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Microstructures and grain refinement performance of rapidly solidified Al–Ti–C master alloys
Journal of Alloys and Compounds 339 (2002) 180–188
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Microstructures and grain refinement performance of rapidly solidified Al–Ti–C master alloys Zhonghua Zhang, Xiufang Bian*, Zhenqing Wang, Xiangfa Liu, Yan Wang The Key Laboratory of Liquid Structure and Heredity of Materials, Ministry of Education, Shandong University, 73 Jingshi Road, Jinan 250061, PR China Received 14 August 2001; accepted 21 November 2001
Abstract In the present work, the microstructures and grain refining performance of melt-spun Al–3.5Ti–0.15C, Al–5Ti–0.3C and Al–10Ti–1C (in wt.%) alloys have been investigated, using X-ray diffraction (XRD), scanning electron microscopy (SEM), and grain refining tests. All three types of ingot-like alloys contain coarse slablike TiAl 3 and fine TiC particles. The microstructures of melt-spun Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys are composed of spherical or near-spherical dispersed TiC particles and a-Al supersaturated solid solution, while melt-spun Al–10Ti–1C alloy consists of small blocky TiAl 3 , spherical or near-spherical TiC particles and a-Al solid solution. Furthermore, most of the TiAl 3 and TiC particles are distributed in the edge zone of the chill surface of the rapidly solidified alloy ribbons. Rapid solidification process (RSP) does not improve the grain refinement performance of Al–Ti–C master alloys. The reasons for the formation of microstructures and the grain refining performance of melt-spun Al–Ti–C master alloys have also been discussed. 2002 Elsevier Science B.V. All rights reserved. Keywords: Al–Ti–C alloys; Rapid solidification; Scanning electron microscopy; X-ray diffraction
1. Introduction Grain refinement has been an important technique for improving the soundness of aluminium products for many years. The addition of grain refiners, usually master alloys containing potent nucleant particles, promotes formation of a fine equiaxed macrostructure by deliberately suppressing the growth of columnar and twin columnar grains. The finer grain size reduces the size of defects such as microporosity and second-phase particles, producing improved mechanical properties [1]. There are three main groups of commercial grain refiners: Al–Ti, Al–Ti–B and Al–Ti–C [2]. The Al–Ti–B master alloys containing soluble TiAl 3 and insoluble TiB 2 particles have been dominant for 20 years and cover a wide range of chemical compositions. The Al–Ti grain refiners containing soluble TiAl 3 particles are usually used in combination with Al–Ti–B alloys. Since one of the earliest systematic studies was reported by Cibula in 1949 [3], much work has been done in an attempt to understand *Corresponding author. E-mail addresses: zh [email protected] ] [email protected] (X. Bian).
(Z.
Zhang),
the fundamental mechanisms through which Al–Ti–B grain refiners are effective [4–7]. Furthermore, commercial producers of master alloys have made considerable progress over the years in improving the consistency and general performance characteristics of their Al–Ti–B products. However, problems with Al–Ti–B alloys can be agglomeration of the borides, blockage of filters, defects during subsequent forming operations and poisoning by certain elements like Zr, V and Cr. Al–Ti–C grain refiners contain soluble TiAl 3 and insoluble TiC particles and are not affected by poisoning by Zr, etc. Therefore, the process of grain refinement by means of Al–Ti–C is of special importance in some Al alloys for plastic working where the application of Al–Ti–B refiners is inconvenient due to various reasons. Therefore, more attention has been paid to the preparation, microstructure and performance of Al–Ti– C master alloys [3,8–11]. Rapid solidification process (RSP) of metals and alloys involves exceptionally high rates of cooling (10 4 –10 8 K s 21 ) during solidification from the molten state [12]. The levels of undercooling achievable at such high rates of cooling lead to significant and often potentially beneficial modifications to rapidly solidified microstructures compared with those produced under conventional conditions.
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 01 )01965-X
Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
181
These include the refinement of the as-solidified microstructure, the refinement of the scale of segregation, increases in the equilibrium solid solubilities of solute elements and the formation of novel metastable crystalline and amorphous phases [13,14]. A number of studies [15– 17] have been undertaken to investigate the microstructure of rapidly solidified Al–Ti and Al–Ti–B alloys, but little information is available on the rapidly solidified Al–Ti–C alloys. The purpose of the present work is to investigate the effects of RSP on the microstructures and grain refining performance of several rapidly solidified Al–Ti–C master alloys and thus gain an improved understanding of the mechanisms of solidification. Fig. 1. XRD patterns of ingot-like Al–Ti–C alloys.
2. Experimental
3. Results
The alloys used in this work, Al–10Ti–1C, Al–5Ti– 0.3C and Al–3.5Ti–0.15C (nominal compositions in wt.%) were prepared using a novel and economic method reported in the literature [18]. The real compositions of the alloys are presented in Table 1, using chemical analysis for Ti and combustion method for C [10]. The master alloys were remelted by high-frequency induction heating in a controlled inert atmosphere in quartz crucible. The melts were first heated to a certain temperature, held at that temperature for 15 min, and then cast into continuous ribbons using the single roller melt spinning technique under a partial argon atmosphere. A wheel speed of 1000 revolutions per minute (rpm), corresponding to a tangential speed of |18 m s 21 , was adopted to make the ribbons 30–50 mm in thickness and 3–5 mm in width. The phase and microstructure identification of the ribbons was carried out using X-ray diffraction (XRD) and scanning electron microscopy (SEM). The chill surface of the ribbons was slightly etched using 0.1% HF for SEM observations. The grain refining performance of both ingot-like and melt-spun Al–Ti–C alloys was tested using additions of 0.005 wt.% Ti and a holding time of 60 s. Dry sand mould was used in testing process. In each set of grain refining tests, the performance of melt-spun alloy was compared directly with that of ingot-like alloy, using a single batch of commercial purity Al as the base alloy to which additions were made. Grain size was obtained from the average of diameters of 20 grains in the test samples which were slightly etched to give grain contrast.
Fig. 1 shows the XRD patterns of ingot-like Al–Ti–C master alloys. All three types of alloys are composed of three phases: a-Al solid solution, TiAl 3 and TiC, regardless of the alloy compositions. The amount of TiAl 3 and TiC in the Al–10Ti–1C is much more than the other two alloys, and the amount of TiC in the Al–3.5Ti–0.15C is minimum, according to the peaks of TiAl 3 and TiC on the XRD patterns. The XRD patterns of melt-spun Al–Ti–C alloys are quite different from those of ingot-like alloys, illustrated in Fig. 2. The rapidly solidified alloys containing lower Ti and C, viz. Al–3.5Ti–0.15C and Al–5Ti– 0.3C, are composed of two phases: a-Al supersaturated solid solution and TiC. No reflections of TiAl 3 are found on the XRD patterns of the two alloys, which means that most of the excess Ti is supersaturated in a-Al solid solution, except the Ti contained in TiC. The rapidly solidified Al–10Ti–1C is still composed of three phases: a-Al solid solution, TiAl 3 and TiC, similar to the ingotlike alloy. On the whole, the peaks of TiAl 3 and TiC on the patterns of rapidly solidified alloys become lower and broader than those of ingot-like alloys.
Table 1 Chemical composition of Al–Ti–C alloys Alloy
Ti (wt.%)
C (wt.%)
Al (wt.%)
Al–10Ti–1C Al–5Ti–0.3C Al–3.5Ti–0.15C
9.64 5.10 3.56
0.75 0.25 0.10
Bal. Bal. Bal.
Fig. 2. XRD patterns of rapidly solidified Al–Ti–C alloys.
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The SEM microstructures of ingot-like Al–Ti–C alloys are shown in Fig. 3. The microstructures of the three alloys consist of coarse, slablike TiAl 3 and fine, particle-like TiC embedded in a-Al solid solution, which is in agreement with the XRD results in Fig. 1. The size of TiAl 3 slabs is 30–90 mm in length and 3–8 mm in width. Furthermore, some of the TiAl 3 slabs are parallel to each other. The microstructures of melt-spun Al–Ti–C alloys are shown in Fig. 4, and are quite different from those of
ingot-like alloys. The chill surface of the ribbons was slightly etched to show clearly the morphologies and distribution of TiC and TiAl 3 particles. Fine particles are uniformly dispersed in the edge zone of the chill surface of Al–3.5Ti–0.15C and Al–5Ti–0.3C alloy ribbons, as seen in Fig. 4(a) and (b), respectively. In combination with XRD patterns shown in Fig. 2, most of the particles are TiC, spherical or near-spherical in morphology and less than 1 mm in size, as seen in Fig. 4(d) and (e). No TiAl 3
Fig. 3. SEM micrographs of ingot-like Al–Ti–C alloys: (a) Al–3.5Ti–0.15C; (b) Al–5Ti–0.3C; (c) Al–10Ti–1C.
Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
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Fig. 4. SEM micrographs of rapidly solidified Al–Ti–C alloys; a–c, g–i represent the edge zone and central zone of chill surface of the ribbons, respectively. (a, g) Al–3Ti–0.15C; (b, h) Al–5Ti–0.3C; (c, i) Al–10Ti–1C; (d, e, f) enlarged graphs corresponding to a, b, c.
particles are found in Al–3.5Ti–0.15C, however, a small amount of blocky TiAl 3 particles are observed in Al–5Ti– 0.3C, as indicated by an arrow in Fig. 4(e). Because of the very small amount of TiAl 3 , they cannot be detected on the XRD pattern of rapidly solidified Al–5Ti–0.3C. Different from Al–3.5Ti–0.15C and Al–5Ti–0.3C, a large number of blocky TiAl 3 particles (indicated by arrows) and spherical or near-spherical TiC particles are observed in Al–10Ti–1C, shown in Fig. 4(c). This is in good agreement with the XRD result of rapidly solidified Al–10Ti– 1C. Furthermore, both TiAl 3 and TiC are zonally distributed in a-Al matrix. It can be seen from Fig. 4(a–c) that the amount of TiC particles dramatically increases with increasing Ti and C concentration. The central zone of chill surface of melt-spun Al–Ti–C alloys is different. No TiC or TiAl 3 particles can be observed in the central zone of chill surface of Al–3.5Ti–0.15C alloy ribbons, and a small
amount of TiC particles appears in that of Al–5Ti–0.3C alloy ribbons, as seen in Fig. 4(g) and (h). Fig. 4(i) shows that a number of TiC and TiAl 3 particles can still be observed in the central zone of chill surface of Al–10Ti– 1C alloy ribbons, although the amount is much less than that in edge zone of the alloy. Fig. 5 shows the macrographs of commercial purity Al unrefined and refined by Al–Ti–C and Al–Ti–B master alloys. Fig. 5(a–c) show the macrographs of commercial purity Al refined by ingot-like Al–3.5Ti–0.15C, Al–5Ti– 0.3C and Al–10Ti–1C alloys with additions of 0.005 wt.% Ti, respectively. In comparison, the macrographs of commercial purity Al refined by rapidly solidified Al–Ti–C master alloys are shown in Fig. 5(d–f). The macrographs of commercial purity Al unrefined and refined by Al–5Ti– 1B master alloy with the same addition are also presented, as shown in Fig. 5(g) and (h), respectively. The commer-
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Fig. 4. (continued)
Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
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Fig. 5. Macrographs of commercial purity Al refined by ingot-like (a–c) and melt-spun (d–f) Al–Ti–C alloys, unrefined (g) and by ingot-like Al–5Ti–1B alloy (h). (a, d) Al–3.5Ti–0.15C; (b, e) Al–5Ti–0.3C; (c, f) Al–10Ti–1C.
cial purity unrefined Al is composed of very coarse equiaxed grains. The addition of Al–Ti–C master alloys can markedly refine the commercial purity Al, as well as the Al–5Ti–1B master alloy. Furthermore, rapid solidifica-
tion can neither improve nor even slightly reduce the refinement performance of Al–Ti–C alloys, as seen in Fig. 5(d–f) compared with Fig. 5(a–c), respectively. Fig. 6 shows the variation of average grain size of commercial
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Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
Fig. 6. Variation of average grain size of commercial purity Al unrefined and refined by various master alloys.
purity Al unrefined and refined by ingot-like Al–Ti–C and Al–Ti–B alloys and melt-spun Al–Ti–C alloys. It can also be seen clearly in Fig. 6 that rapid solidification has no significant effect on the refinement performance of Al–Ti– C master alloys.
4. Discussion All three types of ingot-like Al–Ti–C alloys, viz. Al– 3.5Ti–0.15C, Al–5Ti–0.3C and Al–10Ti–1C, are composed of coarse, slablike TiAl 3 and fine TiC particles, although the amount of the particles increases with increasing Ti and C concentrations in the alloys. The microstructures of the rapidly solidified Al–3.5Ti–0.15C and Al–5Ti–0.3C comprise TiC particles embedded in a-Al solid solution, according to XRD and SEM results. In other words, the formation of TiAl 3 in these two alloys is suppressed. However, the microstructure of rapidly solidified Al–10Ti–1C is still composed of three phases: a-Al solid solution, particle-like TiAl 3 and TiC phases. Moreover, the distribution of TiC and TiAl 3 particles on the chill surface of the ribbons is not uniform. Most of the particles are distributed in the edge zone of the ribbons. The reasons for this are explained below. The stoichiometric ratio (Ti:C, in wt.%) to form TiC phase is estimated to be 4.0. Under equilibrium conditions, one part of Ti combines with C to form TiC, with the remaining Ti existing in the form of TiAl 3 in the rod-like Al–Ti–C alloys. In the present study, RSP has a significant effect on the microstructure of melt-spun Al–Ti–C alloys. TiAl 3 is soluble phase and TiC is insoluble phase [2]. Slablike TiAl 3 dissolved and clusters of TiC dispersed and remained stable at high temperature before casting into ribbons by the melt-spinning technique. In the following rapid solidification process, TiC particles in the melt were retained in the ribbons. Moreover, TiC particles in the ribbons did not agglomerate to form clusters because of the
very high cooling rate. In other words, RSP has no effect on the amount, morphology and size but has on the dispersion of TiC phase. The diameter of the nozzle used in the present work was |0.8 mm, much less than the width of the ribbons obtained (3–5 mm). The droplets sprayed from the nozzle first moved breadthwise on the wheel surface, as indicated in Fig. 7. Therefore, the solidification of the central zone of the ribbons is preliminary to that of the edge zone. Because of the lack of wettability of titanium carbide with molten aluminium [10], TiC particles in the Al–Ti–C melts were excluded to the solidification front during the rapid solidification process. Hence, almost all the TiC particles are distributed in the edge zone of the chill surface of Al–3.5Ti–0.15C and Al–5Ti–0.3C alloy ribbons. For the Al–10Ti–1C alloy ribbons, some particles still remained in the central zone owing to the great amount of TiC and TiAl 3 . Cooling rate plays a very important role in the formation of microstructures of melt-spun Al–Ti–C alloys. The cooling rate R was roughly estimated to be |5310 5 K s 21 from the following equation [19] R 5 h(T 1 2 T 0 ) /CP r t
(1)
where h is the heat transfer coefficient between the ejected molten alloy and the rotating substrate, T 1 and T 0 are the molten and the substrate temperature, respectively, r the density and CP the specific heat of the alloy, and t the thickness of the solidified ribbon specimen. In this work, a value of 1.0 cal?(cm 2 K s)21 was assigned to h for the copper roller [20]. The minimum thermodynamic condition for the suppression of the equilibrium primary intermetallic phase TiAl 3 (b phase) and the formation of supersaturated a solid solution requires that the melt be undercooled to temperatures below the temperature T 0 at which solid and liquid phases of similar compositions have equal free energies [21]. Assuming that the minimum necessary undercooling is achieved, the formation of supersaturated a phase may occur either under conditions in which metastable local equilibrium is preserved at the solid– liquid interface or by non-equilibrium, partitionless solidification from the undercooled melt [14]. Past experimental
Fig. 7. Schematic of movement of droplet.
Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
evidence favors the non-equilibrium case, which involves massive solidification of single-phase solid solution behind a stable planar solid–liquid interface moving at a velocity sufficient to suppress solute partitioning through the process of solute trapping at the interface. In the present study, the very high cooling rate (|5310 5 K s 21 ) ensures that the minimum necessary undercooling is achieved for the complete suppression of the equilibrium primary intermetallic phase TiAl 3 and the formation of supersaturated a-Al solid solution for the Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys. In the case of Al–5Ti–0.3C alloy, except for the Ti to form TiC, the remaining Ti (|4.1 wt.%) is supersaturated in a-Al solid solution and this is in good agreement with previous results [14,22]. With increasing solute content, however, it is expected to be increasingly difficult to achieve the level of undercooling required for partitionless solidification. Attributed to the high Ti content in the Al–10Ti–1C alloy, the level of undercooling is not enough to suppress the formation of TiAl 3 phase. Therefore, both TiC and TiAl 3 particles can be observed in the melt-spun Al–10Ti–1C alloy. Cibula [3] postulated the carbide-boride particle theory which assumed that both TiC and TiB 2 particles are virtually insoluble in aluminium melt and can act as heterogeneous nucleants. In the present work, the addition of Al–Ti–C master alloys can markedly refine commercial purity Al. With the same addition, the grain refinement performance of ingot-like Al–Ti–C alloys is similar to that of commercial Al–5Ti–1B alloy rod. For commercial Al– 5Ti–1B rod alloys, there seems to be sufficient evidence that TiB 2 is the totally dominating phase for the grain refinement [23]. The improved efficiency of Al–Ti–C alloys compared to Al–Ti shows that TiC particles are necessary and play a significant role in the grain refinement. In the former work, it has been reported that both the excess Ti (in the form of TiAl 3 ) and TiC phase play an important role in the grain refinement performance of ingot-like Al–Ti–C master alloys [11]. Two conditions need to be fulfilled in order to obtain efficient grain refinement: (i) a sufficient number of potential nuclei must be present in the melt; and (ii) a large fraction of the potential nuclei must be activated. The grain refinement performance of Al–Ti–C alloys relates not only to the amount of TiC and TiAl 3 in the alloys, but also to the manufacturing process of the alloys. The present results show that the nucleation rate of Al–Ti–C alloys is very low, similar to that of Al–Ti–B alloy. However, the nucleation rate of Al–3.5Ti–0.15C alloy is slightly higher than that of Al–5Ti–0.3C and Al–10Ti–1C alloys. RSP does not improve the grain refinement performance of the Al–Ti–C alloys. Remelting the alloys has a slightly deleterious effect on their performance. The good performance of the melt-spun Al–3.5Ti–0.15C and Al–5Ti– 0.3C alloys containing only TiC particles shows that TiAl 3 particles are not necessary in the rapidly solidified Al–Ti– C alloys, but it does not mean that Ti is not necessary for
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nucleation. This raises a question about the reliability of the peritectic theory and peritectic hulk-theory [24,25]. Furthermore, the ingot-like Al–10Ti–1C master alloy containing slablike TiAl 3 can perform as well as the melt-spun alloy containing blocky TiAl 3 , indicating that the morphology of aluminide is not intrinsically important in influencing the grain refinement performance of the master alloys.
5. Conclusions 1. The microstructures of ingot-like Al–3.5Ti–0.15C, Al– 5Ti–0.3C and Al–10Ti–1C alloys are composed of three phases: a-Al solid solution, coarse slablike TiAl 3 and fine TiC particles. In comparison, the microstructures of melt-spun Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys are composed of two phases: a-Al supersaturated solid solution and spherical or near-spherical dispersed TiC particles, while the microstructure of melt-spun Al–10Ti–1C alloy comprises small blocky TiAl 3 , spherical or near-spherical dispersed TiC particles and a-Al supersaturated solid solution. 2. Most of the TiC and TiAl 3 particles are distributed in the edge zone of the chill surface of melt-spun Al–Ti– C alloys. No particles are observed in the central zone of melt-spun Al–3.5Ti–0.15C alloy and a certain amount of TiC and TiAl 3 particles is distributed in the central zone of melt-spun Al–10Ti–1C alloy. 3. Both ingot-like and melt-spun Al–Ti–C master alloys possess good grain refinement performance. RSP has no obvious effect on the grain refinement performance of Al–Ti–C alloys. For the experimental conditions used in the grain refining tests, the morphologies of TiAl 3 are not of great importance in influencing the performance of grain refiners.
Acknowledgements The authors are grateful for the support of the National Natural Science Foundation and Natural Science Foundation of Shandong Province under grant No. 50171037 and L2000F01.
References [1] M. Easton, D. St John, Metall. Mater. Trans. A 30 (1999) 1613– 1623. [2] H.E. Vatne, Aluminium 75 (1 / 2) (1999) 84–90. [3] A. Cibula, J. Inst. Met. 76 (1949–1950) 321–360. [4] D.G. McCartney, Int. Mater. Rev. 34 (1989) 247–260. [5] M.M. Guzowski, G.K. Sigworth, D.A. Sentner, Metall. Trans. 18A (1987) 603–619. [6] X.F. Liu, X.F. Bian, J.J. Ma, Mater. Sci. Forum 331–337 (2000) 385–389.
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[7] X.F. Liu, Z.G. Zhang, Z. Gao, X.F. Bian, J. Miner., Metals and Mater. Soc. 52 (2000) 47–48. [8] W. Schneider, P. Hoefs, M. Kearns, W. Reif, Aluminium 74 (7–8) (1998) 536–544. [9] A. Cibula, J. Inst. Metals 80 (1951–1952) 1–16. [10] A. Banerji, W. Reif, Metall. Trans. A 16 (1985) 2065–2068. [11] Z.Q. Wang, X.F. Liu, X.F. Bian, Foundry 50 (2001) 316, (in Chinese). [12] H. Jones, J. Mater. Sci. 19 (1984) 1043. [13] R.W.K. Honeycombe, in: B. Cantor (Ed.), Rapidly Quenched Metals III, Vol. 1, The Metals Society, London, 1978, p. 73. [14] A. Majumdar, B.C. Muddle, Mater. Sci. Eng. A 169 (1993) 135– 147. [15] S. Hori, H. Tai, Y. Narita, in: S. Steeb, H. Warlimont (Eds.), Rapidly Quenched Metals 5, Elsevier, Amsterdam, 1985, pp. 911–914. [16] A. Majumdar, R.H. Mair, B.C. Muddle, Mater. Res. Soc. Symp. Proc. 80 (1987) 253–260.
[17] M.G. Chu, Mater. Sci. Eng. A 179 / 180 (1994) 669–675. [18] X.F. Liu, Z.Q. Wang, Z.G. Zhang, X.F. Bian, A method to produce Al–Ti–C master alloy, China patent, Appl. No. 00123953.8, Appl. Date: 2000.10.20. [19] H. Jones, in: H. Herman (Ed.), Ultrarapid Quenching of Liquid Alloys, Treatise on Materials Science and Technology, Vol. 20, Academic Press, New York, 1981, p. 29. [20] M.N. Burden, H. Jones, J. Inst. Met. 98 (1970) 249. [21] J.C. Baker, J.W. Cahn, Acta Metall. 17 (1971) 575. [22] J. Bletry, J. Phys. Chem. Solids 31 (1970) 1263. [23] C.D. Mayes, G.J. Tatlock, D.G. McCartney, in: L.D. Peachley, D.B. Williams (Eds.), Proceedings of the 12th International Congress on Electron Microscopy, San Francisco, 1990, p. 4. [24] F.A. Crossley, L.F. Mondolfo, Trans. AIME (J. Metals) 191 (1951) 1143. ¨ [25] L. Backerud, P. Gustafson, M. Johnsson, Aluminium 67 (1991) 910.
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Microstructures and grain refinement performance of rapidly solidified Al–Ti–C master alloys Zhonghua Zhang, Xiufang Bian*, Zhenqing Wang, Xiangfa Liu, Yan Wang The Key Laboratory of Liquid Structure and Heredity of Materials, Ministry of Education, Shandong University, 73 Jingshi Road, Jinan 250061, PR China Received 14 August 2001; accepted 21 November 2001
Abstract In the present work, the microstructures and grain refining performance of melt-spun Al–3.5Ti–0.15C, Al–5Ti–0.3C and Al–10Ti–1C (in wt.%) alloys have been investigated, using X-ray diffraction (XRD), scanning electron microscopy (SEM), and grain refining tests. All three types of ingot-like alloys contain coarse slablike TiAl 3 and fine TiC particles. The microstructures of melt-spun Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys are composed of spherical or near-spherical dispersed TiC particles and a-Al supersaturated solid solution, while melt-spun Al–10Ti–1C alloy consists of small blocky TiAl 3 , spherical or near-spherical TiC particles and a-Al solid solution. Furthermore, most of the TiAl 3 and TiC particles are distributed in the edge zone of the chill surface of the rapidly solidified alloy ribbons. Rapid solidification process (RSP) does not improve the grain refinement performance of Al–Ti–C master alloys. The reasons for the formation of microstructures and the grain refining performance of melt-spun Al–Ti–C master alloys have also been discussed. 2002 Elsevier Science B.V. All rights reserved. Keywords: Al–Ti–C alloys; Rapid solidification; Scanning electron microscopy; X-ray diffraction
1. Introduction Grain refinement has been an important technique for improving the soundness of aluminium products for many years. The addition of grain refiners, usually master alloys containing potent nucleant particles, promotes formation of a fine equiaxed macrostructure by deliberately suppressing the growth of columnar and twin columnar grains. The finer grain size reduces the size of defects such as microporosity and second-phase particles, producing improved mechanical properties [1]. There are three main groups of commercial grain refiners: Al–Ti, Al–Ti–B and Al–Ti–C [2]. The Al–Ti–B master alloys containing soluble TiAl 3 and insoluble TiB 2 particles have been dominant for 20 years and cover a wide range of chemical compositions. The Al–Ti grain refiners containing soluble TiAl 3 particles are usually used in combination with Al–Ti–B alloys. Since one of the earliest systematic studies was reported by Cibula in 1949 [3], much work has been done in an attempt to understand *Corresponding author. E-mail addresses: zh [email protected] ] [email protected] (X. Bian).
(Z.
Zhang),
the fundamental mechanisms through which Al–Ti–B grain refiners are effective [4–7]. Furthermore, commercial producers of master alloys have made considerable progress over the years in improving the consistency and general performance characteristics of their Al–Ti–B products. However, problems with Al–Ti–B alloys can be agglomeration of the borides, blockage of filters, defects during subsequent forming operations and poisoning by certain elements like Zr, V and Cr. Al–Ti–C grain refiners contain soluble TiAl 3 and insoluble TiC particles and are not affected by poisoning by Zr, etc. Therefore, the process of grain refinement by means of Al–Ti–C is of special importance in some Al alloys for plastic working where the application of Al–Ti–B refiners is inconvenient due to various reasons. Therefore, more attention has been paid to the preparation, microstructure and performance of Al–Ti– C master alloys [3,8–11]. Rapid solidification process (RSP) of metals and alloys involves exceptionally high rates of cooling (10 4 –10 8 K s 21 ) during solidification from the molten state [12]. The levels of undercooling achievable at such high rates of cooling lead to significant and often potentially beneficial modifications to rapidly solidified microstructures compared with those produced under conventional conditions.
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 01 )01965-X
Z. Zhang et al. / Journal of Alloys and Compounds 339 (2002) 180 – 188
181
These include the refinement of the as-solidified microstructure, the refinement of the scale of segregation, increases in the equilibrium solid solubilities of solute elements and the formation of novel metastable crystalline and amorphous phases [13,14]. A number of studies [15– 17] have been undertaken to investigate the microstructure of rapidly solidified Al–Ti and Al–Ti–B alloys, but little information is available on the rapidly solidified Al–Ti–C alloys. The purpose of the present work is to investigate the effects of RSP on the microstructures and grain refining performance of several rapidly solidified Al–Ti–C master alloys and thus gain an improved understanding of the mechanisms of solidification. Fig. 1. XRD patterns of ingot-like Al–Ti–C alloys.
2. Experimental
3. Results
The alloys used in this work, Al–10Ti–1C, Al–5Ti– 0.3C and Al–3.5Ti–0.15C (nominal compositions in wt.%) were prepared using a novel and economic method reported in the literature [18]. The real compositions of the alloys are presented in Table 1, using chemical analysis for Ti and combustion method for C [10]. The master alloys were remelted by high-frequency induction heating in a controlled inert atmosphere in quartz crucible. The melts were first heated to a certain temperature, held at that temperature for 15 min, and then cast into continuous ribbons using the single roller melt spinning technique under a partial argon atmosphere. A wheel speed of 1000 revolutions per minute (rpm), corresponding to a tangential speed of |18 m s 21 , was adopted to make the ribbons 30–50 mm in thickness and 3–5 mm in width. The phase and microstructure identification of the ribbons was carried out using X-ray diffraction (XRD) and scanning electron microscopy (SEM). The chill surface of the ribbons was slightly etched using 0.1% HF for SEM observations. The grain refining performance of both ingot-like and melt-spun Al–Ti–C alloys was tested using additions of 0.005 wt.% Ti and a holding time of 60 s. Dry sand mould was used in testing process. In each set of grain refining tests, the performance of melt-spun alloy was compared directly with that of ingot-like alloy, using a single batch of commercial purity Al as the base alloy to which additions were made. Grain size was obtained from the average of diameters of 20 grains in the test samples which were slightly etched to give grain contrast.
Fig. 1 shows the XRD patterns of ingot-like Al–Ti–C master alloys. All three types of alloys are composed of three phases: a-Al solid solution, TiAl 3 and TiC, regardless of the alloy compositions. The amount of TiAl 3 and TiC in the Al–10Ti–1C is much more than the other two alloys, and the amount of TiC in the Al–3.5Ti–0.15C is minimum, according to the peaks of TiAl 3 and TiC on the XRD patterns. The XRD patterns of melt-spun Al–Ti–C alloys are quite different from those of ingot-like alloys, illustrated in Fig. 2. The rapidly solidified alloys containing lower Ti and C, viz. Al–3.5Ti–0.15C and Al–5Ti– 0.3C, are composed of two phases: a-Al supersaturated solid solution and TiC. No reflections of TiAl 3 are found on the XRD patterns of the two alloys, which means that most of the excess Ti is supersaturated in a-Al solid solution, except the Ti contained in TiC. The rapidly solidified Al–10Ti–1C is still composed of three phases: a-Al solid solution, TiAl 3 and TiC, similar to the ingotlike alloy. On the whole, the peaks of TiAl 3 and TiC on the patterns of rapidly solidified alloys become lower and broader than those of ingot-like alloys.
Table 1 Chemical composition of Al–Ti–C alloys Alloy
Ti (wt.%)
C (wt.%)
Al (wt.%)
Al–10Ti–1C Al–5Ti–0.3C Al–3.5Ti–0.15C
9.64 5.10 3.56
0.75 0.25 0.10
Bal. Bal. Bal.
Fig. 2. XRD patterns of rapidly solidified Al–Ti–C alloys.
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The SEM microstructures of ingot-like Al–Ti–C alloys are shown in Fig. 3. The microstructures of the three alloys consist of coarse, slablike TiAl 3 and fine, particle-like TiC embedded in a-Al solid solution, which is in agreement with the XRD results in Fig. 1. The size of TiAl 3 slabs is 30–90 mm in length and 3–8 mm in width. Furthermore, some of the TiAl 3 slabs are parallel to each other. The microstructures of melt-spun Al–Ti–C alloys are shown in Fig. 4, and are quite different from those of
ingot-like alloys. The chill surface of the ribbons was slightly etched to show clearly the morphologies and distribution of TiC and TiAl 3 particles. Fine particles are uniformly dispersed in the edge zone of the chill surface of Al–3.5Ti–0.15C and Al–5Ti–0.3C alloy ribbons, as seen in Fig. 4(a) and (b), respectively. In combination with XRD patterns shown in Fig. 2, most of the particles are TiC, spherical or near-spherical in morphology and less than 1 mm in size, as seen in Fig. 4(d) and (e). No TiAl 3
Fig. 3. SEM micrographs of ingot-like Al–Ti–C alloys: (a) Al–3.5Ti–0.15C; (b) Al–5Ti–0.3C; (c) Al–10Ti–1C.
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Fig. 4. SEM micrographs of rapidly solidified Al–Ti–C alloys; a–c, g–i represent the edge zone and central zone of chill surface of the ribbons, respectively. (a, g) Al–3Ti–0.15C; (b, h) Al–5Ti–0.3C; (c, i) Al–10Ti–1C; (d, e, f) enlarged graphs corresponding to a, b, c.
particles are found in Al–3.5Ti–0.15C, however, a small amount of blocky TiAl 3 particles are observed in Al–5Ti– 0.3C, as indicated by an arrow in Fig. 4(e). Because of the very small amount of TiAl 3 , they cannot be detected on the XRD pattern of rapidly solidified Al–5Ti–0.3C. Different from Al–3.5Ti–0.15C and Al–5Ti–0.3C, a large number of blocky TiAl 3 particles (indicated by arrows) and spherical or near-spherical TiC particles are observed in Al–10Ti–1C, shown in Fig. 4(c). This is in good agreement with the XRD result of rapidly solidified Al–10Ti– 1C. Furthermore, both TiAl 3 and TiC are zonally distributed in a-Al matrix. It can be seen from Fig. 4(a–c) that the amount of TiC particles dramatically increases with increasing Ti and C concentration. The central zone of chill surface of melt-spun Al–Ti–C alloys is different. No TiC or TiAl 3 particles can be observed in the central zone of chill surface of Al–3.5Ti–0.15C alloy ribbons, and a small
amount of TiC particles appears in that of Al–5Ti–0.3C alloy ribbons, as seen in Fig. 4(g) and (h). Fig. 4(i) shows that a number of TiC and TiAl 3 particles can still be observed in the central zone of chill surface of Al–10Ti– 1C alloy ribbons, although the amount is much less than that in edge zone of the alloy. Fig. 5 shows the macrographs of commercial purity Al unrefined and refined by Al–Ti–C and Al–Ti–B master alloys. Fig. 5(a–c) show the macrographs of commercial purity Al refined by ingot-like Al–3.5Ti–0.15C, Al–5Ti– 0.3C and Al–10Ti–1C alloys with additions of 0.005 wt.% Ti, respectively. In comparison, the macrographs of commercial purity Al refined by rapidly solidified Al–Ti–C master alloys are shown in Fig. 5(d–f). The macrographs of commercial purity Al unrefined and refined by Al–5Ti– 1B master alloy with the same addition are also presented, as shown in Fig. 5(g) and (h), respectively. The commer-
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Fig. 4. (continued)
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Fig. 5. Macrographs of commercial purity Al refined by ingot-like (a–c) and melt-spun (d–f) Al–Ti–C alloys, unrefined (g) and by ingot-like Al–5Ti–1B alloy (h). (a, d) Al–3.5Ti–0.15C; (b, e) Al–5Ti–0.3C; (c, f) Al–10Ti–1C.
cial purity unrefined Al is composed of very coarse equiaxed grains. The addition of Al–Ti–C master alloys can markedly refine the commercial purity Al, as well as the Al–5Ti–1B master alloy. Furthermore, rapid solidifica-
tion can neither improve nor even slightly reduce the refinement performance of Al–Ti–C alloys, as seen in Fig. 5(d–f) compared with Fig. 5(a–c), respectively. Fig. 6 shows the variation of average grain size of commercial
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Fig. 6. Variation of average grain size of commercial purity Al unrefined and refined by various master alloys.
purity Al unrefined and refined by ingot-like Al–Ti–C and Al–Ti–B alloys and melt-spun Al–Ti–C alloys. It can also be seen clearly in Fig. 6 that rapid solidification has no significant effect on the refinement performance of Al–Ti– C master alloys.
4. Discussion All three types of ingot-like Al–Ti–C alloys, viz. Al– 3.5Ti–0.15C, Al–5Ti–0.3C and Al–10Ti–1C, are composed of coarse, slablike TiAl 3 and fine TiC particles, although the amount of the particles increases with increasing Ti and C concentrations in the alloys. The microstructures of the rapidly solidified Al–3.5Ti–0.15C and Al–5Ti–0.3C comprise TiC particles embedded in a-Al solid solution, according to XRD and SEM results. In other words, the formation of TiAl 3 in these two alloys is suppressed. However, the microstructure of rapidly solidified Al–10Ti–1C is still composed of three phases: a-Al solid solution, particle-like TiAl 3 and TiC phases. Moreover, the distribution of TiC and TiAl 3 particles on the chill surface of the ribbons is not uniform. Most of the particles are distributed in the edge zone of the ribbons. The reasons for this are explained below. The stoichiometric ratio (Ti:C, in wt.%) to form TiC phase is estimated to be 4.0. Under equilibrium conditions, one part of Ti combines with C to form TiC, with the remaining Ti existing in the form of TiAl 3 in the rod-like Al–Ti–C alloys. In the present study, RSP has a significant effect on the microstructure of melt-spun Al–Ti–C alloys. TiAl 3 is soluble phase and TiC is insoluble phase [2]. Slablike TiAl 3 dissolved and clusters of TiC dispersed and remained stable at high temperature before casting into ribbons by the melt-spinning technique. In the following rapid solidification process, TiC particles in the melt were retained in the ribbons. Moreover, TiC particles in the ribbons did not agglomerate to form clusters because of the
very high cooling rate. In other words, RSP has no effect on the amount, morphology and size but has on the dispersion of TiC phase. The diameter of the nozzle used in the present work was |0.8 mm, much less than the width of the ribbons obtained (3–5 mm). The droplets sprayed from the nozzle first moved breadthwise on the wheel surface, as indicated in Fig. 7. Therefore, the solidification of the central zone of the ribbons is preliminary to that of the edge zone. Because of the lack of wettability of titanium carbide with molten aluminium [10], TiC particles in the Al–Ti–C melts were excluded to the solidification front during the rapid solidification process. Hence, almost all the TiC particles are distributed in the edge zone of the chill surface of Al–3.5Ti–0.15C and Al–5Ti–0.3C alloy ribbons. For the Al–10Ti–1C alloy ribbons, some particles still remained in the central zone owing to the great amount of TiC and TiAl 3 . Cooling rate plays a very important role in the formation of microstructures of melt-spun Al–Ti–C alloys. The cooling rate R was roughly estimated to be |5310 5 K s 21 from the following equation [19] R 5 h(T 1 2 T 0 ) /CP r t
(1)
where h is the heat transfer coefficient between the ejected molten alloy and the rotating substrate, T 1 and T 0 are the molten and the substrate temperature, respectively, r the density and CP the specific heat of the alloy, and t the thickness of the solidified ribbon specimen. In this work, a value of 1.0 cal?(cm 2 K s)21 was assigned to h for the copper roller [20]. The minimum thermodynamic condition for the suppression of the equilibrium primary intermetallic phase TiAl 3 (b phase) and the formation of supersaturated a solid solution requires that the melt be undercooled to temperatures below the temperature T 0 at which solid and liquid phases of similar compositions have equal free energies [21]. Assuming that the minimum necessary undercooling is achieved, the formation of supersaturated a phase may occur either under conditions in which metastable local equilibrium is preserved at the solid– liquid interface or by non-equilibrium, partitionless solidification from the undercooled melt [14]. Past experimental
Fig. 7. Schematic of movement of droplet.
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evidence favors the non-equilibrium case, which involves massive solidification of single-phase solid solution behind a stable planar solid–liquid interface moving at a velocity sufficient to suppress solute partitioning through the process of solute trapping at the interface. In the present study, the very high cooling rate (|5310 5 K s 21 ) ensures that the minimum necessary undercooling is achieved for the complete suppression of the equilibrium primary intermetallic phase TiAl 3 and the formation of supersaturated a-Al solid solution for the Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys. In the case of Al–5Ti–0.3C alloy, except for the Ti to form TiC, the remaining Ti (|4.1 wt.%) is supersaturated in a-Al solid solution and this is in good agreement with previous results [14,22]. With increasing solute content, however, it is expected to be increasingly difficult to achieve the level of undercooling required for partitionless solidification. Attributed to the high Ti content in the Al–10Ti–1C alloy, the level of undercooling is not enough to suppress the formation of TiAl 3 phase. Therefore, both TiC and TiAl 3 particles can be observed in the melt-spun Al–10Ti–1C alloy. Cibula [3] postulated the carbide-boride particle theory which assumed that both TiC and TiB 2 particles are virtually insoluble in aluminium melt and can act as heterogeneous nucleants. In the present work, the addition of Al–Ti–C master alloys can markedly refine commercial purity Al. With the same addition, the grain refinement performance of ingot-like Al–Ti–C alloys is similar to that of commercial Al–5Ti–1B alloy rod. For commercial Al– 5Ti–1B rod alloys, there seems to be sufficient evidence that TiB 2 is the totally dominating phase for the grain refinement [23]. The improved efficiency of Al–Ti–C alloys compared to Al–Ti shows that TiC particles are necessary and play a significant role in the grain refinement. In the former work, it has been reported that both the excess Ti (in the form of TiAl 3 ) and TiC phase play an important role in the grain refinement performance of ingot-like Al–Ti–C master alloys [11]. Two conditions need to be fulfilled in order to obtain efficient grain refinement: (i) a sufficient number of potential nuclei must be present in the melt; and (ii) a large fraction of the potential nuclei must be activated. The grain refinement performance of Al–Ti–C alloys relates not only to the amount of TiC and TiAl 3 in the alloys, but also to the manufacturing process of the alloys. The present results show that the nucleation rate of Al–Ti–C alloys is very low, similar to that of Al–Ti–B alloy. However, the nucleation rate of Al–3.5Ti–0.15C alloy is slightly higher than that of Al–5Ti–0.3C and Al–10Ti–1C alloys. RSP does not improve the grain refinement performance of the Al–Ti–C alloys. Remelting the alloys has a slightly deleterious effect on their performance. The good performance of the melt-spun Al–3.5Ti–0.15C and Al–5Ti– 0.3C alloys containing only TiC particles shows that TiAl 3 particles are not necessary in the rapidly solidified Al–Ti– C alloys, but it does not mean that Ti is not necessary for
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nucleation. This raises a question about the reliability of the peritectic theory and peritectic hulk-theory [24,25]. Furthermore, the ingot-like Al–10Ti–1C master alloy containing slablike TiAl 3 can perform as well as the melt-spun alloy containing blocky TiAl 3 , indicating that the morphology of aluminide is not intrinsically important in influencing the grain refinement performance of the master alloys.
5. Conclusions 1. The microstructures of ingot-like Al–3.5Ti–0.15C, Al– 5Ti–0.3C and Al–10Ti–1C alloys are composed of three phases: a-Al solid solution, coarse slablike TiAl 3 and fine TiC particles. In comparison, the microstructures of melt-spun Al–3.5Ti–0.15C and Al–5Ti–0.3C alloys are composed of two phases: a-Al supersaturated solid solution and spherical or near-spherical dispersed TiC particles, while the microstructure of melt-spun Al–10Ti–1C alloy comprises small blocky TiAl 3 , spherical or near-spherical dispersed TiC particles and a-Al supersaturated solid solution. 2. Most of the TiC and TiAl 3 particles are distributed in the edge zone of the chill surface of melt-spun Al–Ti– C alloys. No particles are observed in the central zone of melt-spun Al–3.5Ti–0.15C alloy and a certain amount of TiC and TiAl 3 particles is distributed in the central zone of melt-spun Al–10Ti–1C alloy. 3. Both ingot-like and melt-spun Al–Ti–C master alloys possess good grain refinement performance. RSP has no obvious effect on the grain refinement performance of Al–Ti–C alloys. For the experimental conditions used in the grain refining tests, the morphologies of TiAl 3 are not of great importance in influencing the performance of grain refiners.
Acknowledgements The authors are grateful for the support of the National Natural Science Foundation and Natural Science Foundation of Shandong Province under grant No. 50171037 and L2000F01.
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