Ta after air cooling from 1360 C. The nominal cooling rate was 9 C s1.
also involves diffusion as shown by several studies. Thus atom probe analysis has been used to show that the extremely thin a and g lamellae, which are formed when cooling Tie 48Al at 103e104 C s1, have significant compositional differences even when the lamellae were only a few nanometres thick [30e32]. This indicates that diffusion is involved as soon as the g-nuclei are formed. Furthermore, analysis of the interfacial structures of the a2/g lamellae showed that glide and climb of dislocations must be involved and that glide of Shockley partials alone cannot generate these interfaces; diffusion is required as an additional process [33]. Although there are no such detailed studies on the feathery microstructure, it has been shown by Dimiduk and Vasudevan that the feathery
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5. Summary
Fig. 4. Schematic illustration of CCT curves with different regions. M stands for massive, F for feathery, L for lamellar and W for Widmansta¨tten. This diagram clearly indicates that cooling at intermediate rates will generate mixtures of lamellar-type microstructures and massively transformed regions.
microstructure is an a2 þ g two-phase microstructure, with significant compositional differences between the phases as is the lamellar microstructure [9]. The formation of these diffusion-related phases will be affected by the diffusivity of the solutes. The fact that Nb and Ta both slow down the formation of feathery and lamellar microstructures implies that it may be related to their low diffusivity. There are not many published studies on diffusion in TiAl alloys, but it has been shown that Nb is a slow diffuser in both TiAl and Ti3Al with a diffusion coefficient about an order lower than that of Ti [34,35]. Unfortunately, no data is available for diffusion in high temperature a, but judging from their diffusion behaviours in a2 and g phases, it is reasonable to believe Nb diffuses more slowly than Ti in the a-phase too. It could be expected that Ta is an even slower diffuser than Nb in these phases from its relative position in the Periodic Table of Elements to Nb’s. The driving force for diffusion during these transformations is the partitioning of Nb and Ta between the g and a2. The partition coefficient for Nb is about 1.2, with the Nb content being higher in the g-phase [36] so that a significant amount of diffusion of Nb into the gamma phase is required during the lamellar/feathery/Widmansta¨tten transformations. Ta has a partition coefficient of 0.84 [36] and Ta is rich in the alpha phase. Since the mass transportation of alloying elements is delayed due to the low diffusivity, the formation of feathery/lamellar microstructures is suppressed. In Fig. 4 it can be seen that the feathery/lamellar transformation starts at a higher temperature than massive transformation. The suppression of feathery/lamellar transformation gives way to the formation of massive transformation. This should be the way how the massive regime is extended by elements like Nb and Ta. Alloying elements Nb and Ta have another advantage over the others. They are weak beta stabilising elements and will not change the phase constituents significantly at the concentrations considered here [37]. Thus, it is easy to maintain a considerably wide a-phase field which is essential for engineering alloys designed to be grain refined through massive transformation.
Continuous cooling phase transformation behaviour of some TiAl alloys has been studied using Jominy end quenching and air cooling of wedge-shaped samples. These experiments provide a direct correlation between the continuously changing microstructures and cooling rate over a wide range of cooling rates. The results show that the massive regime varies significantly with alloy composition and heavy elements like Nb and Ta can effectively extend the massive regime to the low cooling rate side. The mechanism proposed here is that alloying elements like Nb and Ta have low diffusivity in TiAl alloys, which suppresses the formation of diffusionrelated feathery/lamellar transformation, allowing massive transformation to occur. It is also shown that Ta has a greater effect than Nb in extending the massive regime to low cooling rate side in TiAl alloys, which leads to the possibility to develop alloys able to undergo massive transformation during air cooling in large-sized samples. Acknowledgement The authors would like to thank Prof MH Loretto for his stimulating discussions throughout this work and for comments on the manuscript. DH would like to thank the EPSRC for financial support through various contracts for many years. Support from Nb Products Company is also gratefully acknowledged. References [1] McQuay PA, Dimiduk DM, Semiatin SL. Scripta Metall Mater 1991;25:1689. [2] McQuay PA, Dimiduk DM, Lipsitt HA, Semiatin SL. In: Froes FH, Caplan I, editors. Titanium’92 science and technology. Warrendale, PA: TMS; 1993. p. 1041. [3] Wang P, Vasudevan VK. In: Baker I, Darolia R, Whittenberger JD, Yoo M, editors. High-temperature ordered intermetallic alloys V. Pittsburgh, PA: MRS; 1993. p. 229. [4] Ramanath G, Vasudevan VK. In: Baker I, Darolia R, Whittenberger JD, Yoo M, editors. High-temperature ordered intermetallic alloys V. Pittsburgh, PA: MRS; 1993. p. 223. [5] Jones SA, Kaufman MJ. Acta Metall Mater 1993;41:387. [6] Takeyama M, Kumagai T, Nakamura M, Kikuchi M. In: Darolia R, Lewandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 167. [7] Denquin A, Naka S. In: Kim Y-W, Wagner R, Yamaguchi M, editors. Gamma titanium aluminides. Warrendale, PA: TMS; 1995. p. 141. [8] Kumagai T, Abe E, Nakamura M. Metall Trans 1998;29A:19. [9] Dimiduk DM, Vasudevan VK. In: Kim Y-W, Dumiduk DM, Loretto MH, editors. Gamma titanium aluminides 1999. Warrendale, PA: TMS; 1999. p. 239. [10] Prasad U, Chaturvedi MC. Metall Trans 2003;34A:2053. [11] Hu D, Blenkinsop PA, Loretto MH. In: Gornin IV, Ushkov SS, editors. Titanium’99. St Petersburg: CRISM; 2000. p. 290. [12] Wang P, Viswanathan GB, Vasudevan VK. Metall Mater Trans 1992;23A:691. [13] Wang P, Vasudevan VK. Scripta Metall Mater 1992;27:89. [14] Denquin A, Naka S. Acta Metall Mater 1996;44:353. [15] Dey SR, Bouzy E, Hazotte A. Intermetallics 2006;14:444. [16] Zhang WJ, Francesconi L, Evangelista E. Mater Lett 1996;27:135. [17] Zhao L, Au P, Beddoes JC, Wallace W. US patent US5653828, 1995.
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