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Mechanical alloying of Fe–Mn and Fe–Mn–Si
Materials Science and Engineering A271 (1999) 8 – 13 www.elsevier.com/locate/msea
Mechanical alloying of Fe–Mn and Fe–Mn–Si T. Liu b, H.Y. Liu a,*, Z.T. Zhao c, R.Z. Ma c, T.D. Hu b, Y.N. Xie b a
Laboratory of Semiconductor Materials Science, Institute of Semiconductors, Chinese Academy of Science, Beijing 100083, People’s Republic of China b Beijing Synchrotron Radiation Facility, Institute of High Energy Physics, Chinese Academy of Science, Beijing 100083, People’s Republic of China c Department of Materials Physics, Uni6ersity of Science and Technology, Chinese Academy of Science, Beijing 100083, People’s Republic of China Received 7 May 1998; received in revised form 22 September 1998
Abstract The ball milling of Fe–24Mn and Fe–24Mn–6Si mixed powders has been performed by the high energy ball milling technique. By employing X-ray diffraction and Mo¨ssbauer measurements, the composition evolution during the milling process has been investigated. The results indicate the formation of paramagnetic Fe – Mn or Fe – Mn – Si alloys with a metastable fcc phase as final products, which imply that the Fe and Mn proceed a co-diffusion mechanism through the surface of fragmented powders. The thermal stability and composition evolution of the as-milled alloys were discussed comparing with the bulk alloy. © 1999 Published by Elsevier Science S.A. All rights reserved. Keywords: As-milled alloys; Ball milling; Mo¨ssbauer measurements; X-ray diffraction
1. Introduction Solid state amorphization reaction (SSAR) induced by mechanical alloying (MA) or mechanical milling (MM) has recently been an intensive research subject because of the great potential for producing amorphous materials or nanostructured materials for technological applications. However, the mechanism related to the process of MM or MA is still ambiguous [1,2]. In the past years mechanical alloying of binary metallic elements, or metal– metalloid systems have been widely investigated [3–6]. To our knowledge there have been no previous studies of mechanically alloyed Fe–Mn and Fe–Mn–Si reported. The Fe – Mn based alloys are of great importance for stainless steel, non-magnetism steel and which are also found wide application as important parts in Fe-based shape memory alloys. In the present study powders of Fe – 24Mn and Fe– 24Mn–6Si mixtures were ball-milled, the alloying concerned mechanism of the samples along with the milling process was investigated through X-ray diffraction pat* Corresponding author. E-mail address: [email protected] (H.Y. Liu)
terns and Mo¨ssbauer spectra. Because of its localized character, Mo¨ssbauer can give information about local environment of an Fe atom. Both Fe and Mn are transition metals and are located closely in the periodic table and the two elements have small mixing enthalpy, DHmix, of only +1 kJ mol − 1 [7]. Therefore, the chemical driving force for Fe–Mn alloying can be neglected. Also the atomic sizes of both elements in their bcc forms are very similar. Thus, one would expect that Fe–Mn is a non-fast diffusion system. The bcc, hcp and fcc phases coexist in the Fe–Mn binary alloy, depending on the Mn concentration, and can transform into each other by martensitic transformation.
2. Experimental details The Fe–Mn and Fe–Mn–Si alloys with normal compositions of Fe–24Mn and Fe–24Mn–6Si (wt.%), respectively, were prepared by appropriate masses of pure metal powders of iron (99.5%), maganese (99.7%) and silicon (99.9%). Ball milling of the powder mixtures was performed with a vibration mill using hardened
0921-5093/99/$ - see front matter © 1999 Published by Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 0 2 2 - 3
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
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steel vials and balls. The ball-to-powder weight ratio was :20:1. The vial was loaded and sealed in an argon atmosphere. Annealing of the as-milled Fe – 24Mn–6Si was carried out in a tube furnace at 500, 650 and 700°C for 1 h, under argon atmosphere protection. X-ray diffraction measurements on the as-milled powders were carried out in a Rigaku D/max-RB X-ray diffractometer using the Cu – Ka radiation. Transmission 57Fe Mo¨ssbauer spectra were collected with a conventional constant-acceleration spectrometer using a 57 Co(Pd) source at room temperature. Velocity calibration was carried out using an a-Fe foil. The Mo¨ssbauer spectra were fitted by a program developed by Xu et al. [8]. Differential thermal analysis (DTA) was performed in the DTA-1700 (Perkin Elmer) at a heating rate of 10 K min − 1.
3. Results and discussion
3.1. Mechanical alloying of Fe– Mn Fig. 1 shows X-ray diffraction patterns of the as-filed and ball-milled Fe – Mn powders for 8, 27, and 67 h, respectively. Upon milling the diffraction intensities corresponding to the a-Fe and a-Mn tends to decrease and diminish gradually. In the same time a new phase appears with slightly different lattice parameters from those of a-Fe and a-Mn appears, which are identified as fcc Fe–Mn alloy. Apparently the ball-milling has induced the alloying of Fe and Mn. The diffraction peaks of Fe–Mn alloy are broadened due to the refinement of Fe–Mn particles. However, no trace of amorphization is noted even after 67 h of milling, this may be ascribed to the non-fast diffusion system that Fe and
Fig. 1. XRD patterns of Fe–24Mn, milled for 8, 24 and 67 h, respectively. The 67 h milled pattern indicates the appearance of fcc Fe– Mn.
Fig. 2. Mo¨ssbauer spectra of Fe – 24Mn, milled for 8, 24, 46 and 67 h, respectively. In the 8 h spectrum the sextet subspectra with a small hyperfine field and two singlet subspectra, employed to fit the overall spectra is shown.
Mn belong, that is, the condition for amorphization of binary alloys could not be satisfied [9]. In order to further investigate the local variation associated mechanism with the milling, the 57Fe Mo¨ssbauer spectra of the milled samples were obtained as shown in Fig. 2. A portion of Fe powders has fragmented into fine particles with paramagnetic behaviors upon 8 h milling, as demonstrated by the coexistence of ferromagnetic sextet and paramagnetic singlet peaks. At 27 h the fraction of the ferromagnetic phase becomes more unimportant and tends to diminish after 46 h of milling. The spectrum of 67 h corresponds only to the singlet line of the paramagnetic Fe–Mn phase. Even though the mixing enthalpy of Fe–Mn is not large, mechanical alloying supplies sufficient driving force for Fe–Mn alloying. Upon milling, the powders were fragmented into fine particles, as a result, Fe and Mn diffuse into each other, forming an fcc phase of Fe–Mn, and induces the material change from ferromagnetic to paramagnetic. Considering the character of the central singlet line, we employed two singlet subspectra to fit the paramagnetic phase, as indicated in Fig. 2 for the 8 h spectrum, the two subspectra can be regarded to stand for different local environments around Fe atoms. The fitted results of Mo¨ssbauer spectra for the Fe– Mn alloys milled for different time are listed in Table 1. The value of isomer shift (IS) for the paramagnetic phase is averaged over the two subspectra. As for the
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
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Table 1 Results of fitting Mo¨ssbauer spectra of Fe–Mn samples milled for 8, 27, 46 and 67 ha Time (h)
8 27 46 67 a
Sextet 1
Sextet 2
Singlet
IS
QS
Bhf
A
IS
QS
Bhf
A
IS
A
0.01 0.13 – –
0.01 −0.01 – –
32.8 33.0 – –
61.6 14.1 – –
0.13 0.23 – –
−0.12 −0.22 – –
23.7 28.3 – –
13.9 10.1 – –
−0.09 −0.07 −0.07 −0.07
24.6 75.8 100 100
IS: mm s−1, QS: mm s−1, Bhf: T, A: %, A: area fraction of subspectra.
ferromagnetic phase, besides the sextet subspectrum of a-Fe, another subspectrum with Voigt line can be resolved and fitted with a hyperfine field of 23.7 T according to the sink of the sextet peak. Upon milling some positions originally occupied by Fe atoms, especially for those Fe atoms at the surface of particles, may be taken place now by Mn atoms due to the diffusion, and thus leads to the decrease of hyperfine field, as seen in Table 1. Besides this, the IS and quadrupole splitting (QS) are quite different from those of the a-Fe. On milling one may expect that the Fe atoms on the surface may suffer from intense bump and the lattice distortion of Fe lattice due to the mismatch of lattice constants between Fe and Mn atoms, and thus leading to the deviation of 3d electron between Fe and Mn. In the same time Fe atoms diffuse into Mn particles. The Fe atoms surrounded by Mn may transform into paramagnetic Fe and make contribution to the fraction of paramagnetic phase. As the process goes on, i.e. attachment of Fe and Mn particles surface diffusion bulk diffusion, the fraction of paramagnetic Fe–Mn phase increases. On the other hand, the phase compositions undergo a great change such the bcc Fe transforms into fcc g-Fe(Mn) solid solution. Meanwhile the Fe atoms dissolve soluted in the a-Mn lattice and form the gMn(Fe). According to the TEM of Fe40Cu60 observed by Wu et al. [10], the a-Fe(Cu) was found to transform into g-Fe(Cu) upon milling. The authors explained this change as due to the martensitic transformation. However, the current results indicate that the phase transformation of Fe – Mn is most probably diffusion-controlled, though, the Fe – Mn is very likely to be a non-fast diffusion system.
found that its value changes little upon milling for both Fe–24Mn and Fe–24Mn–6Si. The IS for Fe–Mn–Si sample increases from −0.07 to − 0.01 mm s − 1 relative to that of Fe–Mn due to the addition of Si. Recently we have performed the Mo¨ssbauer measurements for bulk Fe–30Mn–6Si alloys prepared by melting and quenched from high temperature (900°C) [11]. The XRD patterns indicate the quenched strip-shaped sample consists of only fcc austenitic phase. By fitting the singlet line of Mo¨ssbauer spectra the IS value is given as − 0.01 mm s − 1, which is very close to that of
3.2. Mechanical alloying of Fe– Mn – Si The XRD patterns and Mo¨ssbauer spectra of Fe– 24Mn–6Si samples milled for 8, 27, 46 and 67 h are shown in Figs. 3 and 4, respectively. The results of fitting the Mo¨ssbauer spectra are listed in Table 2. Similar to that of Fe – 24Mn, the final product consists of the paramagnetic Fe – 24Mn – 6Si alloy with a metastable fcc phase. Considering the fitted IS, we
Fig. 3. XRD patterns of Fe – 24Mn – 6Si, milled for 0, 24, 46 and 67 h, respectively. The 67 h milled pattern indicates the appearance of fcc Fe – Mn.
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
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6Si on the atomic level has not induced the amorphization of alloys as shown in the XRD patterns by the characteristic broadening of diffraction peaks. As we know, the g phase in Fe–Mn–Si shape memory alloy can only be stabilized at temperatures higher than about 900°C. At room temperature o phase precipitates due to g o martensitic transformation. However, the ball milling of Fe–Mn–Si reveals that ultrafine particles of g phase are relatively stable, which is quite different from the bulk alloy. One possible explanation is that the interfacial energy and strain energy have changed by mechanical alloying due to the formation of numerous structural defects. As a result, the fcc phase can be stabilized from the thermodynamic point of view. Similar explanation can also be adopted for Fe– 24Mn. According to the phase diagram, the stable Fe–Mn is composed of a and g phases at room temperature. However the XRD patterns indicate only the existence of g phase. In fact even cooling the milled Fe–Mn–Si samples down to a extremely low temperature of 1.5 K, no trace of o phase was found in the XRD patterns. We have investigated the reverse martensitic transformation of Fe–30Mn–6Si alloy [13] in a previous investigation. It is known that strain play a role as inductor of g o matensitic transformation, but our recent results indicated that intense bending would not induce the formation of a plenty of martensite; on the contrary, the g o martensitic transformation was suppressed. This is also true for ball-milling of Fe–Mn and Fe–Mn–Si, implying the strain-induced martensite is closely correlated with process. Similar results have been obtained by Hideshi et al. [14] when studied the martensitic transformation in Fe–Mn alloys. The critical point for martensite transforming into autensite was found to decrease while reducing the grain size of Fe–Mn alloys from 100 to 10 mm, the g o martensitic transformation was suppressed.
Fig. 4. Mo¨ssbauer spectra of Fe–24Mn–6Si, milled for 0, 24, 46 and 67 h, respectively. In the 8 h spectra the sextet subspectra with a small hyperfine field and two singlet subspectra, employed to fit the overall spectra are shown.
the ball-milled Fe – 24Mn – 6Si under investigation, despite the difference in Mn concentration. It seems that Mn has a less effect on the electronic arrangement of Fe atoms. With the addition of Si, the Si may take the place of Fe and forms a substitutional solid solution. As a result, the 3d electrons of Fe deviate from Si atoms and increase the IS value. A similar analysis was used to the Fe–Si alloys system studied by Stearn [12]. The results indicate that ball-milling has induced an intermixing of Fe, Mn and Si on an atomic level, and the Fe atoms seem to take the lattice position of Mn but not segregate to the grain boundary. It should be noted that after 46 h of ball-milling the system has reached a metastable equilibrium with a fcc phase. The intermixing of Fe – 24Mn and Fe –24Mn–
3.3. Thermal stability of as-milled Fe–Mn–Si The thermal effect associated with structural relaxation and grain growth of the ball-milled Fe–24Mn– 6Si samples was investigated by the DTA measured for
Table 2 Results of fitting Mo¨ssbauer spectra of Fe–Mn–Si samples milled for 8, 27, 46 and 67 ha Time (h)
8 27 46 67 a
Sextet 1
Sextet 2
Singlet
IS
QS
Bhf
A
IS
QS
Bhf
A
IS
A
0.01 0.00 −0.02 –
0.01 0.01 −0.01 –
33.2 33.1 33.0 –
47.4 14.6 2.8 –
0.05 −0.13 – –
0.08 0.24 – –
24.9 26.6 – –
38.4 8.4 – –
0.07 −0.01 −0.01 −0.01
14.2 77.0 97.2 100
IS: mm s−1, QS: mm s−1, Bhf: T, A: %, A: area fraction of subspectra.
12
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
Fig. 5. DTA thermogram of Fe–24Mn–6Si powder milled for 67 h (heat rate 40 K min − 1).
the powders milled for the 67 h sample at a heating rate of 40 K min − 1, as shown in Fig. 5. Two broad exothermal peaks are observed in the temperature range 510– 620°C and 670–730°C, respectively. A heat treatment was carried out for the as-milled sample at 500, 650 and 700°C, respectively, for 1 h. Fig. 6 shows the XRD patterns of the annealed samples. It shows that additional peaks emerge in the patterns of the fcc phase upon annealing, which are identified as a martensitic phase. The sharp peaks marked by X are hard to identify, which can probably be attributed to the contamination by the substrate. At 650°C the a phase has become dominant. However, for temperatures higher than 700°C, the fcc phase again plays a dominant role. Besides, the grain size of Fe – Mn powders shows a trace of growth-up with increasing temperatures when one take account of the broadening of diffraction peaks. According to the Fe – Mn phase diagram, the a phase or the g phase can be stabilized at certain temperature and alloy composition. Therefore, the first exothermal peak in the DTA trace can be ascribed to the ga transformation, and the second to the a g transformation, which is in agreement with the phase diagram of bulk Fe – Mn alloy.
Fig. 6. Room temperature XRD patterns of annealed Fe–24Mn–6Si powders at 500, 650 and 700°C, respectively, for 1 h.
Fe–Mn–Si. The metastable g phase of Fe–Mn or Fe–Mn–Si was found to be stable final product of ball-milling processes. No trace of o phase was detected as did in the bulk Fe–Mn–Si alloys. However, the as-milled Fe–Mn–Si powders show a similar thermal stability as in the phase diagram of bulk alloys.
References 4. Conclusions The Fe–Mn and Fe – Mn – Si alloys have been prepared by mechanical alloying of Fe – 24Mn and Fe– 24Mn–6Si at ambient temperature. The XRD patterns and Mo¨ssbauer spectra reveal that Fe and Mn atoms diffuse into each other through surface into bulk of ultrafine particles, and leading to the structural transformation and magnetic change upon milling of the Fe–Mn and Fe– Mn – Si powder mixtures. The phase transformation of the as-milled alloys shows a quite difference compared to the bulk alloy of Fe –Mn and
[1] A.W. Weeber, H. Bakker, Physica B 153 (1988) 93. [2] H. Miura, S. Isa, K. Omuro, J. Non-Cryst. Solids 117 (1990) 741. [3] A.R. Yavari, P.J. Desre, T. Benameur, Phys. Rev. Lett. 68 (1992) 2235. [4] J.G. Cabanas Moreno, V.M. Lopez, H.A. Calderon, J.C. Rendon-Angeles, Scripta Metall. Mater. 28 (1993) 645. [5] D.L. Zhang, J. Mater. Sci. Lett. 14 (1995) 1508. [6] O. Drbohlav, A.R. Yavari, Acta Metall. Mater. 43 (1995) 1799. [7] A.R. Miedema, P.F. de Chatel, F.R. de Boer, Physica B 100 (1980) 1. [8] Z.X. Xu, Z.D. Xu, H.S. Yang, R.Z. Ma, Chinese Sci. Bull. 38 (1993) 1767.
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13 [9] H.J. Fecht, G. Han, Z. Fu, W.L. Johnson, J. Appl. Phys. 67 (1990) 1744. [10] Y.K. Wu, J.Y. Huang, A.Q. He, K.Y. Hu, X.M. Meng, Acta Metall. Sinica 29 (1993) B546 (in Chinese). [11] Z.T. Zhao, T. Liu, H.Y. Liu, in press.
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[12] M.B. Stearn, Phy. Rev. B 129 (1963) 1136. [13] Z.T. Zhao, T. Liu, G.W. Liu, R.Z. Ma, J. Mater. Sci. Lett. 15 (1996) 1427. [14] H. Nakatsu, S. Takaki, J. Japan Inst. Metals 60 (1996) 141 (in Japanese).
Mechanical alloying of Fe–Mn and Fe–Mn–Si T. Liu b, H.Y. Liu a,*, Z.T. Zhao c, R.Z. Ma c, T.D. Hu b, Y.N. Xie b a
Laboratory of Semiconductor Materials Science, Institute of Semiconductors, Chinese Academy of Science, Beijing 100083, People’s Republic of China b Beijing Synchrotron Radiation Facility, Institute of High Energy Physics, Chinese Academy of Science, Beijing 100083, People’s Republic of China c Department of Materials Physics, Uni6ersity of Science and Technology, Chinese Academy of Science, Beijing 100083, People’s Republic of China Received 7 May 1998; received in revised form 22 September 1998
Abstract The ball milling of Fe–24Mn and Fe–24Mn–6Si mixed powders has been performed by the high energy ball milling technique. By employing X-ray diffraction and Mo¨ssbauer measurements, the composition evolution during the milling process has been investigated. The results indicate the formation of paramagnetic Fe – Mn or Fe – Mn – Si alloys with a metastable fcc phase as final products, which imply that the Fe and Mn proceed a co-diffusion mechanism through the surface of fragmented powders. The thermal stability and composition evolution of the as-milled alloys were discussed comparing with the bulk alloy. © 1999 Published by Elsevier Science S.A. All rights reserved. Keywords: As-milled alloys; Ball milling; Mo¨ssbauer measurements; X-ray diffraction
1. Introduction Solid state amorphization reaction (SSAR) induced by mechanical alloying (MA) or mechanical milling (MM) has recently been an intensive research subject because of the great potential for producing amorphous materials or nanostructured materials for technological applications. However, the mechanism related to the process of MM or MA is still ambiguous [1,2]. In the past years mechanical alloying of binary metallic elements, or metal– metalloid systems have been widely investigated [3–6]. To our knowledge there have been no previous studies of mechanically alloyed Fe–Mn and Fe–Mn–Si reported. The Fe – Mn based alloys are of great importance for stainless steel, non-magnetism steel and which are also found wide application as important parts in Fe-based shape memory alloys. In the present study powders of Fe – 24Mn and Fe– 24Mn–6Si mixtures were ball-milled, the alloying concerned mechanism of the samples along with the milling process was investigated through X-ray diffraction pat* Corresponding author. E-mail address: [email protected] (H.Y. Liu)
terns and Mo¨ssbauer spectra. Because of its localized character, Mo¨ssbauer can give information about local environment of an Fe atom. Both Fe and Mn are transition metals and are located closely in the periodic table and the two elements have small mixing enthalpy, DHmix, of only +1 kJ mol − 1 [7]. Therefore, the chemical driving force for Fe–Mn alloying can be neglected. Also the atomic sizes of both elements in their bcc forms are very similar. Thus, one would expect that Fe–Mn is a non-fast diffusion system. The bcc, hcp and fcc phases coexist in the Fe–Mn binary alloy, depending on the Mn concentration, and can transform into each other by martensitic transformation.
2. Experimental details The Fe–Mn and Fe–Mn–Si alloys with normal compositions of Fe–24Mn and Fe–24Mn–6Si (wt.%), respectively, were prepared by appropriate masses of pure metal powders of iron (99.5%), maganese (99.7%) and silicon (99.9%). Ball milling of the powder mixtures was performed with a vibration mill using hardened
0921-5093/99/$ - see front matter © 1999 Published by Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 0 2 2 - 3
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
9
steel vials and balls. The ball-to-powder weight ratio was :20:1. The vial was loaded and sealed in an argon atmosphere. Annealing of the as-milled Fe – 24Mn–6Si was carried out in a tube furnace at 500, 650 and 700°C for 1 h, under argon atmosphere protection. X-ray diffraction measurements on the as-milled powders were carried out in a Rigaku D/max-RB X-ray diffractometer using the Cu – Ka radiation. Transmission 57Fe Mo¨ssbauer spectra were collected with a conventional constant-acceleration spectrometer using a 57 Co(Pd) source at room temperature. Velocity calibration was carried out using an a-Fe foil. The Mo¨ssbauer spectra were fitted by a program developed by Xu et al. [8]. Differential thermal analysis (DTA) was performed in the DTA-1700 (Perkin Elmer) at a heating rate of 10 K min − 1.
3. Results and discussion
3.1. Mechanical alloying of Fe– Mn Fig. 1 shows X-ray diffraction patterns of the as-filed and ball-milled Fe – Mn powders for 8, 27, and 67 h, respectively. Upon milling the diffraction intensities corresponding to the a-Fe and a-Mn tends to decrease and diminish gradually. In the same time a new phase appears with slightly different lattice parameters from those of a-Fe and a-Mn appears, which are identified as fcc Fe–Mn alloy. Apparently the ball-milling has induced the alloying of Fe and Mn. The diffraction peaks of Fe–Mn alloy are broadened due to the refinement of Fe–Mn particles. However, no trace of amorphization is noted even after 67 h of milling, this may be ascribed to the non-fast diffusion system that Fe and
Fig. 1. XRD patterns of Fe–24Mn, milled for 8, 24 and 67 h, respectively. The 67 h milled pattern indicates the appearance of fcc Fe– Mn.
Fig. 2. Mo¨ssbauer spectra of Fe – 24Mn, milled for 8, 24, 46 and 67 h, respectively. In the 8 h spectrum the sextet subspectra with a small hyperfine field and two singlet subspectra, employed to fit the overall spectra is shown.
Mn belong, that is, the condition for amorphization of binary alloys could not be satisfied [9]. In order to further investigate the local variation associated mechanism with the milling, the 57Fe Mo¨ssbauer spectra of the milled samples were obtained as shown in Fig. 2. A portion of Fe powders has fragmented into fine particles with paramagnetic behaviors upon 8 h milling, as demonstrated by the coexistence of ferromagnetic sextet and paramagnetic singlet peaks. At 27 h the fraction of the ferromagnetic phase becomes more unimportant and tends to diminish after 46 h of milling. The spectrum of 67 h corresponds only to the singlet line of the paramagnetic Fe–Mn phase. Even though the mixing enthalpy of Fe–Mn is not large, mechanical alloying supplies sufficient driving force for Fe–Mn alloying. Upon milling, the powders were fragmented into fine particles, as a result, Fe and Mn diffuse into each other, forming an fcc phase of Fe–Mn, and induces the material change from ferromagnetic to paramagnetic. Considering the character of the central singlet line, we employed two singlet subspectra to fit the paramagnetic phase, as indicated in Fig. 2 for the 8 h spectrum, the two subspectra can be regarded to stand for different local environments around Fe atoms. The fitted results of Mo¨ssbauer spectra for the Fe– Mn alloys milled for different time are listed in Table 1. The value of isomer shift (IS) for the paramagnetic phase is averaged over the two subspectra. As for the
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
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Table 1 Results of fitting Mo¨ssbauer spectra of Fe–Mn samples milled for 8, 27, 46 and 67 ha Time (h)
8 27 46 67 a
Sextet 1
Sextet 2
Singlet
IS
QS
Bhf
A
IS
QS
Bhf
A
IS
A
0.01 0.13 – –
0.01 −0.01 – –
32.8 33.0 – –
61.6 14.1 – –
0.13 0.23 – –
−0.12 −0.22 – –
23.7 28.3 – –
13.9 10.1 – –
−0.09 −0.07 −0.07 −0.07
24.6 75.8 100 100
IS: mm s−1, QS: mm s−1, Bhf: T, A: %, A: area fraction of subspectra.
ferromagnetic phase, besides the sextet subspectrum of a-Fe, another subspectrum with Voigt line can be resolved and fitted with a hyperfine field of 23.7 T according to the sink of the sextet peak. Upon milling some positions originally occupied by Fe atoms, especially for those Fe atoms at the surface of particles, may be taken place now by Mn atoms due to the diffusion, and thus leads to the decrease of hyperfine field, as seen in Table 1. Besides this, the IS and quadrupole splitting (QS) are quite different from those of the a-Fe. On milling one may expect that the Fe atoms on the surface may suffer from intense bump and the lattice distortion of Fe lattice due to the mismatch of lattice constants between Fe and Mn atoms, and thus leading to the deviation of 3d electron between Fe and Mn. In the same time Fe atoms diffuse into Mn particles. The Fe atoms surrounded by Mn may transform into paramagnetic Fe and make contribution to the fraction of paramagnetic phase. As the process goes on, i.e. attachment of Fe and Mn particles surface diffusion bulk diffusion, the fraction of paramagnetic Fe–Mn phase increases. On the other hand, the phase compositions undergo a great change such the bcc Fe transforms into fcc g-Fe(Mn) solid solution. Meanwhile the Fe atoms dissolve soluted in the a-Mn lattice and form the gMn(Fe). According to the TEM of Fe40Cu60 observed by Wu et al. [10], the a-Fe(Cu) was found to transform into g-Fe(Cu) upon milling. The authors explained this change as due to the martensitic transformation. However, the current results indicate that the phase transformation of Fe – Mn is most probably diffusion-controlled, though, the Fe – Mn is very likely to be a non-fast diffusion system.
found that its value changes little upon milling for both Fe–24Mn and Fe–24Mn–6Si. The IS for Fe–Mn–Si sample increases from −0.07 to − 0.01 mm s − 1 relative to that of Fe–Mn due to the addition of Si. Recently we have performed the Mo¨ssbauer measurements for bulk Fe–30Mn–6Si alloys prepared by melting and quenched from high temperature (900°C) [11]. The XRD patterns indicate the quenched strip-shaped sample consists of only fcc austenitic phase. By fitting the singlet line of Mo¨ssbauer spectra the IS value is given as − 0.01 mm s − 1, which is very close to that of
3.2. Mechanical alloying of Fe– Mn – Si The XRD patterns and Mo¨ssbauer spectra of Fe– 24Mn–6Si samples milled for 8, 27, 46 and 67 h are shown in Figs. 3 and 4, respectively. The results of fitting the Mo¨ssbauer spectra are listed in Table 2. Similar to that of Fe – 24Mn, the final product consists of the paramagnetic Fe – 24Mn – 6Si alloy with a metastable fcc phase. Considering the fitted IS, we
Fig. 3. XRD patterns of Fe – 24Mn – 6Si, milled for 0, 24, 46 and 67 h, respectively. The 67 h milled pattern indicates the appearance of fcc Fe – Mn.
T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
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6Si on the atomic level has not induced the amorphization of alloys as shown in the XRD patterns by the characteristic broadening of diffraction peaks. As we know, the g phase in Fe–Mn–Si shape memory alloy can only be stabilized at temperatures higher than about 900°C. At room temperature o phase precipitates due to g o martensitic transformation. However, the ball milling of Fe–Mn–Si reveals that ultrafine particles of g phase are relatively stable, which is quite different from the bulk alloy. One possible explanation is that the interfacial energy and strain energy have changed by mechanical alloying due to the formation of numerous structural defects. As a result, the fcc phase can be stabilized from the thermodynamic point of view. Similar explanation can also be adopted for Fe– 24Mn. According to the phase diagram, the stable Fe–Mn is composed of a and g phases at room temperature. However the XRD patterns indicate only the existence of g phase. In fact even cooling the milled Fe–Mn–Si samples down to a extremely low temperature of 1.5 K, no trace of o phase was found in the XRD patterns. We have investigated the reverse martensitic transformation of Fe–30Mn–6Si alloy [13] in a previous investigation. It is known that strain play a role as inductor of g o matensitic transformation, but our recent results indicated that intense bending would not induce the formation of a plenty of martensite; on the contrary, the g o martensitic transformation was suppressed. This is also true for ball-milling of Fe–Mn and Fe–Mn–Si, implying the strain-induced martensite is closely correlated with process. Similar results have been obtained by Hideshi et al. [14] when studied the martensitic transformation in Fe–Mn alloys. The critical point for martensite transforming into autensite was found to decrease while reducing the grain size of Fe–Mn alloys from 100 to 10 mm, the g o martensitic transformation was suppressed.
Fig. 4. Mo¨ssbauer spectra of Fe–24Mn–6Si, milled for 0, 24, 46 and 67 h, respectively. In the 8 h spectra the sextet subspectra with a small hyperfine field and two singlet subspectra, employed to fit the overall spectra are shown.
the ball-milled Fe – 24Mn – 6Si under investigation, despite the difference in Mn concentration. It seems that Mn has a less effect on the electronic arrangement of Fe atoms. With the addition of Si, the Si may take the place of Fe and forms a substitutional solid solution. As a result, the 3d electrons of Fe deviate from Si atoms and increase the IS value. A similar analysis was used to the Fe–Si alloys system studied by Stearn [12]. The results indicate that ball-milling has induced an intermixing of Fe, Mn and Si on an atomic level, and the Fe atoms seem to take the lattice position of Mn but not segregate to the grain boundary. It should be noted that after 46 h of ball-milling the system has reached a metastable equilibrium with a fcc phase. The intermixing of Fe – 24Mn and Fe –24Mn–
3.3. Thermal stability of as-milled Fe–Mn–Si The thermal effect associated with structural relaxation and grain growth of the ball-milled Fe–24Mn– 6Si samples was investigated by the DTA measured for
Table 2 Results of fitting Mo¨ssbauer spectra of Fe–Mn–Si samples milled for 8, 27, 46 and 67 ha Time (h)
8 27 46 67 a
Sextet 1
Sextet 2
Singlet
IS
QS
Bhf
A
IS
QS
Bhf
A
IS
A
0.01 0.00 −0.02 –
0.01 0.01 −0.01 –
33.2 33.1 33.0 –
47.4 14.6 2.8 –
0.05 −0.13 – –
0.08 0.24 – –
24.9 26.6 – –
38.4 8.4 – –
0.07 −0.01 −0.01 −0.01
14.2 77.0 97.2 100
IS: mm s−1, QS: mm s−1, Bhf: T, A: %, A: area fraction of subspectra.
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T. Liu et al. / Materials Science and Engineering A271 (1999) 8–13
Fig. 5. DTA thermogram of Fe–24Mn–6Si powder milled for 67 h (heat rate 40 K min − 1).
the powders milled for the 67 h sample at a heating rate of 40 K min − 1, as shown in Fig. 5. Two broad exothermal peaks are observed in the temperature range 510– 620°C and 670–730°C, respectively. A heat treatment was carried out for the as-milled sample at 500, 650 and 700°C, respectively, for 1 h. Fig. 6 shows the XRD patterns of the annealed samples. It shows that additional peaks emerge in the patterns of the fcc phase upon annealing, which are identified as a martensitic phase. The sharp peaks marked by X are hard to identify, which can probably be attributed to the contamination by the substrate. At 650°C the a phase has become dominant. However, for temperatures higher than 700°C, the fcc phase again plays a dominant role. Besides, the grain size of Fe – Mn powders shows a trace of growth-up with increasing temperatures when one take account of the broadening of diffraction peaks. According to the Fe – Mn phase diagram, the a phase or the g phase can be stabilized at certain temperature and alloy composition. Therefore, the first exothermal peak in the DTA trace can be ascribed to the ga transformation, and the second to the a g transformation, which is in agreement with the phase diagram of bulk Fe – Mn alloy.
Fig. 6. Room temperature XRD patterns of annealed Fe–24Mn–6Si powders at 500, 650 and 700°C, respectively, for 1 h.
Fe–Mn–Si. The metastable g phase of Fe–Mn or Fe–Mn–Si was found to be stable final product of ball-milling processes. No trace of o phase was detected as did in the bulk Fe–Mn–Si alloys. However, the as-milled Fe–Mn–Si powders show a similar thermal stability as in the phase diagram of bulk alloys.
References 4. Conclusions The Fe–Mn and Fe – Mn – Si alloys have been prepared by mechanical alloying of Fe – 24Mn and Fe– 24Mn–6Si at ambient temperature. The XRD patterns and Mo¨ssbauer spectra reveal that Fe and Mn atoms diffuse into each other through surface into bulk of ultrafine particles, and leading to the structural transformation and magnetic change upon milling of the Fe–Mn and Fe– Mn – Si powder mixtures. The phase transformation of the as-milled alloys shows a quite difference compared to the bulk alloy of Fe –Mn and
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