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Compositional dependence of amorphization of M–C–Si (M=Fe, Co or Ni) materials by mechanical alloying
Journal of Materials Processing Technology 143–144 (2003) 256–260
Compositional dependence of amorphization of M–C–Si (M = Fe, Co or Ni) materials by mechanical alloying Hidenori Ogawa∗ , Harumatsu Miura College of Industrial Technology, 1-27-1 Nishikoya Amagasaki, 661-0047 Japan
Abstract Mechanical alloying (MA) of elemental powder mixtures of Mx (C, Si)100−x (M = Fe, Co or Ni) with x = 65–85 at.% was carried out in an argon atmosphere by use of a planetary ball mill. In the MA processing, the addition of silicon to M–C materials highly enhanced amorphization of the MA materials. X-ray diffraction (XRD) and differential scanning calorimetry (DSC) results of MA-processed materials showed that after 720 ks of processing, Cox (C, Si)100−x samples can be amorphized in the compositional range of x = 65–80 at.%, whereas the amorphous range of Mx (C, Si)100−x (M = Fe or Ni) accessible under the same milling condition lies in the range of x = 70–75 at.%. This seems to be associated with the difference in crystal plasticity of the host metals M, in view of the fact that the metals Co, Fe, and Ni have hexagonal close-packed (h.c.p.), body-centered cubic (b.c.c.), and face-centered cubic (f.c.c.) structures at ambient temperature, respectively. Prolonged MA processing made it possible to amorphize Fe80 (C, Si)20 materials whose amorphization were not attained by liquid quenching process [Metall. Trans. A 18 (1987) 715]. The crystallization temperature Tx and Curie point Tc of MA-processed amorphous Fe70 (C, Si)30 samples, measured by DSC, raised with increasing Si content. For instance, the values of Tx and Tc of the Fe70 C10 Si20 sample were 761 and 624 K, respectively. © 2003 Published by Elsevier B.V. Keywords: M–C–Si (M = Fe, Co or Ni) materials; Mechanical alloying; Amorphization; Compositional dependence; Host metal crystal plasticity; Host metal concentration
1. Introduction Fe–C–Si amorphous alloys are particularly attractive as one of the soft magnetic materials, because their raw materials are almost inexhaustible and contain no expensive elements like boron. Inoue et al. [1] that showed Fe–C–Si amorphous alloys with relatively higher metalloid concentrations (30–35 at.%) can be prepared by a liquid quenching process. They reported, however, that because of such high metalloid concentrations, the amorphous alloys exhibit low magnetization and high coercive force [2] in comparison with other Fe-based amorphous alloys with lower metalloid concentrations such as amorphous Fe–C–B (10–25 at.%) and Fe–C–P (17–27 at.%) alloys [3]. On the other hand, since Koch et al. [4] made the first definitive demonstration that structurally amorphous alloys of Ni–Nb can be produced by mechanical alloying (MA) from crystalline elemental powder mixtures using a ball mill, amorphous alloy powders of various binary metal–metal systems such as Ni–Ti, Ni–Zr, and Cu–Zr have successfully been produced by the MA process as well [5,6]. However, ∗ Corresponding author. Tel.: +81-6-6431-7173; fax: +81-6-6431-6243. E-mail address: [email protected] (H. Ogawa).
0924-0136/$ – see front matter © 2003 Published by Elsevier B.V. doi:10.1016/S0924-0136(03)00306-6
for metal–metalloid systems, technologically most important group, only a few experimental data on amorphization by MA are available as briefly reviewed in our earlier paper [7]. As a matter of fact, it is said that amorphous alloys of metal–metalloid systems such as Fe–B and Fe–C are more difficult to prepare by the MA processing [8]. One reason for this is probably that the dissolution of such metalloid elements into iron phase does not readily occur, as conjectured from their extremely small solid solubilities indicated in the respective binary phase diagrams. In the present work, MA of M–C (M = Fe, Co or Ni) materials containing silicon (element with a much greater solid solubility than carbon in the host metal M) was performed with the aim of producing amorphous MA powder materials and of studying the composition-dependent amorphization in each system. The results obtained from the MA experiments will also be discussed from the crystal plasticity viewpoint of the host metals.
2. Experimental details Pure elemental powders with particle size in the range 63–149 m were used as starting materials for MA
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experiments. The powder mixtures of each desired average starting composition Mx (C, Si)100−x (M = Fe, Co or Ni) with x = 65–85 at.% were milled in an O-ring sealed cylindrical hardened tool steel vial (75 mm inside diameter×90 mm high) together with hardened steel balls in an argon atmosphere. The mass of the powder samples was 25 g and the ball-to-powder sample mass ratio was 11.27:1. The MA experiments were performed using a planetary ball mill which is the same as employed previously [9]. To avoid oxidation of the MA materials, the steel vials were not opened during the experiments. In the present work, the vial rotation speed was adjusted to 406 rpm. Amorphicity of milled powder samples was examined by XRD (Co K␣ radiation, λ = 0.179021 nm) and differential scanning calorimetry (DSC) (heating rate, 1.67 × 10−1 K/s).
3. Results and discussion
Fig. 2. XRD patterns for MA powders of the samples: (a) Fe75 C25 (at.%), (b) Fe75 C20 Si10 (at.%), (c) Fe75 C15 Si10 (at.%), (d) Fe75 C12.5 Si12.5 (at.%), (e) Fe75 C10 Si15 (at.%), and (f) Fe75 C5 Si20 (at.%) processed for 720 ks.
Figs. 1–4 show the X-ray diffraction (XRD) patterns of the Mx (C, Si)100−x MA powders with different metalloid concentrations (Fe80 (C, Si)20 , Fe75 (C, Si)25 , Fe70 (C, Si)30 , and Fe65 (C, Si)35 (at.%)) processed for 720 ks. The only Fe80 C10 Si10 (at.%) sample of these materials was processed for 1080 ks in order to examine the effect of processing time on solid-state reactions in the MA powders (Fig. 1(b)); in the 720 ks-processed Fe80 C10 Si10 MA sample, it was observed that a Bragg peak of ␣-Fe was superimposed on the broad diffraction pattern indicating the formation of an amorphous phase (not shown in Fig. 1). From these XRD results, it is obvious that solid-state reactions in MA powders depends greatly on the concentration of total metalloid components (CTM). Amorphization of the CTM = 20 and 35 samples does not fully proceed even after 720 ks of processing, in contrast to the CTM = 25 and 30 MA samples whose amorphization markedly prevails. Fig. 3. XRD patterns for MA powders of the samples: (a) Fe70 C25 Si5 (at.%), (b) Fe70 C20 Si10 (at.%), (c) Fe70 C15 Si15 (at.%), (d) Fe70 C10 Si20 (at.%), and (e) Fe70 C7.5 Si22.5 (at.%) processed for 720 ks.
Fig. 1. XRD patterns for: (a) Fe80 C15 Si5 (at.%), (b) Fe80 C10 Si10 (at.%), and (c) Fe80 C5 Si15 (at.%) MA powders. The MA samples (a) and (c) were processed for 720 ks, and the sample (b) was processed for 1080 ks. Co K␣ radiation (λ = 0.179021 nm).
Fig. 4. XRD patterns for MA powders of the samples: (a) Fe65 C25 Si10 (at.%), (b) Fe65 C15 Si20 (at.%), and (c) Fe65 C10 Si25 (at.%) processed for 720 ks.
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Fig. 2 illustrates the XRD patterns for MA powders of: (a) Fe75 C25 , (b) Fe75 C20 Si5 , (c) Fe75 C15 Si10 , (d) Fe75 C12.5 Si12.5 , (e) Fe75 C10 Si15 , and (f) Fe75 C5 Si20 (at.%) processed for 720 ks. In curve (a) of Fig. 2, the diffraction peak of ␣-Fe is largely broadened up to 720 ks of MA due to the refinement of iron grains and an increase in internal strain. Some peaks of Fe3 C are also visible in the diffraction pattern. So, the dissolution of carbon into the Fe phase of the sample does not seem to be so great also as conjectured from the work of Tanaka et al. [10] in which only a slight extension of solid solubility of carbon in ␣-Fe was formed in the ball-milled Fe–C materials. Curves (b) and (c) of Fig. 2 apparently shows that solid-state reactions in the Fe75 C20 Si5 and Fe75 C15 Si10 MA powders are strongly promoted due to processing for 720 ks so that each of the MA powder samples reveals a halo pattern typical of amorphous materials. Also, in the 720 ks-processed Fe70 (C, Si)30 MA samples, such effects of silicon more clearly appear as shown in Fig. 3 where a halo pattern characteristic of amorphous materials is seen in each of the XRD patterns for the MA samples of: (a) Fe70 C25 Si5 , (b) Fe70 C20 Si10 , (c) Fe70 C15 Si15 , and (d) Fe70 C10 Si10 . However, it is indicated that amorphization does not fully proceed in the MA powders represented by curves (d)–(f) of Fig. 2 and curve (e) of Fig. 3. The reason for this will be discussed later. When the thermal stability of SiC, i.e., the high melting point and large negative value of the free energy of formation [11] is taken into account under the present milling conditions, the MA processing is considered to generate a great attractive interaction between the silicon and carbon atoms in the MA materials having such high metalloid concentrations (25–30 at.%). Therefore, it seems that when the addition of silicon with a great solid solubility in iron is made to Fe–C materials, carbon atoms in the MA samples come to dissolve readily in the iron phase together with such silicon atoms during MA processing, i.e., intermixing of the atomic species iron, silicon, and carbon occurs on an atomic scale, leading to the formation of an amorphous phase. Similar MA experiments for Cox (C, Si)100−x and Nix (C, Si)100−x (x = 65–85 at.%) were carried out. Several XRD patterns are exhibited in Fig. 5. It is noted that amorphization reaction in Co–C–Si and Ni–C–Si specimens most readily takes place near the composition of x = 70–75 at.%, i.e., CTM = 30–25 as well, as in the case of Fe–C–Si materials. The M–C–Si MA specimens found to be amorphous by X-ray analysis were also examined by DSC. Each of the DSC thermograms showed a sharp exothermic reaction that can be attributed to the crystallization of amorphous phase formed in the MA powder sample. Some typical DSC curves for the M–C–Si MA samples are presented in Fig. 6. The XRD and DSC results for the 720 ks-processed M–C–Si MA samples, obtained in the present work, are summarized in Fig. 7. It is noteworthy that the widest composition range of amorphization is seen for the Co–C–Si MA sample whose host metal has the hexagonal close-packed
Fig. 5. XRD patterns for MA powders of the samples Co75 C20 Si5 , Co70 C20 Si10 , Co70 C10 Si20 , and Ni70 C10 Si20 (at.%) processed for 720 ks.
(h.c.p.) crystal structure. Thus it appears that the difference in the amorphous regions of the M–C–Si MA materials depicted in the Si (at.%) vs. C (at.%) diagrams of Fig. 7 is associated with the plasticity of the host metals M with the
Fig. 6. DSC curves for Co70 C10 Si20 , Fe70 C20 Si10 , and Fe75 C15 Si10 (at.%) MA powders processed for 720 ks, and Fe80 C10 Si10 (at.%) MA powders processed for 1080 ks. The letters Tx and Tc denote the crystallization temperature and Curie point, respectively. The inset illustrates a typical endothermic peak accompanying the magnetic transformation at a temperature of Tc . Heating rate: 1.67 × 10−1 K/s.
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Fig. 7. Compositional dependence of amorphization of M–C–Si (M = Fe, Co or Ni) MA materials processed for 720 ks.
different crystal structures, in view of the fact that the metals Fe, Co, and Ni have body-centered cubic (b.c.c.), h.c.p., and face-centered cubic (f.c.c.) structures at ambient temperature, respectively. In general, the rate of work-hardening of h.c.p. metals under plastic deformation is much greater at room temperature, due to the limited number of slip systems available, than that of f.c.c. and b.c.c. metals. MA process is said to be a repetitive impact process that results in deformation, work-hardening, fracturing, fragmentation, and cold welding of powder particles trapped between the colliding balls in the mill. Accordingly, it seems that MA materials with a high work-hardening rate like Co more rapidly cause crystalline grains of milled powder particles to fracture and refine, thereby also generating a large concentration of clean crystalline interfaces. So MA materials, themselves, with a higher rate of work-hardening are expected to more strongly promote solid-state reactions, during MA, such as chemical alloying and amorphization through generating such a great concentration of clean interface and extremely increasing energies stored in milled crystalline products. This appears to be reflected in the diagrams shown in Fig. 7. Furthermore, in Fig. 7 it is noted that any M–C–Si MA materials are easy to amorphize in the carbon-rich composition range as compared with the silicon-rich one. Amorphization of MA materials is regarded as attainable when recovery processes, which compete with amorphization in MA products, such as nucleation and growth of thermodynamically preferred crystalline intermetallic compounds are suppressed during MA processing. Table 1 presents the enthalpies of formation, H for , of the binary compounds M3 C
and M3 Si (M = Fe, Co or Ni) in the M–C and M–Si systems [12]. In view of the values of H for in Table 1 indicating that M3 C is thermodynamically less stable, suppression of recovery processes leading to formation of M3 C may be easy in comparison with that of M3 Si during MA processing. This seems to be reflected in the difference in amorphization in the carbon-rich and silicon-rich side regions. Although the silicide Fe3 Si [13] has a main diffraction peak, in the displayed 2θ range, at 2θ angle of 0.9331 rad (53.46◦ ) [14] which is located near that of ␣-Fe showing a main peak at 2θ angle of 0.9133 rad (52.33◦ ), the diffraction peak of Fe3 Si was not recognized in XRD analyses of MA powders in silicon-rich side regions. One reason for this is probably that even though the Fe3 Si phase is formed in the MA powders, the diffraction peak may be broadened to a great
Table 1 Enthalpy of formation, H for (kJ mol−1 ), of binary compounds M3 C and M3 Si (M = Fe, Co or Ni) in the M–C and M–Si systems Host metals, M
Fe Co Ni
Binary compounds M3 C
M3 Si
−1 +4 +7
−21 −24 −26
Fig. 8. Variation of crystallization temperature Tx and Curie point Tc with silicon concentration in amorphous Fe70 (C, Si)30 MA samples processed for 720 ks. Open circles denote the published data on amorphous liquid-quenched (LQ) ribbons with the same compositions as the present system.
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Amorphous Fe80 (C, Si)20 MA samples thus obtained are expected to exhibit better magnetic properties like higher magnetization characteristics.
4. Conclusions
Fig. 9. Effect of the prolonged MA processing an amorphization of Fe–C–Si MA materials. The circles (䊉), ( ), and ( ) denote amorphous (Am) compositions after MA processing for 720, 1080 and 1440 ks, respectively.
extent, due to grain refinement to nanosize levels after MA processing leading to an extremely severe deformation, so that the broadened diffraction peak of Fe3 Si is superimposed on that of ␣-Fe to show an apparent broadened single peak. However, more data are needed on this point. The crystallization temperature Tx and Curie point Tc of the amorphous Fe70 (C, Si)30 MA samples, respectively, were determined from the extrapolated onset temperatures (i.e., the value of Tx is given as the intersection of a line extrapolated through the straight line portion of the DSC curve towards its peak with the base line) and endothermic peaks in the DSC curves represented in Fig. 6. The results are given in Fig. 8 where the values of Tx and Tc are depicted as a function of silicon concentration in the MA Fe70 (C, Si)30 materials. The Tx and Tc values of the MA samples are comparable with the data, labeled by open circles in Fig. 8, available in the literature [1] on amorphous, rapidly quenched ribbons of the same compositions as the present systems. Though there appears a maximum value of Tx at a silicon concentration of 10 at.% in the graph, the reason is not clear. The work on this is still in progress. Fig. 9 illustrates the effect of the prolonged MA processing (1080–1440 ks) on amorphization of Fe–C–Si MA samples. With increasing processing time, the amorphized region extends toward lower total metalloid concentration range, i.e., higher iron concentration. The Fe80 (C, Si)20 MA samples with higher iron concentration are possible to be amorphized after 1080–1440 ks milling (Figs. 1 and 6), in contrast to the observation for the liquid quenching process.
(1) The addition of silicon to M–C materials highly enhances amorphization of the MA materials. (2) Amorphization in MA materials of Mx (C, Si)100−x (at.%) (M = Fe, Co, and Ni) depends greatly on the crystal structure of the host metal M. The widest composition range of amorphization was seen for the Co–C–Si MA system whose host metal has h.c.p. structure. (3) Amorphization reaction most readily occurs near the composition of x = 70–75 at.% in any MA samples. (4) Any MA materials of M–C–Si were easy to amorphize in C-rich composition range in comparison with the Si-rich one. (5) Prolonged MA processing made it possible to amorphize Fe80 (C, Si)20 materials whose amorphization was not attained by the liquid quenching process.
References [1] A. Inoue, S. Furukawa, T. Masumoto, Metall. Trans. A 18 (1987) 715. [2] H. Fujimori, in: T. Masumoto (Ed.), Materials Science of Amorphous Metals, Ohmu Pub. Co., Tokyo, 1982, p. 97. [3] M. Naka, T. Masumoto, Sci. Rep. Res. Inst. Tohoku Univ. 27 (1979) 118. [4] C.C. Koch, O.B. Calvin, C.G. McKamey, J.O. Scarbrough, Appl. Phys. Lett. 43 (1983) 1017. [5] L. Schults, E. Hellstern, in: M. Tenhover, L.E. Tanner, W.L. Johnson (Eds.), Science and Technology of Rapidly Quenched Alloys, Materials Research Society Symposium Proceedings, vol. 80, Materials Research Society, Pittsburgh, PA, 1987, p. 3. [6] C. Politis, W.L. Johnson, J. Appl. Phys. 60 (1986) 1147. [7] H. Miura, S. Isa, K. Omuro, Jpn. J. Appl. Phys. 29 (1990) L339. [8] R.B. Schwarz, in: J.D. Embury, G.R. Purdy (Eds.), Advances in Phase Transitions, Pergamon Press, Oxford, 1988, p. 166. [9] K. Omuro, H. Miura, Appl. Phys. Lett. 60 (1992) 1433. [10] T. Tanaka, K.N. Ishihara, P.H. Shingu, Metall. Trans. A 23 (1992) 2431. [11] C.G. Harman, W.G. Mixer, US Atomic Energy Communications, Publication No. BMI-748, 1952. [12] F.R. de Boer, R. Boom, W.C.M. Mattens, A.R. Miedema, A.K. Niessen, Cohesion in Metals, North-Holland, Amsterdam, 1988, p. 219. [13] M. Hansen, Constitution of Binary Alloys, McGraw-Hill, New York, 1958, p. 711. [14] Powder Diffraction File, Alphabetical Index, Inorganic Phases, International Center for Diffraction Data, Swarthmore, PA, 1985, p. 704.
Compositional dependence of amorphization of M–C–Si (M = Fe, Co or Ni) materials by mechanical alloying Hidenori Ogawa∗ , Harumatsu Miura College of Industrial Technology, 1-27-1 Nishikoya Amagasaki, 661-0047 Japan
Abstract Mechanical alloying (MA) of elemental powder mixtures of Mx (C, Si)100−x (M = Fe, Co or Ni) with x = 65–85 at.% was carried out in an argon atmosphere by use of a planetary ball mill. In the MA processing, the addition of silicon to M–C materials highly enhanced amorphization of the MA materials. X-ray diffraction (XRD) and differential scanning calorimetry (DSC) results of MA-processed materials showed that after 720 ks of processing, Cox (C, Si)100−x samples can be amorphized in the compositional range of x = 65–80 at.%, whereas the amorphous range of Mx (C, Si)100−x (M = Fe or Ni) accessible under the same milling condition lies in the range of x = 70–75 at.%. This seems to be associated with the difference in crystal plasticity of the host metals M, in view of the fact that the metals Co, Fe, and Ni have hexagonal close-packed (h.c.p.), body-centered cubic (b.c.c.), and face-centered cubic (f.c.c.) structures at ambient temperature, respectively. Prolonged MA processing made it possible to amorphize Fe80 (C, Si)20 materials whose amorphization were not attained by liquid quenching process [Metall. Trans. A 18 (1987) 715]. The crystallization temperature Tx and Curie point Tc of MA-processed amorphous Fe70 (C, Si)30 samples, measured by DSC, raised with increasing Si content. For instance, the values of Tx and Tc of the Fe70 C10 Si20 sample were 761 and 624 K, respectively. © 2003 Published by Elsevier B.V. Keywords: M–C–Si (M = Fe, Co or Ni) materials; Mechanical alloying; Amorphization; Compositional dependence; Host metal crystal plasticity; Host metal concentration
1. Introduction Fe–C–Si amorphous alloys are particularly attractive as one of the soft magnetic materials, because their raw materials are almost inexhaustible and contain no expensive elements like boron. Inoue et al. [1] that showed Fe–C–Si amorphous alloys with relatively higher metalloid concentrations (30–35 at.%) can be prepared by a liquid quenching process. They reported, however, that because of such high metalloid concentrations, the amorphous alloys exhibit low magnetization and high coercive force [2] in comparison with other Fe-based amorphous alloys with lower metalloid concentrations such as amorphous Fe–C–B (10–25 at.%) and Fe–C–P (17–27 at.%) alloys [3]. On the other hand, since Koch et al. [4] made the first definitive demonstration that structurally amorphous alloys of Ni–Nb can be produced by mechanical alloying (MA) from crystalline elemental powder mixtures using a ball mill, amorphous alloy powders of various binary metal–metal systems such as Ni–Ti, Ni–Zr, and Cu–Zr have successfully been produced by the MA process as well [5,6]. However, ∗ Corresponding author. Tel.: +81-6-6431-7173; fax: +81-6-6431-6243. E-mail address: [email protected] (H. Ogawa).
0924-0136/$ – see front matter © 2003 Published by Elsevier B.V. doi:10.1016/S0924-0136(03)00306-6
for metal–metalloid systems, technologically most important group, only a few experimental data on amorphization by MA are available as briefly reviewed in our earlier paper [7]. As a matter of fact, it is said that amorphous alloys of metal–metalloid systems such as Fe–B and Fe–C are more difficult to prepare by the MA processing [8]. One reason for this is probably that the dissolution of such metalloid elements into iron phase does not readily occur, as conjectured from their extremely small solid solubilities indicated in the respective binary phase diagrams. In the present work, MA of M–C (M = Fe, Co or Ni) materials containing silicon (element with a much greater solid solubility than carbon in the host metal M) was performed with the aim of producing amorphous MA powder materials and of studying the composition-dependent amorphization in each system. The results obtained from the MA experiments will also be discussed from the crystal plasticity viewpoint of the host metals.
2. Experimental details Pure elemental powders with particle size in the range 63–149 m were used as starting materials for MA
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experiments. The powder mixtures of each desired average starting composition Mx (C, Si)100−x (M = Fe, Co or Ni) with x = 65–85 at.% were milled in an O-ring sealed cylindrical hardened tool steel vial (75 mm inside diameter×90 mm high) together with hardened steel balls in an argon atmosphere. The mass of the powder samples was 25 g and the ball-to-powder sample mass ratio was 11.27:1. The MA experiments were performed using a planetary ball mill which is the same as employed previously [9]. To avoid oxidation of the MA materials, the steel vials were not opened during the experiments. In the present work, the vial rotation speed was adjusted to 406 rpm. Amorphicity of milled powder samples was examined by XRD (Co K␣ radiation, λ = 0.179021 nm) and differential scanning calorimetry (DSC) (heating rate, 1.67 × 10−1 K/s).
3. Results and discussion
Fig. 2. XRD patterns for MA powders of the samples: (a) Fe75 C25 (at.%), (b) Fe75 C20 Si10 (at.%), (c) Fe75 C15 Si10 (at.%), (d) Fe75 C12.5 Si12.5 (at.%), (e) Fe75 C10 Si15 (at.%), and (f) Fe75 C5 Si20 (at.%) processed for 720 ks.
Figs. 1–4 show the X-ray diffraction (XRD) patterns of the Mx (C, Si)100−x MA powders with different metalloid concentrations (Fe80 (C, Si)20 , Fe75 (C, Si)25 , Fe70 (C, Si)30 , and Fe65 (C, Si)35 (at.%)) processed for 720 ks. The only Fe80 C10 Si10 (at.%) sample of these materials was processed for 1080 ks in order to examine the effect of processing time on solid-state reactions in the MA powders (Fig. 1(b)); in the 720 ks-processed Fe80 C10 Si10 MA sample, it was observed that a Bragg peak of ␣-Fe was superimposed on the broad diffraction pattern indicating the formation of an amorphous phase (not shown in Fig. 1). From these XRD results, it is obvious that solid-state reactions in MA powders depends greatly on the concentration of total metalloid components (CTM). Amorphization of the CTM = 20 and 35 samples does not fully proceed even after 720 ks of processing, in contrast to the CTM = 25 and 30 MA samples whose amorphization markedly prevails. Fig. 3. XRD patterns for MA powders of the samples: (a) Fe70 C25 Si5 (at.%), (b) Fe70 C20 Si10 (at.%), (c) Fe70 C15 Si15 (at.%), (d) Fe70 C10 Si20 (at.%), and (e) Fe70 C7.5 Si22.5 (at.%) processed for 720 ks.
Fig. 1. XRD patterns for: (a) Fe80 C15 Si5 (at.%), (b) Fe80 C10 Si10 (at.%), and (c) Fe80 C5 Si15 (at.%) MA powders. The MA samples (a) and (c) were processed for 720 ks, and the sample (b) was processed for 1080 ks. Co K␣ radiation (λ = 0.179021 nm).
Fig. 4. XRD patterns for MA powders of the samples: (a) Fe65 C25 Si10 (at.%), (b) Fe65 C15 Si20 (at.%), and (c) Fe65 C10 Si25 (at.%) processed for 720 ks.
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Fig. 2 illustrates the XRD patterns for MA powders of: (a) Fe75 C25 , (b) Fe75 C20 Si5 , (c) Fe75 C15 Si10 , (d) Fe75 C12.5 Si12.5 , (e) Fe75 C10 Si15 , and (f) Fe75 C5 Si20 (at.%) processed for 720 ks. In curve (a) of Fig. 2, the diffraction peak of ␣-Fe is largely broadened up to 720 ks of MA due to the refinement of iron grains and an increase in internal strain. Some peaks of Fe3 C are also visible in the diffraction pattern. So, the dissolution of carbon into the Fe phase of the sample does not seem to be so great also as conjectured from the work of Tanaka et al. [10] in which only a slight extension of solid solubility of carbon in ␣-Fe was formed in the ball-milled Fe–C materials. Curves (b) and (c) of Fig. 2 apparently shows that solid-state reactions in the Fe75 C20 Si5 and Fe75 C15 Si10 MA powders are strongly promoted due to processing for 720 ks so that each of the MA powder samples reveals a halo pattern typical of amorphous materials. Also, in the 720 ks-processed Fe70 (C, Si)30 MA samples, such effects of silicon more clearly appear as shown in Fig. 3 where a halo pattern characteristic of amorphous materials is seen in each of the XRD patterns for the MA samples of: (a) Fe70 C25 Si5 , (b) Fe70 C20 Si10 , (c) Fe70 C15 Si15 , and (d) Fe70 C10 Si10 . However, it is indicated that amorphization does not fully proceed in the MA powders represented by curves (d)–(f) of Fig. 2 and curve (e) of Fig. 3. The reason for this will be discussed later. When the thermal stability of SiC, i.e., the high melting point and large negative value of the free energy of formation [11] is taken into account under the present milling conditions, the MA processing is considered to generate a great attractive interaction between the silicon and carbon atoms in the MA materials having such high metalloid concentrations (25–30 at.%). Therefore, it seems that when the addition of silicon with a great solid solubility in iron is made to Fe–C materials, carbon atoms in the MA samples come to dissolve readily in the iron phase together with such silicon atoms during MA processing, i.e., intermixing of the atomic species iron, silicon, and carbon occurs on an atomic scale, leading to the formation of an amorphous phase. Similar MA experiments for Cox (C, Si)100−x and Nix (C, Si)100−x (x = 65–85 at.%) were carried out. Several XRD patterns are exhibited in Fig. 5. It is noted that amorphization reaction in Co–C–Si and Ni–C–Si specimens most readily takes place near the composition of x = 70–75 at.%, i.e., CTM = 30–25 as well, as in the case of Fe–C–Si materials. The M–C–Si MA specimens found to be amorphous by X-ray analysis were also examined by DSC. Each of the DSC thermograms showed a sharp exothermic reaction that can be attributed to the crystallization of amorphous phase formed in the MA powder sample. Some typical DSC curves for the M–C–Si MA samples are presented in Fig. 6. The XRD and DSC results for the 720 ks-processed M–C–Si MA samples, obtained in the present work, are summarized in Fig. 7. It is noteworthy that the widest composition range of amorphization is seen for the Co–C–Si MA sample whose host metal has the hexagonal close-packed
Fig. 5. XRD patterns for MA powders of the samples Co75 C20 Si5 , Co70 C20 Si10 , Co70 C10 Si20 , and Ni70 C10 Si20 (at.%) processed for 720 ks.
(h.c.p.) crystal structure. Thus it appears that the difference in the amorphous regions of the M–C–Si MA materials depicted in the Si (at.%) vs. C (at.%) diagrams of Fig. 7 is associated with the plasticity of the host metals M with the
Fig. 6. DSC curves for Co70 C10 Si20 , Fe70 C20 Si10 , and Fe75 C15 Si10 (at.%) MA powders processed for 720 ks, and Fe80 C10 Si10 (at.%) MA powders processed for 1080 ks. The letters Tx and Tc denote the crystallization temperature and Curie point, respectively. The inset illustrates a typical endothermic peak accompanying the magnetic transformation at a temperature of Tc . Heating rate: 1.67 × 10−1 K/s.
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Fig. 7. Compositional dependence of amorphization of M–C–Si (M = Fe, Co or Ni) MA materials processed for 720 ks.
different crystal structures, in view of the fact that the metals Fe, Co, and Ni have body-centered cubic (b.c.c.), h.c.p., and face-centered cubic (f.c.c.) structures at ambient temperature, respectively. In general, the rate of work-hardening of h.c.p. metals under plastic deformation is much greater at room temperature, due to the limited number of slip systems available, than that of f.c.c. and b.c.c. metals. MA process is said to be a repetitive impact process that results in deformation, work-hardening, fracturing, fragmentation, and cold welding of powder particles trapped between the colliding balls in the mill. Accordingly, it seems that MA materials with a high work-hardening rate like Co more rapidly cause crystalline grains of milled powder particles to fracture and refine, thereby also generating a large concentration of clean crystalline interfaces. So MA materials, themselves, with a higher rate of work-hardening are expected to more strongly promote solid-state reactions, during MA, such as chemical alloying and amorphization through generating such a great concentration of clean interface and extremely increasing energies stored in milled crystalline products. This appears to be reflected in the diagrams shown in Fig. 7. Furthermore, in Fig. 7 it is noted that any M–C–Si MA materials are easy to amorphize in the carbon-rich composition range as compared with the silicon-rich one. Amorphization of MA materials is regarded as attainable when recovery processes, which compete with amorphization in MA products, such as nucleation and growth of thermodynamically preferred crystalline intermetallic compounds are suppressed during MA processing. Table 1 presents the enthalpies of formation, H for , of the binary compounds M3 C
and M3 Si (M = Fe, Co or Ni) in the M–C and M–Si systems [12]. In view of the values of H for in Table 1 indicating that M3 C is thermodynamically less stable, suppression of recovery processes leading to formation of M3 C may be easy in comparison with that of M3 Si during MA processing. This seems to be reflected in the difference in amorphization in the carbon-rich and silicon-rich side regions. Although the silicide Fe3 Si [13] has a main diffraction peak, in the displayed 2θ range, at 2θ angle of 0.9331 rad (53.46◦ ) [14] which is located near that of ␣-Fe showing a main peak at 2θ angle of 0.9133 rad (52.33◦ ), the diffraction peak of Fe3 Si was not recognized in XRD analyses of MA powders in silicon-rich side regions. One reason for this is probably that even though the Fe3 Si phase is formed in the MA powders, the diffraction peak may be broadened to a great
Table 1 Enthalpy of formation, H for (kJ mol−1 ), of binary compounds M3 C and M3 Si (M = Fe, Co or Ni) in the M–C and M–Si systems Host metals, M
Fe Co Ni
Binary compounds M3 C
M3 Si
−1 +4 +7
−21 −24 −26
Fig. 8. Variation of crystallization temperature Tx and Curie point Tc with silicon concentration in amorphous Fe70 (C, Si)30 MA samples processed for 720 ks. Open circles denote the published data on amorphous liquid-quenched (LQ) ribbons with the same compositions as the present system.
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Amorphous Fe80 (C, Si)20 MA samples thus obtained are expected to exhibit better magnetic properties like higher magnetization characteristics.
4. Conclusions
Fig. 9. Effect of the prolonged MA processing an amorphization of Fe–C–Si MA materials. The circles (䊉), ( ), and ( ) denote amorphous (Am) compositions after MA processing for 720, 1080 and 1440 ks, respectively.
extent, due to grain refinement to nanosize levels after MA processing leading to an extremely severe deformation, so that the broadened diffraction peak of Fe3 Si is superimposed on that of ␣-Fe to show an apparent broadened single peak. However, more data are needed on this point. The crystallization temperature Tx and Curie point Tc of the amorphous Fe70 (C, Si)30 MA samples, respectively, were determined from the extrapolated onset temperatures (i.e., the value of Tx is given as the intersection of a line extrapolated through the straight line portion of the DSC curve towards its peak with the base line) and endothermic peaks in the DSC curves represented in Fig. 6. The results are given in Fig. 8 where the values of Tx and Tc are depicted as a function of silicon concentration in the MA Fe70 (C, Si)30 materials. The Tx and Tc values of the MA samples are comparable with the data, labeled by open circles in Fig. 8, available in the literature [1] on amorphous, rapidly quenched ribbons of the same compositions as the present systems. Though there appears a maximum value of Tx at a silicon concentration of 10 at.% in the graph, the reason is not clear. The work on this is still in progress. Fig. 9 illustrates the effect of the prolonged MA processing (1080–1440 ks) on amorphization of Fe–C–Si MA samples. With increasing processing time, the amorphized region extends toward lower total metalloid concentration range, i.e., higher iron concentration. The Fe80 (C, Si)20 MA samples with higher iron concentration are possible to be amorphized after 1080–1440 ks milling (Figs. 1 and 6), in contrast to the observation for the liquid quenching process.
(1) The addition of silicon to M–C materials highly enhances amorphization of the MA materials. (2) Amorphization in MA materials of Mx (C, Si)100−x (at.%) (M = Fe, Co, and Ni) depends greatly on the crystal structure of the host metal M. The widest composition range of amorphization was seen for the Co–C–Si MA system whose host metal has h.c.p. structure. (3) Amorphization reaction most readily occurs near the composition of x = 70–75 at.% in any MA samples. (4) Any MA materials of M–C–Si were easy to amorphize in C-rich composition range in comparison with the Si-rich one. (5) Prolonged MA processing made it possible to amorphize Fe80 (C, Si)20 materials whose amorphization was not attained by the liquid quenching process.
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