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Mechanical alloying of an Fe-based bulk metallic glass composition (Fe60Co8Zr10Mo5W2B15) ... E-mail address: csuryana@mail.ucf.edu (C. Suryanarayana).
Journal of Alloys and Compounds 389 (2005) 121–126

An unusual phase transformation during mechanical alloying of an Fe-based bulk metallic glass composition Umesh Patil, Soon-Jik Hong, C. Suryanarayana∗ Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA Received 7 July 2004; accepted 3 August 2004

Abstract Mechanical alloying of an Fe-based bulk metallic glass composition (Fe60 Co8 Zr10 Mo5 W2 B15 ) was conducted starting from blended elemental powders. A supersaturated solid solution of all the alloying elements in Fe has formed in the initial stages of milling. This ␣-Fe phase became amorphous on subsequent milling up to 20–30 h. Further milling, unexpectedly, led to primary crystallization of the amorphous phase resulting in the formation of an (␣-Fe) phase with a slightly smaller lattice parameter (and lower solute content) than that in the early stages of milling. External annealing of the amorphous phase powder at 700 ◦ C for 1 h led to complete crystallization of the amorphous phase into a mixture of phases including ␣-Fe, ZrFe2 , Fe2 B, Fe3 B, MoB2 , FeMo2 B2 , and at least another unidentified phase. Possible reasons for such an unusual amorphous → crystalline transformation on milling in the bulk metallic glass composition are discussed and it was shown that this transformation is due to the combination of defect-assisted diffusion and a slight rise in the powder temperature. © 2004 Elsevier B.V. All rights reserved. Keywords: Bulk metallic glass; X-ray diffraction; Fe-based alloy; Mechanical alloying; Phase transformation

1. Introduction Bulk metallic glasses are novel engineering materials which show an unusual combination of properties. They can be formed at very low solidification rates (0.1–100 K/s), have a high glass-forming ability, high thermal stability, and offer an opportunity to investigate the properties of materials in the supercooled region in addition to its completely glassy or crystalline states. Such metallic glasses have been synthesized in Zr-, Ln-, Mg-, Pd-, Al-, and Fe-based alloy systems. The structure, properties, and applications of bulk metallic glasses have been reviewed earlier [1–5]. Amongst these bulk metallic glasses, Fe-based alloys are of considerable interest for their potential magnetic and mechanical applications [6]. Even though fully glassy rods could be synthesized up to 72 mm in diameter in a Pd–Cu–Ni–P alloy [7], the maximum ∗

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section thickness that could be achieved in Fe-based alloys was only about 6 mm in the (Fe, Co)–Zr–(Mo, W)–B system [8], limiting large-scale high-volume applications for these materials [9]. Thus, alternative methods have been sought to be able to produce Fe-based glassy alloys in bulk quantities. Mechanical alloying (MA) is one such potential method. MA is a high-energy ball milling process involving repeated cold welding, fracturing, and rewelding of powder particles. This technique has been shown to be able to synthesize a variety of metastable phases including glassy alloys [10,11]. A few bulk metallic glass compositions have also been synthesized by MA [12]. Even though some investigators have reported on crystallization of the starting amorphous phase on MA, termed as mechanical crystallization [13], and others have reported crystalline-amorphous-crystalline cyclic transformations [14–16], we demonstrate an unusual amorphous → crystalline transformation of an amorphous phase synthesized by MA from blended elemental powders in an Fe-based bulk metallic glass composition.

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2. Experimental procedure Appropriate amounts of the pure elemental powders (99.9% purity) of Fe, Co, Zr, Mo, W, and B were mixed together to obtain the desired composition of Fe60 Co8 Zr10 Mo5 W2 B15 (the subscripts represent the atomic percentage of each of the elements in the powder mix). MA was conducted in a SPEX 8000 mixer mill using hardened steel vial and balls. For each experiment, 10 g of the blended elemental powder mix and 100 g of stainless balls were loaded into the milling container, thus maintaining a ball-to-powder weight ratio of 10:1 during milling. Loading and unloading of the powders was always conducted in an argon-filled glove box to minimize powder contamination. The progress of phase transformations in the powders with milling time was followed by X-ray diffraction (XRD) using a Rigaku Xray diffractometer using Cu K␣ radiation (λ = 0.15406 nm) at 40 kV and 40 mA settings. The XRD patterns were recorded in the 2θ range of 30–70◦ . To ensure that there were no problems with fluorescence, the XRD patterns were also recorded using Co K␣ radiation. No difference was noticed between the two sets of XRD patterns (except for the peak positions). Identification of the phases present and calculation of the lattice parameters were done using standard XRD procedures [17,18].

Fig. 1. X-ray diffraction patterns of the blended elemental powder mixture of Fe60 Co8 Zr10 Mo5 W2 B15 milled for different times in a SPEX mill.

3. Results and discussion 3.1. As-milled powder Fig. 1 shows the XRD patterns of the Fe60 Co8 Zr10 Mo5 W2 B15 powder blend as a function of milling time, up to 50 h. The initial powder blend shows all the expected major diffraction peaks. Since the crystal structure is the same and the lattice parameters are very close to each other for W and Mo, their peak positions could not be distinguished. With increasing milling time, the intensity of the diffraction peaks decreased and their width increased, due to particle refinement (reduction of crystallite size) and increase of lattice strain. This trend continued to beyond about 15 h of milling time. The diffraction pattern of the powder milled for 20 h shows mostly a broad halo, suggesting formation of an amorphous phase. However, a relatively low intensity peak may be noticed at 2θ ∼ = 40.5◦ , which by comparison with standard XRD patterns, has been identified to be either Mo or W. Thus, the amorphous phase formed at 20 h of milling time is lean in W and/or Mo. On continued milling, however, this W/Mo peak intensity decreased and was almost completely absent in the powder milled for 30 h. But, at longer milling times, a few diffraction peaks started appearing and this trend continued till 50 h, the maximum milling time employed in the present investigation. Fig. 2 compares the XRD patterns of the powder milled for 20 and 50 h in the SPEX mill. One can easily notice the sharp ␣-Fe peak in the powder milled for 50 h, whereas only a broad halo is present in the powder

Fig. 2. Comparison of the X-ray diffraction patterns of the Fe60 Co8 Zr10 Mo5 W2 B15 powder mixture milled for 20 and 50 h. Note that while the powder milled for 20 h shows only a reasonably broad halo, indicating formation of an amorphous phase, that milled for 50 h shows a clear ␣-Fe peak suggesting crystallization of the amorphous phase.

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milled for 20 h. These additional diffraction peaks in the powder milled for 50 h were identified as arising from the ␣-Fe phase. This observation suggests that the amorphous phase has started to crystallize on long-term milling. 3.2. Lattice parameters A close observation of the diffraction patterns from powders milled for different times (Fig. 1) indicates that the ␣-Fe peaks have shifted to lower 2θ angles with increasing milling time, suggesting an increase of the lattice parameter. Even though measurement of the lattice parameter is not very accurate, due to peak broadening, the trend clearly shows that the lattice parameter increases up to about 20 h and then starts to decrease (Fig. 3). This variation in the lattice parameter of the ␣-Fe solid solution can be related to its composition. In the initial stages of milling, the solute elements seem to be dissolving in the Fe lattice forming a supersaturated solid solution. Since all the solute elements have a larger atomic size (Co: 0.2506 nm; Zr: 0.3186 nm; Mo: 0.2726 nm; W: 0.2742 nm) than Fe (0.2482 nm), their dissolution is expected to increase the lattice parameter of the ␣-Fe solid solution. Boron atoms with a size of 0.188 nm will dissolve interstitially in the Fe lattice and this will also dilate the lattice. Thus, it is easy to understand that the lattice parameter has increased in the early stages of milling. Once the threshold amount of distortion of the lattice is achieved due to the size mismatch between the solvent and solute atoms, the crystalline lattice becomes destabilized and an amorphous phase is expected to form [19,20]. Further, accumulation of lattice defects may have aided formation of the amorphous phase. As noted above, milling of the powder beyond the formation of the amorphous phase (20–30 h) resulted in the reappearance of diffraction peaks of the ␣-Fe phase, suggesting that the amorphous phase has started to crystallize. It is well known, that metallic glasses, whether formed at very high cooling rates as by traditional melt spinning techniques or by relatively slow rates as in bulk metallic glasses, begin to crys-

Fig. 3. Lattice parameter of the ␣-Fe phase as a function of milling time for the Fe60 Co8 Zr10 Mo5 W2 B15 powder mixture.

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tallize in the primary mode through formation of the primarymetal-based solid solution [4,21–23]. That is, crystallization of the homogeneous amorphous phase in the Fe-based glasses occurs by the formation of the (␣-Fe) phase (with a composition different from the ␣-Fe phase formed prior to the amorphous phase) and the remaining amorphous phase acquiring a different composition (Amorphous ). This new Amorphous phase may subsequently crystallize by either the eutectic or polymorphous modes. Thus, the Fe-based amorphous phase synthesized by MA can be expected to transform to the equilibrium phases in two stages: amorphous phase → (␣-Fe) + amorphous and amorphous → equilibrium phases Formation of the (␣-Fe) phase in the present case confirms that primary crystallization has occurred here. It should also be noted that the amount of the (␣-Fe) solid solution formed is small in quantity, as evidenced by the relatively low intensity of the peaks. Thus, the solute elements got repartitioned into the amorphous phase (containing a higher solute content) and the (␣-Fe) solid solution (with a lower solute content). Consequently, the lattice parameter of the ␣-Fe solid solution started to decrease, once crystallization has started. 3.3. External heat treatment Since the amorphous phase contains several metals and one nonmetal component, it is most unlikely that complete crystallization of the amorphous phase will result in the formation of only the ␣-Fe phase. To confirm that the (␣-Fe) phase is the product of only the primary crystallization of the amorphous phase and that it is not the only phase in the final product, the amorphous powder (milled for 20 h) was heat treated externally for 1 h at 700 ◦ C. While the powder milled for 20 h shows mostly a broad halo (except for a low intensity Mo/W peak), the powder heat treated for 1 h at 700 ◦ C (Fig. 4) shows a number of sharp diffraction peaks. Comparison with standard XRD patterns of the possible phases that could be

Fig. 4. X-ray diffraction pattern of the Fe60 Co8 Zr10 Mo5 W2 B15 powder mixture milled for 20 h in a SPEX mill and annealed externally for 1 h at 700 ◦ C.

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present in an alloy of this composition suggests that the annealed powder now contains several phases such as (␣-Fe) , Fe3 B, Fe2 B, MoB2 , FeMo2 B2 , and ZrFe2 . A few additional peaks have also been noted indicating that at least another unidentified phase exists. A similar phase constitution was also observed in the bulk metallic glass of the same composition, but processed by the solidification route [24]. Since (␣-Fe) is the major phase in the crystallization product, and it is the first phase to form as a result of crystallization of the amorphous phase, it may be assumed that the presence of the (␣-Fe) phase in the milled powder suggests start of crystallization of the amorphous phase on milling the powder for 50 h; almost complete crystallization has occurred on annealing the powder at 700 ◦ C. 3.4. Possible reasons for the primary crystallization of the amorphous phase Four possibilities can be visualized for this unusual transformation in the present case. One is the temperature rise of the powder during MA. It has been clearly shown in several instances [10,11,25] that the powder temperature increases during milling. Even though some investigators assume a very large instantaneous local temperature rise, the global temperature rise was never reported to be more than about 200 ◦ C. Since the crystallization temperature of the present Fe-based metallic glass has been reported to be 950 K [24], it is unlikely that formation of the ␣-Fe phase during milling is due to an increase of the powder temperature to a value above that of the crystallization temperature of the amorphous phase. The second possibility is powder contamination. The MA literature contains several examples of impurity-stabilized crystalline phases, which have formed after the formation of the amorphous phase [11,26,27]. Since the amorphous phase in the present alloy powder has crystallized into all the expected intermetallic phases (in addition to the ␣-Fe phase) on external heat treatment, the amorphous phase can be assumed not to contain any substantial amount of impurities. Thus, this possibility can also be discounted. The third possibility is inverse melting, a phenomenon observed in Ti–Cr alloys [28,29]. Subsequently this transformation was also reported to occur in several transitionmetal (Ti, Zr, Nb, and Ta) based binary alloys, especially those synthesized by MA [30,31]. In this inverse melting phenomenon, heating a homogeneous metastable b.c.c. solid solution formed by MA to higher temperatures, produces an amorphous phase polymorphously. On further heating to still higher temperatures, the amorphous phase crystallizes into the b.c.c. phase again, i.e., the phenomenon is reversible. The stabilization of the amorphous phase below the inverse melting temperature was explained on the basis that the amorphous phase has a higher chemical short- and medium-range order than the corresponding crystalline phase. Consequently, destruction of this order (disordering of the amorphous structure), e.g., by heating (or continued milling) enhances its free

energy and favors formation of the equilibrium crystalline phase. In a process similar to what has been reported for inverse melting in the b.c.c. solid solutions above, it is possible that the crystalline → amorphous → crystalline phase transformation in the present Fe-based alloy could also be explained on the basis of an inverse melting phenomenon. But, the inverse melting phenomenon in all the systems so far has followed certain requirements. It occurs in binary transition metal alloy systems, starts from a homogeneous crystalline phase, the atomic co-ordinations of the amorphous and the equilibrium crystalline Laves phase are very similar, it is reversible, and for the kinetics of the amorphization reaction to be appreciable, the temperature at which amorphization occurs is reasonably high (746 ± 20 ◦ C in Ti–55 at.% Cr and 810 ◦ C in Nb–55 at.% Cr alloys). Further, the electron-toatom (e/a) ratio of the alloy exhibiting inverse melting was found to be between 5.0 and 5.5. The (metastable) b.c.c. solid solution is stable at e/a values 5.5. Most of the above criteria are not met in the present case and therefore, it may be concluded that the crystalline → amorphous → crystalline phase transformation in the present Fe-based alloy is not due to the inverse melting phenomenon. The last possibility is that MA has caused changes in the relative thermodynamic stabilities of the different phases and phase combinations, due to the introduction of mechanical energy into the system. It is well known that mechanical milling introduces a variety of lattice defects such as vacancies, dislocations, and grain boundaries [10,11] and since the concentration of these defects can be very high [32], the sequence of phase transformations observed in the present investigation can be explained with reference to the relative thermodynamic stabilities of the different phases or phase combinations. Fig. 5 represents a schematic free energy versus composition diagram indicating the relative positions of the amorphous, amorphous , ␣-Fe, and (␣-Fe) phases and the blended elemental powder mixture. In alloy systems with a large negative heat of mixing (as in the present case), the blended elemental powder mixture has obviously a high free energy (point 1) and so on MA, the powder blend is expected to change into a more stable configuration and thus reduce its free energy. Since a crystalline solid solution has a lower free energy than the blended elemental powder mixture, the first transformation product will be the ␣-Fe solid solution phase, corresponding to point 2 in the free energy curve. This ␣-Fe phase is a supersaturated solid solution of Fe with all the solute elements dissolved in it. Continued milling introduces a high density of defects and thus increases the free energy of the crystalline ␣-Fe phase to above that of the amorphous phase. Thus, the ␣-Fe phase becomes destabilized and the amorphous phase begins to be more stable (point 3). The phase mixture of the (␣-Fe) solid solution (containing a lower solute content) and the amorphous phase with a much higher solute content than what was started with (due to the crystallization of the initial amorphous phase) is likely to have a lower free energy than a

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Fig. 5. Hypothetical free energy vs. composition diagram to explain the observed transformations in the present investigation.

single amorphous phase (point 4) and that is why continued milling resulted in the primary crystallization of the amorphous phase. Thus, it could be concluded that formation of the (␣-Fe) phase after the formation of the amorphous phase is neither just because of the excessive temperature rise of the powder, nor because of powder contamination, nor due to inverse melting. Even though the powder temperature is much lower than the crystallization temperature of the amorphous phase, partial crystallization through the primary mode has occurred and this can be ascribed to the increased diffusion of solute atoms due to the presence of a high density of millinginduced defects. A slight powder temperature rise may also have contributed to this transformation. 3.5. Comparison with related observations A cyclic crystalline-amorphous-crystalline transformation has been reported earlier in Co–Ti alloys [14–16]. Starting with the elemental powders, the authors had synthesized a metastable intermetallic phase, which on continued milling became amorphous. This process of crystalline → amorphous → crystalline transformations was observed repeatedly. Even though an amorphous phase has transformed into a crystalline phase in the present investigation also, important differences may be noted between our results and those on Co–Ti alloys. In the case of the Co–Ti alloys, the amorphous phase has transformed into a metastable crystalline b.c.c.Co3 Ti phase, which has again changed into an amorphous phase. On continued milling, the b.c.c.-Co3 Ti phase transformed again into the amorphous phase. A stable crystalline phase had never formed in these alloys on milling; it formed only on externally annealing the amorphous alloy powder. However, in our present investigation, the ␣-Fe phase is an

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almost equilibrium phase (although the solute content is different) and on subsequent annealing the full equilibrium constitution has been achieved. Secondly, we have formed a solid solution phase on MA of the powder blend. Even though formation of an amorphous phase on milling a crystalline phase depends on the increase of free energy as a result of accumulation of lattice defects in the crystalline phase, it is not clear whether an amorphous phase can be produced more easily from the solid solution or the intermetallic phase. The literature has examples of both the situations [10,11]. However, since intermetallic phases are usually ordered and their crystal structures more complex than those of the primary solid solutions, it has been noted that it is easier to amorphize intermetallics than solid solution phases [33]. This is because it is easy to induce a lot of defects such as grain boundaries, dislocations, antiphase boundaries, and disordering of the lattice, all of which raise the free energy of the crystalline intermetallic above that of the amorphous and consequently, the amorphous phase becomes more stable. Thus, we believe that it is more difficult to form the amorphous phase from a solid solution (as in our Fe-based alloy composition) than from an intermetallic as in the Co–Ti system [14–16].

4. Conclusions It has been shown that an unusual phase transformation occurs on MA in an Fe-based alloy of bulk metallic glass composition. Starting with blended elemental powders, a supersaturated solid solution of solute elements in Fe was obtained in the early stages of milling. On milling to 20 h, an amorphous phase had formed, which on continued milling to 50 h, crystallized in a primary mode resulting in the formation of the (␣-Fe) phase and another amorphous phase with a higher solute content. Complete crystallization was achieved on heat treating the amorphous powder externally for 1 h at 700 ◦ C. Reasons for such an unusual transformation have been explained.

Acknowledgments This research work was supported in part by National Science Foundation under Grant # DMR-0314212. One of the authors (CS) is grateful to Dr.-Ing. Thomas Klassen of GKSS Research Center, Geesthacht, Germany, for fruitful discussions.

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