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30%TiB2 composite system during mechanical alloying
Journal of Alloys and Compounds 485 (2009) 724–729
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Formation of nanodispersoids in Fe–Cr–Al/30%TiB2 composite system during mechanical alloying Ritesh Sachan, Jong-Woo Park ∗ Materials Science & Technology Research Division, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea
a r t i c l e
i n f o
Article history: Received 5 February 2009 Received in revised form 8 June 2009 Accepted 9 June 2009 Available online 17 June 2009 Keywords: Metal matrix composite Nanostructured materials Mechanical alloying Electron microprobe analysis TiB2 /Fe–Cr–Al
a b s t r a c t Mechanical alloying is an effective processing technology which allows the synthesis of nanocrystalline composite materials. In this study, a nanocrystalline Fe–Cr–Al/30%TiB2 composite material is synthesized without process control agents by the mechanical alloying process using a planetary high-energy ball mill. Crystalline (or grain) size reduction and dispersion behavior of brittle TiB2 powder during ball milling are investigated together with mechanical alloying behavior of ductile metallic matrix powder for synthesis of the composite. Mechanical alloying between metallic elements is almost completed after approximately 8 h milling. The crystalline size of TiB2 decreases to 36 nm after 48 h of ball milling, while the average particle size of the composite powder increases in comparison with the initial size. Transmission electron microscopy reveals formation of TiB2 nanodispersoids of size around 50 nm after 48 h milling. © 2009 Published by Elsevier B.V.
1. Introduction Titanium diboride (TiB2 ), which is an advanced ceramic material, has shown peculiar physical and mechanical properties such as relatively high melting point, elastic modulus, hardness, strengthto-density ratio, wear resistance and oxidation resistance [1]. Furthermore, pure TiB2 would not be deformed plastically even at very high temperatures due to its intrinsically high Peierls barrier to the dislocation movement [2]. For this reason, TiB2 dispersoids have been used as a reinforcing material for metal matrix composites providing a desirable combination of high-temperature strength and hardness with adequate ductility and fracture toughness. To synthesize metal matrix TiB2 nanocomposites, in situ reaction or self-propagating high-temperature synthesis processes have been used widely. For example, Tanaka et al. made a high-modulus steel reinforced with high volume (10–45 vol.%) of 200–800 nm TiB2 particles through an in situ reaction of Fe–Ti and Fe–B powders [3]. Finer (30–50 nm) TiB2 particles (up to 57 vol.%) in Cu matrix were fabricated by SHS incorporated with high-energy ball milling of Ti, B and Cu powders [4–7]. Much finer (around 17 nm) TiB2 particles in Fe–Cr–Ni matrix were synthesized by mechanically activated self-sustaining reaction during high-energy ball milling of Ti/BN/Fe–Cr–Ni powder mixture [8]. For the synthesis of Cu/TiB2 composites containing low volume (below 3 wt.%) nanoparticles (around 20 nm), in situ reaction technique was used in a melt
∗ Corresponding author. Tel.: +82 53 856 5233; fax: +82 53 856 5228. E-mail address: [email protected] (J.-W. Park). 0925-8388/$ – see front matter © 2009 Published by Elsevier B.V. doi:10.1016/j.jallcom.2009.06.063
[9] or during rapid solidification [10]. Grain refinement of TiB2 to nano-scale during high-energy ball milling was reported for the TiB2 /Al composite systems. Nanocrystalline TiB2 was synthesized in one case by mechanically activated reaction between Ti, Al and B powders [11], and in another case between Al, TiO2 and B2 O3 powders [12]. Fe–Cr–Al/TiB2 composite system will be investigated in the present study because of its promising character for high-temperature applications. Fe–Cr–Al alloy, which is a ferritic stainless steel, is being used in various applications up to temperature as high as 1300 ◦ C due to their excellent oxidation resistance [13], and is selected as the metallic matrix or binder material in this study. The purpose of the current study is to fabricate metal matrix nanocomposite powder containing high volume fraction of nanocrystalline TiB2 directly from coarse powders of raw materials i.e. Fe, Cr and Al powders, using only single process of high-energy ball milling without in situ reaction. Nanocrystallization and dispersion behavior of brittle ceramic powder during ball milling are investigated together with mechanical alloying behavior of ductile metallic powder for synthesis of a Fe–Cr–Al/30%TiB2 nanocomposite. 2. Experimental procedures In the present study, TiB2 , 99% Fe, 99.8% Cr and 99.9% Al with average particle size of 6, 3, 10 and 4 m, respectively, were used as the starting powder. Fe, Cr and Al powders were physically mixed in a composition of Fe–20%Cr–5%Al (percentage by weight is used unless otherwise specified). This powder was further mixed with TiB2 to make a physical mixture of composition Fe–Cr–Al/30%TiB2 . These powder mixtures were prepared by mixing for 20 min in a vibration mixer (Retsch MM-201).
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Fig. 1. SEM images of Fe–Cr–Al/30%TiB2 composite powder after (a) 0 h, (b) 1 h, (c) 8 h and (d) 48 h of ball milling.
A mixture of TiB2 /Fe–Cr–Al of 10 g was charged into a tool steel vial (125 ml internal volume) with an O-ring seal under an Ar atmosphere to avoid oxidation during the milling process. Mechanical alloying was performed at room temperature up to 48 h with a high-energy planetary mill (PM200) at 500 rpm using 11.9 mm diameter chrome steel balls, and the ball-to-powder wt. ratio was maintained approximately 20:1. As-milled powders were characterized by X-ray diffraction (XRD) using CuK␣ radiation (Bruker D8 Advance). The crystalline size was calculated by Sherrer’s formula, taking all values of peak broadening at second major peak in the XRD patterns. Since Fe, Cr and TiB2 show the first major peaks at the same angle 2Â of 44.4◦ , to avoid the ambiguity, the peaks at 2Â around 34.1◦ (the second major peak for TiB2 corresponding to {1 0 0} diffraction) were used for analyzing the crystalline size reduction only in case of TiB2 . The morphology of nanocomposite powders was analyzed by scanning electron microscopy (SEM) operating at 20 kV (Hitachi S2400). Further analysis of average particle size of nanocomposite was carried out by a laser diffraction particle size analyzer (Beckman Coulter LS230). As-polished specimen cross-sections were observed by SEM (JXA-8500F) in back scattered electron (BSE) contrast mode. Elemental mapping of the cross-section of the nanocomposite powder was also performed under thermal field emission electron probe microanalyzer combined with high-resolution SEM (JXA-8500F). The nanostructures of as-milled powders were analyzed by highresolution transmission electron microscopy operating at 300 kV (JEOL 3000F). The TEM specimens were prepared by ion milling (Fishione Ion Miller Model 1010) of the powder sample bonded by resin.
Fig. 2. The distribution curve of the starting material shows a single major peak at the particle size of about 5 m, together with a minor peak around 20 m. One hour of ball milling decreases the number of fine particles, and increases coarse particles, resulting in double major peaks at the particle size around 8 and 22 m, respectively, expanding and broadening the curve to a much bigger size region. This bimodal particle size distribution is consistent with the previous SEM observation (Fig. 1b), demonstrating that the second major peak is related with the agglomerates formed by ball milling. After 8 and 48 h of milling, a single major peak appears again. Very fine particles disappear, and the number of medium size particles around 10 m increases with milling time, while big particles or agglomerates around a few tens m diminish gradually. Fig. 3 shows the XRD patterns of the ball milled Fe–Cr–Al/30%TiB2 powder at different milling intervals. Peaks corresponding to TiB2 (C32 hexagonal), Al (FCC), Fe (BCC) and Cr (BCC) are observed in the diffraction pattern of the starting powder
3. Results Fig. 1 shows the SEM surface morphology of powder particles at different stages of ball milling for synthesis of the Fe–Cr–Al/30%TiB2 composite powder. The starting material, which is a mixture of TiB2 , Fe, Cr and Al powder, exhibits almost globular particles of various size (Fig. 1a). After ball milling starts, some big particles or agglomerates are noted among small particles (Fig. 1b). The agglomerates have rough surface, and the shape of the small particles becomes somewhat irregular. With further milling, the small particles grow in size, and the average particle size increases. After long term milling of 48 h, most of big agglomerates and fine particles have disappeared, leading to more uniform distribution in particle size (Fig. 1d). In addition, the shape of the particles becomes more smooth and granular. Particle size distribution of the composite powder for different milling time has been investigated quantitatively as shown in
Fig. 2. Particle size distribution in ball milled powder samples.
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Fig. 3. XRD patterns for samples milled for (a) 0 h, (b) 1 h, (c) 8 h, (d) 24 h and (e) 48 h.
mixture (Fig. 3a). Peaks corresponding to Al disappear after 1 h of milling (Fig. 3b). Possible explanations of Al peak disappearance could be either Al element is alloyed with Fe and Cr, or XRD is unable to detect as amount of Al is comparatively small. Gradual broadening of the peaks takes place together with weakening of the peak intensity on further milling, which is mainly related with the refinement of crystalline size during ball milling [16,17]. This XRD analysis exhibits that the finally ball milled powder consists of a fine crystalline mixture of a Fe–Cr–Al metallic solid solution (BCC) and a TiB2 ceramic component. Mechanical alloying behavior of Fe–Cr–Al metallic elements and the TiB2 ceramic component has also been investigated by back scattered images and elemental mapping on the cross-section of the composite powder. The back scattered images of the compos-
ite powder after 1, 8, 48 h of milling are shown in Fig. 4 together with the corresponding elemental mappings of Ti and B. In the back scattered images, agglomerated powder particles consist of a large number of small dark spots imbedded in the bright matrix area. As darker spots belong to the lighter atomic number elements in the back scattered image, Ti-rich constituents in the agglomerate are observed as dark spots, whereas Fe–Cr–Al matrix phase is the bright area. The same Ti-rich regions are also observed in the elemental mapping images of Ti. Elemental concentration at different areas is color-coded according to the legend shown in the each image. Presence of B-rich regions is evident at the same positions as Ti-rich regions in Fig. 4g–i, though the detection of B being a light element is comparatively difficult. Furthermore, the Ti-rich region is believed to be TiB2 phase as observed in the above XRD analysis earlier. Moreover the sole presence of chemically stable TiB2 phase is believed as Lu et al. [11,12] have reported that TiB2 remains thermodynamically stable without any decomposition once formed by mechanically activated self-sustaining reaction. It is noted from this analysis that the TiB2 particles in the agglomerates are getting smaller and more finely distributed in the Fe–Cr–Al matrix with subsequent ball milling. The TiB2 particles are dispersed in such a fine manner after 48 h of ball milling that it is hard to distinguish particles clearly by EPMA in this magnification with 97.6 nm resolution. Elemental mapping of matrix components, Fe, Cr and Al, of the composite powder are also investigated after 1, 8 and 48 h of ball milling respectively as shown in Fig. 5. Fe-, Cr- and Al-rich regions disappear with subsequent ball milling, which suggests that alloying has taken place between the constituent elements (Fe, Cr and Al) of matrix phase. It is confirmed from Fig. 5 that alloying between metallic elements had been almost completed after approximately 8 h, as no significant change is observed in the elemental mapping images of 8 and 48 h as-milled composite powders. This result for completion of alloying by EPMA is more precise than the earlier reported XRD results in Fig. 3 where dissolution of Al and Cr in Fe was observed after 1 h of ball milling.
Fig. 4. Back scattered images, and elemental maps for Ti and B of composite powder milled for (a, d, g) 1 h, (b, e, h) 8 h and (c, f, i) 48 h.
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Fig. 5. Elemental mapping images of Fe, Cr and Al by EPMA for the composite powder milled for 1, 8 and 48 h, respectively.
TEM analysis of Fe–Cr–Al/30%TiB2 composite particles ball milled for 48 h is performed for identification of TiB2 dispersoids in the Fe–Cr–Al metallic matrix as shown in Fig. 6. A dispersoid particle having size around 60 nm is shown in Fig. 6a, and the selected area diffraction pattern (SADP) of the encircled area is in Fig. 6b. The diffraction pattern corresponds to (1 1¯ 1) zone axis reflection of TiB2 , confirming that the particle is actually TiB2 nanodispersoid. A lot of nanodispersoids of size around 40–60 nm are observed in the metallic matrix of the bright field image of Fig. 6c. The corresponding dark field image, which is taken with respect to TiB2 (1 0 1) reflection, exhibits that TiB2 crystallites are visible having size around 30–40 nm. 4. Discussion It has been known that a critical balance between cold welding and fracturing is necessary for successful mechanical alloying, which enables powder particles to be always in contact with each other with atomically clean surfaces, minimizing diffusion distance [14]. Soft metallic materials such as aluminum are hardly alloyed mechanically because the process is hindered by excessive cold welding of the powder particles, preventing them from fracturing [15]. Hence, surface or process control agents are often used to annul the forces of cold welding during mechanical alloying. For the current material system consisting of hard ceramic and soft metallic powder, however, mechanical alloying has been completed successfully without surface agents, since agglomeration by cold welding occurs actively only in the early stage of milling, becoming more balanced with fracturing in the later stage. Agglomeration and fracturing behavior of the present composite system is seen more clearly in Fig. 7 where the mean and median particle size of the composite powder material is obtained from Fig. 2. The mean and median particle size first increases and then decreases after attaining a peak with increasing milling time. The whole process of ball milling can be divided into 2 stages, stage I
and stage II, having a boundary along the peak particle size in Fig. 7. It has been known that the degree of cold welding is dependent on the ductility and the ability to cold welding of the powder [14]. In our composite system, cold welding is a dominant mechanism in the stage I, which is responsible for the increase in powder particle size. In the stage II, however, a fracturing process gradually becomes more dominant over agglomeration process with increasing milling time, probably owing to solid solution hardening by mechanical alloying as well as dispersion hardening of the composite powder with milling time, which results in the decrease in the agglomerated particle size. It may not be just a coincidence in the current study that the stage II starts when mechanical alloying between metallic elements has been completed. Humail et al. [16] reported a similar behavior of first domination of cold welding and then fracturing during the ball milling process in the system of 95%W–3.5%Ni–l.5%Fe where hard W powder is much larger in volume than soft Ni and Fe powders. However, they added 1% stearic acid as a process control agent, differently from us using no agent. In Fig. 7, variation in crystalline size of TiB2 with milling time, which is measured from peak broadening in Fig. 3, is also shown along with the particle size of the composite. The crystalline size of TiB2 reduces rapidly in the stage I, and at the same time the average particle size of the composite increases fast as well. Therefore, it might be suggested that the agglomeration between powder particles assists the fragmentation of TiB2 in this stage. Agglomeration between different particles helps in reducing the scattering of particles, thus increases the effective energy transfer from the milling balls. It is also notable in the stage I that reduction rate in the crystalline size of TiB2 becomes slow with gradual reduction in the agglomeration rate, supporting the correlation between the powder agglomeration and the TiB2 fragmentation. In the stage II, the crystalline size of TiB2 continues to decrease at the same reduced rate as the last of the stage I, and the average reduction rate of the crystalline size is much lower than that in the stage I. The crystalline size reduction in the stage II is
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Fig. 6. (a) TEM bright field image, (b) SADP demonstrating TiB2 , (c) bright field image of composite powder showing nanodispersoids, and (d) corresponding dark field image.
accompanied by the reduction in powder particle size. After 48 h milling, the mean or median particle size is around a few hundred times of the crystalline size of TiB2 particles. The crystalline size (36 nm) of TiB2 particles calculated by XRD is in good agreement with the TiB2 crystallites size (30–40 nm) observed in the dark field TEM images, and comparable with the TiB2 particles size (40–60 nm) in the bright field images after 48 h milling. This implies that most of the fractured TiB2 particles are likely to be single or bi-crystal.
Since there is no mechano-chemical reaction involved in the current work, the nanocrystallites of TiB2 are achieved only by mechanical milling using TiB2 , Fe, Cr and Al as starting powder. In addition, TiB2 nanodispersoids are uniformly distributed in the homogeneous ferritic matrix as shown in Figs. 4–6. The TiB2 nanoparticles synthesized by our group have relatively larger crystalline size than the one reported by Tu et al. [9] and Guo et al. [10] having TiB2 of 20 nm crystalline size in Cu/TiB2 system, where they have used in situ reaction in a melt or rapid solidification route. Kim et al. [8] also reported finer crystalline size of TiB2 (around 17 nm) using mechanically activated self-sustaining reaction in TiB2 /TiN/Fe–Cr–Al system than the one reported in present study. However, the observed crystalline size in the present study is considerably finer than the one reported by Shim et al. [17] who synthesized around 100 nm sized TiB2 particle along with TiN, starting with Ti and BN powder particles through self-propagating high-temperature reaction as well as to that of Tanaka et al. [3] who reported 2–3 m sized TiB2 dispersoids in Fe–Cr–Ti ferrite matrix by conventional powder metallurgy route and 200–800 nm dispersed TiB2 synthesized through in situ reaction route. It is also noteworthy in the present study that nanocomposites containing TiB2 dispersoids can be obtained simply by high-energy ball milling without any in situ reaction and process control agents. 5. Conclusions
Fig. 7. Variation in particle size of the Fe–Cr–Al/30%TiB2 composite and crystalline size of TiB2 with milling time.
Nanocrystalline Fe–Cr–Al/30%TiB2 composite powder has been synthesized without process control agents from a micron sized mixture of metallic (Fe, Cr and Al) and ceramic (TiB2 ) powders by
R. Sachan, J.-W. Park / Journal of Alloys and Compounds 485 (2009) 724–729
a high-energy ball milling process. Fine TiB2 dispersoids of size around 50 nm or lower are observed in the metallic matrix of Fe–Cr–Al solid solution. In addition to it, the crystalline size of TiB2 particles decreases with longer ball milling duration, resulting in an average size of 36 nm after 48 h of ball milling. The mean particle size of the composite powder, however, increases three times due to agglomeration of powders in the early stage of ball milling, and then reduces gradually with further milling. Acknowledgements This research was supported by an Institutional R&D Program of KIST. We wish to thank Dr. J.H. Shim and Dr. Y.W. Cho for their kind discussion and the supply of sample preparation facilities. We are also very thankful to Mr. S.I. Baik and Dr. Y.W. Kim at Seoul National University for transmission electron microscopic analysis work. References [1] R.G. Munro, J. Am. Ceram. Soc. 80 (1997) 1919–1928.
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[2] J. Ramberg, C. Woolfe, W. Williams, J. Am. Ceram. Soc. 68 (1985) C78– C79. [3] K. Tanaka, T. Oshima, T. Saito, Tetsu-to-Hagane 84 (10) (1998) 747– 754. [4] Y.S. Kwon, D.V. Dudina, M.A. Korchagin, O.I. Lomovsky, J. Mater. Sci. 39 (2004) 5325–5331. [5] J.S. Kim, Y.S. Kwon, D.V. Dudina, O.I. Lomovsky, M.A. Korchangin, V.I. Mali, J. Mater. Sci. 40 (13) (2005) 3491–3495. [6] J.S. Kim, Y.S. Kwon, O.I. Lomovsky, M.A. Korchagin, V.I. Mali, D.V. Dudina, Mater. Lett. 60 (2006) 3723–3726. [7] J.S. Kim, Y.S. Kwon, O.I. Lomovsky, D.V. Dudina, V.F. Kosarev, S.V. Klinkov, D.H. Kwon, Compos. Sci. Technol. 67 (2007) 2292–2296. [8] J.W. Kim, H.S. Chung, J.H. Shim, J.P. Ahn, Y.W. Cho, K.H. Oh, J. Alloys Compd. 422 (2006) 62–66. [9] J.P. Tu, W. Rong, S.Y. Guo, Y.Z. Yang, Wear 255 (2003) 832–835. [10] M.X. Guo, M.P. Wang, K. Shen, L.F. Cao, Z. Li, Z. Zhang, J. Alloys Compd. 460 (2008) 585–589. [11] L. Lu, M.O. Lai, H.Y. Wang, J. Mater. Sci. 35 (2000) 241–248. [12] L. Lu, M.O. Lai, Y. Su, H.L. Teo, C.F. Feng, Scripta Mater. 45 (2001) 1017– 1023. [13] C. Badini, F. Laurella, Surf. Coat. Technol. 135 (2001) 291–298. [14] L. Lu, M.O. Lai, Mechanical Alloying, Kluwer Academic Publishers, Massachusetts, 1998. [15] P.S. Gilman, W.D. Nix, Metall. Trans. A 12A (1981) 813–824. [16] I.S. Humail, X. Qu, C. Jia, M. Qin, X. He, J. Univ. Sci. Technol. Beijing 13 (5) (2006) 442–445. [17] J.H. Shim, J.S. Byun, Y.W. Cho, Scripta Mater. 47 (2002) 493–497.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Formation of nanodispersoids in Fe–Cr–Al/30%TiB2 composite system during mechanical alloying Ritesh Sachan, Jong-Woo Park ∗ Materials Science & Technology Research Division, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea
a r t i c l e
i n f o
Article history: Received 5 February 2009 Received in revised form 8 June 2009 Accepted 9 June 2009 Available online 17 June 2009 Keywords: Metal matrix composite Nanostructured materials Mechanical alloying Electron microprobe analysis TiB2 /Fe–Cr–Al
a b s t r a c t Mechanical alloying is an effective processing technology which allows the synthesis of nanocrystalline composite materials. In this study, a nanocrystalline Fe–Cr–Al/30%TiB2 composite material is synthesized without process control agents by the mechanical alloying process using a planetary high-energy ball mill. Crystalline (or grain) size reduction and dispersion behavior of brittle TiB2 powder during ball milling are investigated together with mechanical alloying behavior of ductile metallic matrix powder for synthesis of the composite. Mechanical alloying between metallic elements is almost completed after approximately 8 h milling. The crystalline size of TiB2 decreases to 36 nm after 48 h of ball milling, while the average particle size of the composite powder increases in comparison with the initial size. Transmission electron microscopy reveals formation of TiB2 nanodispersoids of size around 50 nm after 48 h milling. © 2009 Published by Elsevier B.V.
1. Introduction Titanium diboride (TiB2 ), which is an advanced ceramic material, has shown peculiar physical and mechanical properties such as relatively high melting point, elastic modulus, hardness, strengthto-density ratio, wear resistance and oxidation resistance [1]. Furthermore, pure TiB2 would not be deformed plastically even at very high temperatures due to its intrinsically high Peierls barrier to the dislocation movement [2]. For this reason, TiB2 dispersoids have been used as a reinforcing material for metal matrix composites providing a desirable combination of high-temperature strength and hardness with adequate ductility and fracture toughness. To synthesize metal matrix TiB2 nanocomposites, in situ reaction or self-propagating high-temperature synthesis processes have been used widely. For example, Tanaka et al. made a high-modulus steel reinforced with high volume (10–45 vol.%) of 200–800 nm TiB2 particles through an in situ reaction of Fe–Ti and Fe–B powders [3]. Finer (30–50 nm) TiB2 particles (up to 57 vol.%) in Cu matrix were fabricated by SHS incorporated with high-energy ball milling of Ti, B and Cu powders [4–7]. Much finer (around 17 nm) TiB2 particles in Fe–Cr–Ni matrix were synthesized by mechanically activated self-sustaining reaction during high-energy ball milling of Ti/BN/Fe–Cr–Ni powder mixture [8]. For the synthesis of Cu/TiB2 composites containing low volume (below 3 wt.%) nanoparticles (around 20 nm), in situ reaction technique was used in a melt
∗ Corresponding author. Tel.: +82 53 856 5233; fax: +82 53 856 5228. E-mail address: [email protected] (J.-W. Park). 0925-8388/$ – see front matter © 2009 Published by Elsevier B.V. doi:10.1016/j.jallcom.2009.06.063
[9] or during rapid solidification [10]. Grain refinement of TiB2 to nano-scale during high-energy ball milling was reported for the TiB2 /Al composite systems. Nanocrystalline TiB2 was synthesized in one case by mechanically activated reaction between Ti, Al and B powders [11], and in another case between Al, TiO2 and B2 O3 powders [12]. Fe–Cr–Al/TiB2 composite system will be investigated in the present study because of its promising character for high-temperature applications. Fe–Cr–Al alloy, which is a ferritic stainless steel, is being used in various applications up to temperature as high as 1300 ◦ C due to their excellent oxidation resistance [13], and is selected as the metallic matrix or binder material in this study. The purpose of the current study is to fabricate metal matrix nanocomposite powder containing high volume fraction of nanocrystalline TiB2 directly from coarse powders of raw materials i.e. Fe, Cr and Al powders, using only single process of high-energy ball milling without in situ reaction. Nanocrystallization and dispersion behavior of brittle ceramic powder during ball milling are investigated together with mechanical alloying behavior of ductile metallic powder for synthesis of a Fe–Cr–Al/30%TiB2 nanocomposite. 2. Experimental procedures In the present study, TiB2 , 99% Fe, 99.8% Cr and 99.9% Al with average particle size of 6, 3, 10 and 4 m, respectively, were used as the starting powder. Fe, Cr and Al powders were physically mixed in a composition of Fe–20%Cr–5%Al (percentage by weight is used unless otherwise specified). This powder was further mixed with TiB2 to make a physical mixture of composition Fe–Cr–Al/30%TiB2 . These powder mixtures were prepared by mixing for 20 min in a vibration mixer (Retsch MM-201).
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Fig. 1. SEM images of Fe–Cr–Al/30%TiB2 composite powder after (a) 0 h, (b) 1 h, (c) 8 h and (d) 48 h of ball milling.
A mixture of TiB2 /Fe–Cr–Al of 10 g was charged into a tool steel vial (125 ml internal volume) with an O-ring seal under an Ar atmosphere to avoid oxidation during the milling process. Mechanical alloying was performed at room temperature up to 48 h with a high-energy planetary mill (PM200) at 500 rpm using 11.9 mm diameter chrome steel balls, and the ball-to-powder wt. ratio was maintained approximately 20:1. As-milled powders were characterized by X-ray diffraction (XRD) using CuK␣ radiation (Bruker D8 Advance). The crystalline size was calculated by Sherrer’s formula, taking all values of peak broadening at second major peak in the XRD patterns. Since Fe, Cr and TiB2 show the first major peaks at the same angle 2Â of 44.4◦ , to avoid the ambiguity, the peaks at 2Â around 34.1◦ (the second major peak for TiB2 corresponding to {1 0 0} diffraction) were used for analyzing the crystalline size reduction only in case of TiB2 . The morphology of nanocomposite powders was analyzed by scanning electron microscopy (SEM) operating at 20 kV (Hitachi S2400). Further analysis of average particle size of nanocomposite was carried out by a laser diffraction particle size analyzer (Beckman Coulter LS230). As-polished specimen cross-sections were observed by SEM (JXA-8500F) in back scattered electron (BSE) contrast mode. Elemental mapping of the cross-section of the nanocomposite powder was also performed under thermal field emission electron probe microanalyzer combined with high-resolution SEM (JXA-8500F). The nanostructures of as-milled powders were analyzed by highresolution transmission electron microscopy operating at 300 kV (JEOL 3000F). The TEM specimens were prepared by ion milling (Fishione Ion Miller Model 1010) of the powder sample bonded by resin.
Fig. 2. The distribution curve of the starting material shows a single major peak at the particle size of about 5 m, together with a minor peak around 20 m. One hour of ball milling decreases the number of fine particles, and increases coarse particles, resulting in double major peaks at the particle size around 8 and 22 m, respectively, expanding and broadening the curve to a much bigger size region. This bimodal particle size distribution is consistent with the previous SEM observation (Fig. 1b), demonstrating that the second major peak is related with the agglomerates formed by ball milling. After 8 and 48 h of milling, a single major peak appears again. Very fine particles disappear, and the number of medium size particles around 10 m increases with milling time, while big particles or agglomerates around a few tens m diminish gradually. Fig. 3 shows the XRD patterns of the ball milled Fe–Cr–Al/30%TiB2 powder at different milling intervals. Peaks corresponding to TiB2 (C32 hexagonal), Al (FCC), Fe (BCC) and Cr (BCC) are observed in the diffraction pattern of the starting powder
3. Results Fig. 1 shows the SEM surface morphology of powder particles at different stages of ball milling for synthesis of the Fe–Cr–Al/30%TiB2 composite powder. The starting material, which is a mixture of TiB2 , Fe, Cr and Al powder, exhibits almost globular particles of various size (Fig. 1a). After ball milling starts, some big particles or agglomerates are noted among small particles (Fig. 1b). The agglomerates have rough surface, and the shape of the small particles becomes somewhat irregular. With further milling, the small particles grow in size, and the average particle size increases. After long term milling of 48 h, most of big agglomerates and fine particles have disappeared, leading to more uniform distribution in particle size (Fig. 1d). In addition, the shape of the particles becomes more smooth and granular. Particle size distribution of the composite powder for different milling time has been investigated quantitatively as shown in
Fig. 2. Particle size distribution in ball milled powder samples.
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Fig. 3. XRD patterns for samples milled for (a) 0 h, (b) 1 h, (c) 8 h, (d) 24 h and (e) 48 h.
mixture (Fig. 3a). Peaks corresponding to Al disappear after 1 h of milling (Fig. 3b). Possible explanations of Al peak disappearance could be either Al element is alloyed with Fe and Cr, or XRD is unable to detect as amount of Al is comparatively small. Gradual broadening of the peaks takes place together with weakening of the peak intensity on further milling, which is mainly related with the refinement of crystalline size during ball milling [16,17]. This XRD analysis exhibits that the finally ball milled powder consists of a fine crystalline mixture of a Fe–Cr–Al metallic solid solution (BCC) and a TiB2 ceramic component. Mechanical alloying behavior of Fe–Cr–Al metallic elements and the TiB2 ceramic component has also been investigated by back scattered images and elemental mapping on the cross-section of the composite powder. The back scattered images of the compos-
ite powder after 1, 8, 48 h of milling are shown in Fig. 4 together with the corresponding elemental mappings of Ti and B. In the back scattered images, agglomerated powder particles consist of a large number of small dark spots imbedded in the bright matrix area. As darker spots belong to the lighter atomic number elements in the back scattered image, Ti-rich constituents in the agglomerate are observed as dark spots, whereas Fe–Cr–Al matrix phase is the bright area. The same Ti-rich regions are also observed in the elemental mapping images of Ti. Elemental concentration at different areas is color-coded according to the legend shown in the each image. Presence of B-rich regions is evident at the same positions as Ti-rich regions in Fig. 4g–i, though the detection of B being a light element is comparatively difficult. Furthermore, the Ti-rich region is believed to be TiB2 phase as observed in the above XRD analysis earlier. Moreover the sole presence of chemically stable TiB2 phase is believed as Lu et al. [11,12] have reported that TiB2 remains thermodynamically stable without any decomposition once formed by mechanically activated self-sustaining reaction. It is noted from this analysis that the TiB2 particles in the agglomerates are getting smaller and more finely distributed in the Fe–Cr–Al matrix with subsequent ball milling. The TiB2 particles are dispersed in such a fine manner after 48 h of ball milling that it is hard to distinguish particles clearly by EPMA in this magnification with 97.6 nm resolution. Elemental mapping of matrix components, Fe, Cr and Al, of the composite powder are also investigated after 1, 8 and 48 h of ball milling respectively as shown in Fig. 5. Fe-, Cr- and Al-rich regions disappear with subsequent ball milling, which suggests that alloying has taken place between the constituent elements (Fe, Cr and Al) of matrix phase. It is confirmed from Fig. 5 that alloying between metallic elements had been almost completed after approximately 8 h, as no significant change is observed in the elemental mapping images of 8 and 48 h as-milled composite powders. This result for completion of alloying by EPMA is more precise than the earlier reported XRD results in Fig. 3 where dissolution of Al and Cr in Fe was observed after 1 h of ball milling.
Fig. 4. Back scattered images, and elemental maps for Ti and B of composite powder milled for (a, d, g) 1 h, (b, e, h) 8 h and (c, f, i) 48 h.
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Fig. 5. Elemental mapping images of Fe, Cr and Al by EPMA for the composite powder milled for 1, 8 and 48 h, respectively.
TEM analysis of Fe–Cr–Al/30%TiB2 composite particles ball milled for 48 h is performed for identification of TiB2 dispersoids in the Fe–Cr–Al metallic matrix as shown in Fig. 6. A dispersoid particle having size around 60 nm is shown in Fig. 6a, and the selected area diffraction pattern (SADP) of the encircled area is in Fig. 6b. The diffraction pattern corresponds to (1 1¯ 1) zone axis reflection of TiB2 , confirming that the particle is actually TiB2 nanodispersoid. A lot of nanodispersoids of size around 40–60 nm are observed in the metallic matrix of the bright field image of Fig. 6c. The corresponding dark field image, which is taken with respect to TiB2 (1 0 1) reflection, exhibits that TiB2 crystallites are visible having size around 30–40 nm. 4. Discussion It has been known that a critical balance between cold welding and fracturing is necessary for successful mechanical alloying, which enables powder particles to be always in contact with each other with atomically clean surfaces, minimizing diffusion distance [14]. Soft metallic materials such as aluminum are hardly alloyed mechanically because the process is hindered by excessive cold welding of the powder particles, preventing them from fracturing [15]. Hence, surface or process control agents are often used to annul the forces of cold welding during mechanical alloying. For the current material system consisting of hard ceramic and soft metallic powder, however, mechanical alloying has been completed successfully without surface agents, since agglomeration by cold welding occurs actively only in the early stage of milling, becoming more balanced with fracturing in the later stage. Agglomeration and fracturing behavior of the present composite system is seen more clearly in Fig. 7 where the mean and median particle size of the composite powder material is obtained from Fig. 2. The mean and median particle size first increases and then decreases after attaining a peak with increasing milling time. The whole process of ball milling can be divided into 2 stages, stage I
and stage II, having a boundary along the peak particle size in Fig. 7. It has been known that the degree of cold welding is dependent on the ductility and the ability to cold welding of the powder [14]. In our composite system, cold welding is a dominant mechanism in the stage I, which is responsible for the increase in powder particle size. In the stage II, however, a fracturing process gradually becomes more dominant over agglomeration process with increasing milling time, probably owing to solid solution hardening by mechanical alloying as well as dispersion hardening of the composite powder with milling time, which results in the decrease in the agglomerated particle size. It may not be just a coincidence in the current study that the stage II starts when mechanical alloying between metallic elements has been completed. Humail et al. [16] reported a similar behavior of first domination of cold welding and then fracturing during the ball milling process in the system of 95%W–3.5%Ni–l.5%Fe where hard W powder is much larger in volume than soft Ni and Fe powders. However, they added 1% stearic acid as a process control agent, differently from us using no agent. In Fig. 7, variation in crystalline size of TiB2 with milling time, which is measured from peak broadening in Fig. 3, is also shown along with the particle size of the composite. The crystalline size of TiB2 reduces rapidly in the stage I, and at the same time the average particle size of the composite increases fast as well. Therefore, it might be suggested that the agglomeration between powder particles assists the fragmentation of TiB2 in this stage. Agglomeration between different particles helps in reducing the scattering of particles, thus increases the effective energy transfer from the milling balls. It is also notable in the stage I that reduction rate in the crystalline size of TiB2 becomes slow with gradual reduction in the agglomeration rate, supporting the correlation between the powder agglomeration and the TiB2 fragmentation. In the stage II, the crystalline size of TiB2 continues to decrease at the same reduced rate as the last of the stage I, and the average reduction rate of the crystalline size is much lower than that in the stage I. The crystalline size reduction in the stage II is
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Fig. 6. (a) TEM bright field image, (b) SADP demonstrating TiB2 , (c) bright field image of composite powder showing nanodispersoids, and (d) corresponding dark field image.
accompanied by the reduction in powder particle size. After 48 h milling, the mean or median particle size is around a few hundred times of the crystalline size of TiB2 particles. The crystalline size (36 nm) of TiB2 particles calculated by XRD is in good agreement with the TiB2 crystallites size (30–40 nm) observed in the dark field TEM images, and comparable with the TiB2 particles size (40–60 nm) in the bright field images after 48 h milling. This implies that most of the fractured TiB2 particles are likely to be single or bi-crystal.
Since there is no mechano-chemical reaction involved in the current work, the nanocrystallites of TiB2 are achieved only by mechanical milling using TiB2 , Fe, Cr and Al as starting powder. In addition, TiB2 nanodispersoids are uniformly distributed in the homogeneous ferritic matrix as shown in Figs. 4–6. The TiB2 nanoparticles synthesized by our group have relatively larger crystalline size than the one reported by Tu et al. [9] and Guo et al. [10] having TiB2 of 20 nm crystalline size in Cu/TiB2 system, where they have used in situ reaction in a melt or rapid solidification route. Kim et al. [8] also reported finer crystalline size of TiB2 (around 17 nm) using mechanically activated self-sustaining reaction in TiB2 /TiN/Fe–Cr–Al system than the one reported in present study. However, the observed crystalline size in the present study is considerably finer than the one reported by Shim et al. [17] who synthesized around 100 nm sized TiB2 particle along with TiN, starting with Ti and BN powder particles through self-propagating high-temperature reaction as well as to that of Tanaka et al. [3] who reported 2–3 m sized TiB2 dispersoids in Fe–Cr–Ti ferrite matrix by conventional powder metallurgy route and 200–800 nm dispersed TiB2 synthesized through in situ reaction route. It is also noteworthy in the present study that nanocomposites containing TiB2 dispersoids can be obtained simply by high-energy ball milling without any in situ reaction and process control agents. 5. Conclusions
Fig. 7. Variation in particle size of the Fe–Cr–Al/30%TiB2 composite and crystalline size of TiB2 with milling time.
Nanocrystalline Fe–Cr–Al/30%TiB2 composite powder has been synthesized without process control agents from a micron sized mixture of metallic (Fe, Cr and Al) and ceramic (TiB2 ) powders by
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a high-energy ball milling process. Fine TiB2 dispersoids of size around 50 nm or lower are observed in the metallic matrix of Fe–Cr–Al solid solution. In addition to it, the crystalline size of TiB2 particles decreases with longer ball milling duration, resulting in an average size of 36 nm after 48 h of ball milling. The mean particle size of the composite powder, however, increases three times due to agglomeration of powders in the early stage of ball milling, and then reduces gradually with further milling. Acknowledgements This research was supported by an Institutional R&D Program of KIST. We wish to thank Dr. J.H. Shim and Dr. Y.W. Cho for their kind discussion and the supply of sample preparation facilities. We are also very thankful to Mr. S.I. Baik and Dr. Y.W. Kim at Seoul National University for transmission electron microscopic analysis work. References [1] R.G. Munro, J. Am. Ceram. Soc. 80 (1997) 1919–1928.
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