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Processing Ti3Al–SiC nanocomposites using high energy mechanical milling
Materials Science and Engineering A 375–377 (2004) 911–916
Processing Ti3Al–SiC nanocomposites using high energy mechanical milling D.L. Zhang∗ , J. Liang, J. Wu Department of Materials and Process Engineering, University of Waikato, Private Bag 3105, Hamilton, New Zealand
Abstract Ti3 Al–10 vol.% SiC and Ti3 Al–20 vol.% SiC composite powders were produced by high energy mechanical milling of Ti, Al and SiC powders, using both coarse and fine SiC powders. For a coarse SiC powder, the milling energy has to be high enough to fracture the SiC particles in-situ. While for a fine SiC powder, a lower energy mill is sufficient to produce a nanocomposite powder. Additionally, nanocrystalline Ti(Al) solution was formed during milling. A simple theoretical analysis shows that there exists a lower limit of the particle size below which further reduction of the particle size cannot be achieved by milling since the stresses required for further particle refinement cannot be reached during milling. During consolidation of the Ti3 Al–SiC nanocomposite powders, the reaction between the Ti3 Al matrix and SiC is difficult to be avoided. © 2003 Elsevier B.V. All rights reserved. Keywords: Titanium aluminides; Metal–ceramic nanocomposites; Mechanical milling; Powder processing; Powder consolidation
1. Introduction Metal–ceramic nanocomposites are a type of materials which combine metal phases and ceramic phases with at least one dimension of the ceramic or the metallic phase domains being in the nanometer scale. One popular example of metal–ceramic nanocomposites is metal matrix nanocomposites consisting of discrete ceramic nanoparticles embedded in a continuous metal matrix [1]. Fig. 1 shows a schematic diagram which was constructed based on a diagram presented in reference [2] to illustrate the relationship between reinforcement particle size (R) and mechanical properties (e.g. yield strength, σ y ). With very fine and shearable intermetallic precipitates, the strengthening effect increases with increase of particle size, as has been well known with precipitation hardening of aluminium alloys. However, these precipitates are not stable at high temperatures, and once the particle size is higher than a certain value, they become non-shearable or “hard” particles, and the strengthening effect becomes weaker with increasing particle size. On the other hand, with particulate reinforced metal–ceramic composites, since the ceramic particles are hard particles, the strengthening effect will increase with decreasing par∗
Corresponding author. Tel.: +64-7-838-4783; fax: +64-7-838-4835. E-mail address: deliang [email protected] (D.L. Zhang).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2003.10.231
ticle size. In the meantime, with a smaller ceramic particle size, the stress concentration level on each particle is lower as there are more particles carrying a given load, making it more difficult to fracture the hard particles. In addition, smaller particles are less prone to having internal defects and thus are more difficult to be fractured. Both of these two factors are favourable for achieving high strength and good fracture toughness. The ceramic particles are stable at high temperatures, so the strengthening effect will hold at high temperatures, leading to a high creep resistance when creep is dislocation controlled. Despite all the expected advantages of metal–ceramic nanocomposites, manufacturing metal–ceramic nanocomposites is a challenging task. The techniques of making metal–ceramic nanocomposites include—in situ growth of ceramic particles through decomposition [3–5], internal oxidation of metal powders [6,7] and mechanical milling of a mixture of metal powder and ceramic powder (e.g. [8–11]). This paper reports and discusses the results of a study on processing Ti3 Al–SiC nanocomposites using high energy mechanical milling and powder metallurgy. This study aims to develop a process for making titanium aluminide based nanocomposites which are expected to have high strength at elevated temperatures, high fracture toughness, high creep resistance and even a higher oxidation resistance than monolithic titanium aluminides.
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D.L. Zhang et al. / Materials Science and Engineering A 375–377 (2004) 911–916
Fig. 1. Schematic diagram illustrating the effect of reinforcement particle size on the mechanical properties of metal–ceramic composites.
2. Experimental technique The starting powders used were Ti powder (99.9 wt.% pure and −100 mesh) and Al powder (99.7 wt.% pure and −325 mesh), and SiC powder (99.9% pure and −400 mesh). The preparation of Ti3 Al–10 vol.% SiC nanocomposite powder was carried out using a SPEX 8000 mill/mixer highenergy ball mill. For each batch of milling, 8 g of the powder mixture was loaded into a hardened steel vial together with four stainless steel bearing balls with a diameter of 12.7 mm in a glove box filled with argon. About 0.1 g of isopropyl alcohol was added into the vial as process control agent. The ball to powder mass ratio was approximately 4:1. The asmilled powders were characterized by using X-ray diffractometry (XRD), optical microscopy and scanning electron microscopy (SEM). The XRD analysis was carried out using a Philip’s X-pert with Cu K␣ radiation. The SEM was performed using a Philip’s XL30 FEG scanning electron microscope and a Hitachi 5400 scanning Electron Microscope equipped with an EDX system. As a comparison, Ti3 Al–20 vol.% SiC nanocomposite powder was also produced using a lower energy planetary mill and a fine SiC powder under argon. The ball to powder mass ratio used was also approximately 4:1. The fine SiC powder was produced by milling the as-received SiC powder for 10 h under argon in the SPEX mill with the same condition as milling the composite powder. The particle size distributions of the as-received coarse powder and the fine powder produced by milling were measured by using a Laser Mastersizer particle size analyser. As-milled com-
posite powders were consolidated by using a conventional hot isostatic pressing (HIP) process.
3. Results and discussion Fig. 2(a) and (b) show the particle size distribution of the as-received coarse SiC powder and the fine powder produced by high energy milling the coarse powder for 10 h, respectively. It was clear that the as-received SiC powder had a mean particle size of about 25 m, while after milling, the mean particle size was reduced to about 0.35 m. In the as-received coarse SiC powder, there was a small fraction of fine powder, while in the finely milled SiC powder, there was a small fraction of coarse powder. The latter was caused by incomplete fracturing of the coarse SiC powder particles during milling. Fig. 3 shows a SEM micrograph of the fine SiC powder. It was evident that the SiC particles in the fine powder were highly equiaxed and the fine particles were heavily agglomerated. The microstructure development of the composite powder particles during milling the Ti3 Al–10 vol.% SiC powder was monitored as shown in Fig. 4. In the early stage of milling, the Ti and Al phases were heavily deformed and extensively cold welded, forming multilayer metal–metal composite structure (Fig. 4(a)), similar to the process of mechanical alloying of Ti and Al reported previously [12,13]. At this point, a small fraction of fractured SiC particles were also incorporated into the powder particles, but the microstructure was far from being uniform. With further milling, this
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Fig. 3. SEM micrograph of the fine SiC powder produced by milling of the coarse SiC powder in the SPEX mill for 10 h.
Fig. 2. Particle size distribution of (a) the as-received coarse SiC powder and (b) the fine SiC powder produced by mechanical milling.
irregular layered structure between Al and Ti became finer and less apparent, and the degree of incorporation of the SiC particles and the metallic phase increased significantly, as shown in Fig. 4(b–d). Simultaneously, SiC particles were extensively fractured, making their average size significantly smaller. By the end of 16 h of milling, the size of most of the SiC particles was in the range of 100–500 nm, as shown
Fig. 4. SEM micrographs of the Ti3 Al–10 vol.% SiC powder particles after different milling durations. (a) 2 h; (b) 4 h; (c) 8 h; and (d) 16 h. Coarse SiC powder was used.
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Fig. 5. Microstructure of powder particles in the Ti3 Al–10 vol.% SiC powder after milling for 16 h. Coarse SiC powder was used.
in Fig. 5. In each powder particle, there were a few SiC particles with size greater than 500 nm, and their size can be reduced by further milling. Fig. 6 shows the XRD patterns of the Ti3 Al–10 vol.% SiC composite powders produced with different milling durations. It was clear that during milling the Al phase gradually alloyed with the Ti phase forming Ti(Al) solid solution, while the grain size of the matrix decreased as shown by the significant broadening of the Ti(Al) peaks. The SiC peaks can still be clearly seen on the pattern of the 16 h milled powder, confirming that the SiC phase did not dissolve as a result of the milling, consistent with the observation based on the SEM examination of the microstructure of the powder particles. As a comparison with the microstructure of the nanocomposite produced using the coarse SiC powder, Fig. 7(a) shows a high magnification SEM micrograph of the powder particles in the 16 h milled Ti3 Al–20 vol.% SiC powder produced using the lower energy planetary mill and the fine SiC powder. The size of the SiC particles in the compos-
Fig. 6. XRD patterns of as-milled Ti3 Al–10 vol.% SiC powders after different milling durations.
Fig. 7. (a) Microstructure of Ti3 Al–20 vol.% SiC nanocomposite powder particles produced by milling for 16 h, using the fine SiC powder; and (b) XRD pattern of the corresponding 16 h milled powder.
ite powder produced using the lower energy planetary mill and fine SiC powder was smaller than that in the composite powder produced by milling for the same time using a high energy ball mill and a coarse SiC powder. An attempt was also made to produce the nanocomposite powder using the planetary mill and the coarse SiC powder, but it was found that the extent of fracturing of the SiC powder particles was very limited, and after 16 h of milling, the average SiC particle size was still well at the micrometer level. XRD pattern of the 16 h planetary milled Ti3 Al–20 vol.% SiC nanocomposite powder (Fig. 7(b)) showed that the Ti(Al) peaks were very broad, the Al peaks disappeared and the SiC peaks were still sharp. This indicates that as a result of the milling, in addition to forming a nanocomposite structure, the metallic matrix became nanocrystalline. Fig. 8 shows SEM micrographs of the bulk Ti3 Al–20 vol.% SiC nanocomposite produced by hipping the 16 h milled nanocomposite powder at 800 ◦ C and under a pressure of 200 MPa. Under this consolidation condition, more than 98% relative density was achieved and the powder was well sintered. Some micrometer sized pores were still evident in the microstructure, as shown in Fig. 8(a). As shown in Fig. 8(b), it appeared that the very fine SiC particles with sizes close to be 100 nm disappeared from the mi-
D.L. Zhang et al. / Materials Science and Engineering A 375–377 (2004) 911–916
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Fig. 8. (a) low magnification and (b) high magnification SEM micrograph of a bulk Ti3 Al–20 vol.% SiC nanocomposite produced by hipping the 16 h milled powder at 800 ◦ C under a pressure of 200 MPa. Fine SiC powder was used.
crostructure, indicating that they may have been dissolved into the matrix. The particles with a larger size were still present in the consolidated material. XRD analysis of the bulk material showed that the reaction between the Ti3 Al matrix and the SiC particles had occurred, forming several different phases including Al2 Ti4 C2 , Ti3 SiC2 and Ti5 Si3 even during hipping at 800 ◦ C. The reaction became more extensive with a higher consolidation temperature. This study demonstrates that Ti3 Al–SiC nanocomposite powder with different SiC volume fractions can be produced by high energy mechanical milling. When a coarse SiC powder is used, it is essential for the impact energy to be high enough to fracture the SiC particles down to nanometer sized level, and therefore high energy mechanical mill such as SPEX mill must be used. On the other hand, when fine SiC powder consisting of nanometer sized SiC particles is used, the impact energy needs only to be sufficient to break the agglomeration of the fine particles. In this case, a lower energy mechanical mill such as the planetary mill can be used. The advantage of using coarse SiC powder and high energy milling is that the matrix/ceramic interfaces are clean as they are produced in situ in an inert atmosphere. If using a fine SiC powder, the interfaces may be contaminated with impurities such as oxygen absorbed on the surface of the fine ceramic particles, and this potential problem could be very severe because of the large surface area to volume ratio of nanoparticles. This means that care must be taken to ensure that the fine ceramic powder is not heavily contaminated during handling. When producing metal–ceramic nanocomposite by fracturing the coarse ceramic powder, it is expected that there exists a lower limit of the ceramic particle size which can be achieved using high energy milling. According to the well established relationship between the defect size (ac ) and the required fracture stress (σ f ) of a ceramic material, as ex-
pressed by the following equation [14]: Kc σf = √ πac
(1)
Where Kc is the fracture toughness of the ceramic. Large ceram particles are likely to contain some internal defects and surface notches. As ceramic materials generally have low fracture toughness, the fracture stress is not very high, making the particles easier to be fractured. However, when the ceramic particles are reduced to a nanometer sized level, the likelihood of having internal defects and surface notches are reduced considerably. In this case, σ f will approach the theoretical strength of the ceramic which is about 1/30 of its Young’s modulus. As an example, the Young’s modulus of SiC is approximately 450 GPa. This means that the impact stress has to be over 15 GPa to fracture a “perfect” SiC particle. It can be imagined that this stress would be very difficult to be achieved with a conventional high energy mechanical mill. So the lower size limit will be defined by the critical point at which the ceramic particles will mostly become “perfect”. Experience and current study suggest that this critical particle size might lie between 50 and 100 nm. In processing metal–ceramic nanocomposite, the difficult task is to avoid reaction between the metallic phase and the ceramic phase. The fact that very fine SiC particles are dissolved during hipping the Ti3 Al–20 vol.% SiC nanocomposite powder at 800 ◦ C illustrate suggest that the level of this difficulty is very high.
4. Conclusions Ti3 Al–10 vol.% SiC and Ti3 Al–20 vol.% SiC nanocomposite powders have been successfully produced using high energy mechanical milling and both coarse and fine SiC powders. With coarse SiC powder, the milling energy has to
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be high enough to fracture the SiC particles in situ. With a fine SiC powder, a lower energy mill is sufficient to produce the nanocomposite powder. During extensive milling, the Ti and Al phases in the matrix are mechanically alloyed, forming nanocrystalline Ti(Al) solution. When producing metal–ceramic nanocomposite powders by high energy milling of metal powders and coarse ceramic powders, there exists a lower limit of the particle size below which reduction of the particle size cannot be achieved using milling. With metal–ceramic nanocomposites, the reaction between the matrix and the ceramic phase is difficult to be avoided. Acknowledgements This work is financially supported by the Foundation for Research, Science and Technology, New Zealand, through a New Economy Research Fund (NERF) grant. Titanox Development Limited, Auckland, New Zealand, is the licenser of the technology for manufacturing titanium based composite powders.
References [1] D.L. Zhang, in: Proceedings of Oxford–Kobe Seminar on Nanocomposite Materials, 11–14th September 2001, Kobe, Japan, pp. 109–119. [2] E. Artzt, Acta Mater. 46 (1998) 5611. [3] K. Tokumitsu, T. Nasu, J. Metastable Nanocryst. Mater. 8 (2000) 562. [4] Z. Chen, Y. Zhang, G.C. Hadjipanayis, Q. Chen, B. Ma, J. Alloys Compd. 287 (1999) 227. [5] J.R. Wilde, A.L. Greer, Mater. Sci. Eng. A 304–306 (2001) 932. [6] D. Preston, N.J. Grant, Trans. TMS-AIME 221 (1961) 164. [7] Z. Shi, D. Wang, J. Mater. Sci. Lett. 17 (1998) 477. [8] D.Y. Ying, D.L. Zhang, Mater. Sci. Eng. A 301 (2001) 90. [9] T.J. Goodwin, S.H. Yoo, P. Matteazzi, J.R. Groza, Nanostruct. Mater. 8 (1997) 559. [10] O. Matteazzi, M. Alcala, Mater. Sci. Eng. A 230 (1997) 161. [11] J. Jose, M.A. Khadar, Acta Mater. 49 (2001) 729. [12] D.L. Zhang, D.Y. Ying, in: D.L. Zhang, K.L. Pickering, X.Y. Xiong (Eds.), Proceedings of the First International Conference on Advanced Materials Processing, Institute of Materials Engineering Australasia, Melbourne, 2000, p. 23. [13] T. Klassen, M. Oehring, R. Boremann, J. Mater. Res. 9 (1994) 47. [14] T.H. Courtney, Mechanical Behaviour of Materials, McGraw-Hill, New York, 1990, p. 408.
Processing Ti3Al–SiC nanocomposites using high energy mechanical milling D.L. Zhang∗ , J. Liang, J. Wu Department of Materials and Process Engineering, University of Waikato, Private Bag 3105, Hamilton, New Zealand
Abstract Ti3 Al–10 vol.% SiC and Ti3 Al–20 vol.% SiC composite powders were produced by high energy mechanical milling of Ti, Al and SiC powders, using both coarse and fine SiC powders. For a coarse SiC powder, the milling energy has to be high enough to fracture the SiC particles in-situ. While for a fine SiC powder, a lower energy mill is sufficient to produce a nanocomposite powder. Additionally, nanocrystalline Ti(Al) solution was formed during milling. A simple theoretical analysis shows that there exists a lower limit of the particle size below which further reduction of the particle size cannot be achieved by milling since the stresses required for further particle refinement cannot be reached during milling. During consolidation of the Ti3 Al–SiC nanocomposite powders, the reaction between the Ti3 Al matrix and SiC is difficult to be avoided. © 2003 Elsevier B.V. All rights reserved. Keywords: Titanium aluminides; Metal–ceramic nanocomposites; Mechanical milling; Powder processing; Powder consolidation
1. Introduction Metal–ceramic nanocomposites are a type of materials which combine metal phases and ceramic phases with at least one dimension of the ceramic or the metallic phase domains being in the nanometer scale. One popular example of metal–ceramic nanocomposites is metal matrix nanocomposites consisting of discrete ceramic nanoparticles embedded in a continuous metal matrix [1]. Fig. 1 shows a schematic diagram which was constructed based on a diagram presented in reference [2] to illustrate the relationship between reinforcement particle size (R) and mechanical properties (e.g. yield strength, σ y ). With very fine and shearable intermetallic precipitates, the strengthening effect increases with increase of particle size, as has been well known with precipitation hardening of aluminium alloys. However, these precipitates are not stable at high temperatures, and once the particle size is higher than a certain value, they become non-shearable or “hard” particles, and the strengthening effect becomes weaker with increasing particle size. On the other hand, with particulate reinforced metal–ceramic composites, since the ceramic particles are hard particles, the strengthening effect will increase with decreasing par∗
Corresponding author. Tel.: +64-7-838-4783; fax: +64-7-838-4835. E-mail address: deliang [email protected] (D.L. Zhang).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2003.10.231
ticle size. In the meantime, with a smaller ceramic particle size, the stress concentration level on each particle is lower as there are more particles carrying a given load, making it more difficult to fracture the hard particles. In addition, smaller particles are less prone to having internal defects and thus are more difficult to be fractured. Both of these two factors are favourable for achieving high strength and good fracture toughness. The ceramic particles are stable at high temperatures, so the strengthening effect will hold at high temperatures, leading to a high creep resistance when creep is dislocation controlled. Despite all the expected advantages of metal–ceramic nanocomposites, manufacturing metal–ceramic nanocomposites is a challenging task. The techniques of making metal–ceramic nanocomposites include—in situ growth of ceramic particles through decomposition [3–5], internal oxidation of metal powders [6,7] and mechanical milling of a mixture of metal powder and ceramic powder (e.g. [8–11]). This paper reports and discusses the results of a study on processing Ti3 Al–SiC nanocomposites using high energy mechanical milling and powder metallurgy. This study aims to develop a process for making titanium aluminide based nanocomposites which are expected to have high strength at elevated temperatures, high fracture toughness, high creep resistance and even a higher oxidation resistance than monolithic titanium aluminides.
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D.L. Zhang et al. / Materials Science and Engineering A 375–377 (2004) 911–916
Fig. 1. Schematic diagram illustrating the effect of reinforcement particle size on the mechanical properties of metal–ceramic composites.
2. Experimental technique The starting powders used were Ti powder (99.9 wt.% pure and −100 mesh) and Al powder (99.7 wt.% pure and −325 mesh), and SiC powder (99.9% pure and −400 mesh). The preparation of Ti3 Al–10 vol.% SiC nanocomposite powder was carried out using a SPEX 8000 mill/mixer highenergy ball mill. For each batch of milling, 8 g of the powder mixture was loaded into a hardened steel vial together with four stainless steel bearing balls with a diameter of 12.7 mm in a glove box filled with argon. About 0.1 g of isopropyl alcohol was added into the vial as process control agent. The ball to powder mass ratio was approximately 4:1. The asmilled powders were characterized by using X-ray diffractometry (XRD), optical microscopy and scanning electron microscopy (SEM). The XRD analysis was carried out using a Philip’s X-pert with Cu K␣ radiation. The SEM was performed using a Philip’s XL30 FEG scanning electron microscope and a Hitachi 5400 scanning Electron Microscope equipped with an EDX system. As a comparison, Ti3 Al–20 vol.% SiC nanocomposite powder was also produced using a lower energy planetary mill and a fine SiC powder under argon. The ball to powder mass ratio used was also approximately 4:1. The fine SiC powder was produced by milling the as-received SiC powder for 10 h under argon in the SPEX mill with the same condition as milling the composite powder. The particle size distributions of the as-received coarse powder and the fine powder produced by milling were measured by using a Laser Mastersizer particle size analyser. As-milled com-
posite powders were consolidated by using a conventional hot isostatic pressing (HIP) process.
3. Results and discussion Fig. 2(a) and (b) show the particle size distribution of the as-received coarse SiC powder and the fine powder produced by high energy milling the coarse powder for 10 h, respectively. It was clear that the as-received SiC powder had a mean particle size of about 25 m, while after milling, the mean particle size was reduced to about 0.35 m. In the as-received coarse SiC powder, there was a small fraction of fine powder, while in the finely milled SiC powder, there was a small fraction of coarse powder. The latter was caused by incomplete fracturing of the coarse SiC powder particles during milling. Fig. 3 shows a SEM micrograph of the fine SiC powder. It was evident that the SiC particles in the fine powder were highly equiaxed and the fine particles were heavily agglomerated. The microstructure development of the composite powder particles during milling the Ti3 Al–10 vol.% SiC powder was monitored as shown in Fig. 4. In the early stage of milling, the Ti and Al phases were heavily deformed and extensively cold welded, forming multilayer metal–metal composite structure (Fig. 4(a)), similar to the process of mechanical alloying of Ti and Al reported previously [12,13]. At this point, a small fraction of fractured SiC particles were also incorporated into the powder particles, but the microstructure was far from being uniform. With further milling, this
D.L. Zhang et al. / Materials Science and Engineering A 375–377 (2004) 911–916
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Fig. 3. SEM micrograph of the fine SiC powder produced by milling of the coarse SiC powder in the SPEX mill for 10 h.
Fig. 2. Particle size distribution of (a) the as-received coarse SiC powder and (b) the fine SiC powder produced by mechanical milling.
irregular layered structure between Al and Ti became finer and less apparent, and the degree of incorporation of the SiC particles and the metallic phase increased significantly, as shown in Fig. 4(b–d). Simultaneously, SiC particles were extensively fractured, making their average size significantly smaller. By the end of 16 h of milling, the size of most of the SiC particles was in the range of 100–500 nm, as shown
Fig. 4. SEM micrographs of the Ti3 Al–10 vol.% SiC powder particles after different milling durations. (a) 2 h; (b) 4 h; (c) 8 h; and (d) 16 h. Coarse SiC powder was used.
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Fig. 5. Microstructure of powder particles in the Ti3 Al–10 vol.% SiC powder after milling for 16 h. Coarse SiC powder was used.
in Fig. 5. In each powder particle, there were a few SiC particles with size greater than 500 nm, and their size can be reduced by further milling. Fig. 6 shows the XRD patterns of the Ti3 Al–10 vol.% SiC composite powders produced with different milling durations. It was clear that during milling the Al phase gradually alloyed with the Ti phase forming Ti(Al) solid solution, while the grain size of the matrix decreased as shown by the significant broadening of the Ti(Al) peaks. The SiC peaks can still be clearly seen on the pattern of the 16 h milled powder, confirming that the SiC phase did not dissolve as a result of the milling, consistent with the observation based on the SEM examination of the microstructure of the powder particles. As a comparison with the microstructure of the nanocomposite produced using the coarse SiC powder, Fig. 7(a) shows a high magnification SEM micrograph of the powder particles in the 16 h milled Ti3 Al–20 vol.% SiC powder produced using the lower energy planetary mill and the fine SiC powder. The size of the SiC particles in the compos-
Fig. 6. XRD patterns of as-milled Ti3 Al–10 vol.% SiC powders after different milling durations.
Fig. 7. (a) Microstructure of Ti3 Al–20 vol.% SiC nanocomposite powder particles produced by milling for 16 h, using the fine SiC powder; and (b) XRD pattern of the corresponding 16 h milled powder.
ite powder produced using the lower energy planetary mill and fine SiC powder was smaller than that in the composite powder produced by milling for the same time using a high energy ball mill and a coarse SiC powder. An attempt was also made to produce the nanocomposite powder using the planetary mill and the coarse SiC powder, but it was found that the extent of fracturing of the SiC powder particles was very limited, and after 16 h of milling, the average SiC particle size was still well at the micrometer level. XRD pattern of the 16 h planetary milled Ti3 Al–20 vol.% SiC nanocomposite powder (Fig. 7(b)) showed that the Ti(Al) peaks were very broad, the Al peaks disappeared and the SiC peaks were still sharp. This indicates that as a result of the milling, in addition to forming a nanocomposite structure, the metallic matrix became nanocrystalline. Fig. 8 shows SEM micrographs of the bulk Ti3 Al–20 vol.% SiC nanocomposite produced by hipping the 16 h milled nanocomposite powder at 800 ◦ C and under a pressure of 200 MPa. Under this consolidation condition, more than 98% relative density was achieved and the powder was well sintered. Some micrometer sized pores were still evident in the microstructure, as shown in Fig. 8(a). As shown in Fig. 8(b), it appeared that the very fine SiC particles with sizes close to be 100 nm disappeared from the mi-
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Fig. 8. (a) low magnification and (b) high magnification SEM micrograph of a bulk Ti3 Al–20 vol.% SiC nanocomposite produced by hipping the 16 h milled powder at 800 ◦ C under a pressure of 200 MPa. Fine SiC powder was used.
crostructure, indicating that they may have been dissolved into the matrix. The particles with a larger size were still present in the consolidated material. XRD analysis of the bulk material showed that the reaction between the Ti3 Al matrix and the SiC particles had occurred, forming several different phases including Al2 Ti4 C2 , Ti3 SiC2 and Ti5 Si3 even during hipping at 800 ◦ C. The reaction became more extensive with a higher consolidation temperature. This study demonstrates that Ti3 Al–SiC nanocomposite powder with different SiC volume fractions can be produced by high energy mechanical milling. When a coarse SiC powder is used, it is essential for the impact energy to be high enough to fracture the SiC particles down to nanometer sized level, and therefore high energy mechanical mill such as SPEX mill must be used. On the other hand, when fine SiC powder consisting of nanometer sized SiC particles is used, the impact energy needs only to be sufficient to break the agglomeration of the fine particles. In this case, a lower energy mechanical mill such as the planetary mill can be used. The advantage of using coarse SiC powder and high energy milling is that the matrix/ceramic interfaces are clean as they are produced in situ in an inert atmosphere. If using a fine SiC powder, the interfaces may be contaminated with impurities such as oxygen absorbed on the surface of the fine ceramic particles, and this potential problem could be very severe because of the large surface area to volume ratio of nanoparticles. This means that care must be taken to ensure that the fine ceramic powder is not heavily contaminated during handling. When producing metal–ceramic nanocomposite by fracturing the coarse ceramic powder, it is expected that there exists a lower limit of the ceramic particle size which can be achieved using high energy milling. According to the well established relationship between the defect size (ac ) and the required fracture stress (σ f ) of a ceramic material, as ex-
pressed by the following equation [14]: Kc σf = √ πac
(1)
Where Kc is the fracture toughness of the ceramic. Large ceram particles are likely to contain some internal defects and surface notches. As ceramic materials generally have low fracture toughness, the fracture stress is not very high, making the particles easier to be fractured. However, when the ceramic particles are reduced to a nanometer sized level, the likelihood of having internal defects and surface notches are reduced considerably. In this case, σ f will approach the theoretical strength of the ceramic which is about 1/30 of its Young’s modulus. As an example, the Young’s modulus of SiC is approximately 450 GPa. This means that the impact stress has to be over 15 GPa to fracture a “perfect” SiC particle. It can be imagined that this stress would be very difficult to be achieved with a conventional high energy mechanical mill. So the lower size limit will be defined by the critical point at which the ceramic particles will mostly become “perfect”. Experience and current study suggest that this critical particle size might lie between 50 and 100 nm. In processing metal–ceramic nanocomposite, the difficult task is to avoid reaction between the metallic phase and the ceramic phase. The fact that very fine SiC particles are dissolved during hipping the Ti3 Al–20 vol.% SiC nanocomposite powder at 800 ◦ C illustrate suggest that the level of this difficulty is very high.
4. Conclusions Ti3 Al–10 vol.% SiC and Ti3 Al–20 vol.% SiC nanocomposite powders have been successfully produced using high energy mechanical milling and both coarse and fine SiC powders. With coarse SiC powder, the milling energy has to
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be high enough to fracture the SiC particles in situ. With a fine SiC powder, a lower energy mill is sufficient to produce the nanocomposite powder. During extensive milling, the Ti and Al phases in the matrix are mechanically alloyed, forming nanocrystalline Ti(Al) solution. When producing metal–ceramic nanocomposite powders by high energy milling of metal powders and coarse ceramic powders, there exists a lower limit of the particle size below which reduction of the particle size cannot be achieved using milling. With metal–ceramic nanocomposites, the reaction between the matrix and the ceramic phase is difficult to be avoided. Acknowledgements This work is financially supported by the Foundation for Research, Science and Technology, New Zealand, through a New Economy Research Fund (NERF) grant. Titanox Development Limited, Auckland, New Zealand, is the licenser of the technology for manufacturing titanium based composite powders.
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