Crystallization kinetics and magnetic properties of Fe66Nb4B30 bulk metallic glass

Journal of Alloys and Compounds 483 (2009) 632–637

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Crystallization kinetics and magnetic properties of Fe66 Nb4 B30 bulk metallic glass M. Stoica a,∗ , S. Kumar a , S. Roth b , S. Ram c , J. Eckert a,1 , G. Vaughan d , A.R. Yavari e a

IFW Dresden, Institute for Complex Materials, P.O. Box 270116, D-01171 Dresden, Germany IFW Dresden, Institute for Metallic Materials, P.O. Box 270116, D-01171 Dresden, Germany c Materials Science Centre, Indian Institute of Technology, Kharagpur 721302, India d European Synchrotron Radiation Facilities ESRF, 38402 Grenoble, France e LTPCM-CNRS, Institut National Politechnique de Grenoble, 1130 Rue de la Piscine, BP 75, Saint Martin d’Hères Campus 38402, France b

a r t i c l e

i n f o

Article history: Received 30 August 2007 Accepted 12 November 2007 Available online 10 December 2008 Keywords: Amorphous materials Rapid solidification Phase transition Thermal analysis

a b s t r a c t Fe-based bulk metallic glasses (BMGs) have a high application potential because of their unique soft magnetic properties, mechanical behaviour and high corrosion resistance. Also, they can be obtained directly in the final shape suitable for use as magnetic sensors, magnetic valves, magnetic clutches etc. in different devices. Fe-based alloys able to form magnetic BMGs are of the type transition metal–metalloid and often contain 5 or more elements. Usually, the metalloid content is around 20 at.%. Recently, a new Febased BMG containing only 3 elements and a very high boron content was synthesized. The preparation of this BMG was done by employing the copper mold casting method and using the fluxing technique. This new BMG is ferromagnetic, with a Curie temperature around 550 K and a saturation magnetization of 105 Am2 /kg. Differential scanning calorimetry (DSC) investigations revealed a reduced glass transition temperature of 0.55 and an extension of the supercooled liquid region of about 31 K, values which indicate a relatively good thermal stability. Despite of numerous studies about Fe-based BMGs, there is still a lack of data about the crystallization kinetics. Also, the intermediate metastable phases, which form upon crystallization from the amorphous state, as well as the mechanism of their formation, are not fully understood. The present work discusses the kinetics of the phase formation using the Kissinger analysis and Johnson–Mehl–Avrami plots, correlated with the results obtained upon X-ray diffraction (XRD) of samples with different metastable structures. Additionally, the magnetic behaviour of different phase(s) is presented. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Generally, Fe-based glassy alloys are well known for their good soft magnetic properties depending on compositions, constituents and subsequent heat treatment of alloy. The soft magnetic properties arise because of absence of any crystalline anisotropy. Since the preparation of amorphous alloys in Fe-metalloid systems which exhibit good soft-magnetic properties in 1974 [1,2], a large number of studies on the development of soft-magnetic amorphous alloys have been carried out for the subsequent decades. However, the shape and dimension of Fe-based amorphous magnetic alloys had been limited to thin ribbon form with thicknesses below 30 ␮m because of the necessity of a high cooling rate of almost 106 K/s for the formation of an amorphous phase [3].

∗ Corresponding author. Tel.: +49 351 4659644; fax: +49 351 4659452. E-mail address: [email protected] (M. Stoica). 1 TU Dresden, Institute of Materials Science, D-01062 Dresden, Germany. 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.11.150

In 1995, a distinct glass transition before crystallization was found in the Fe72 Al5 Ga2 P11 C6 B4 rapidly solidified alloy [4], and an Fe73 Al5 Ga2 P11 C5 B4 ferromagnetic BMG was synthesized through the stabilization of supercooled liquid [5]. Subsequently, a variety of Fe-based ferromagnetic BMGs have been developed because of their potential magnetic applications [6,7]. Now, the development of Fe-based BMGs with high glass forming ability (GFA) has become a very hot research topic not only because of the soft-magnetic properties [8,9] but also of the high fracture strength [10,11] and corrosion resistance [12,13]. Fe-based alloys able to form magnetic BMGs are of the type transition metal–metalloid and often contain 5 or more elements. Usually, the metalloid content is around 20 at.%. In some cases, the magnetic properties of such BMGs can be enhanced by partial devitrification upon heating at a constant rate or by isothermal annealing. The change in magnetic properties is due to structural changes induced upon heating/annealing. Usually, Fe-based BMGs form intermetallic metastable phases at elevated temperatures, which finally transform into crystalline stable phases if the heating goes further. Very recently [14], we reported that the

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ternary Fe66 Nb4 B30 alloy may form BMG by copper mold casting. The BMGs of this composition can be prepared in rod-shape and the maximum diameter for which the alloy is still amorphous is 2 mm. Before casting, the arc-melted master alloy was cleaned by fluxing with B2 O3 . Despite several studies published in the literature about Fe-based BMGs and their magnetic properties, just few of them deal with crystallization behaviour and crystallization kinetics. The aim of this work is to present the crystallization behaviour of the ternary Fe66 Nb4 B30 and to link the structural changes with modification of the magnetic properties. 2. Experimental The ternary Fe66 Nb4 B30 BMG was prepared in several steps, using arc melting, flux-melting in induction furnace and copper mold injection casting. The procedures are described in detail elsewhere [14]. Additionally, amorphous ribbons with the section 4 mm × 30 ␮m were produced using the melt spinning apparatus, at 35 m/s tangential speed of the wheel. The glassy nature of the BMGs and the in situ crystallization behaviour was examined by X-ray diffraction in transmission configuration using a high intensity monochromatic synchrotron beam (with wave length  = 0.0155 nm) at ID11 at ESRF Grenoble. The structure of the as-cast and annealed melt spun ribbons were investigated by means of Bragg–Brentano diffractometer using a Co-K␣ radiation ( = 0.1788 nm). The thermal stability and crystallization kinetic were examined by differential scanning calorimetry (DSC), using both isochronal and isothermal mode. The variation of saturation magnetization with temperature was registered using the Faraday magnetometer, and the Curie temperature was calculated using the Hertzer’s approach [15]. The hysteresis loops and the saturation magnetization were measured by means of vibrating sample magnetometer (VSM).

3. Results and discussion The as-cast structure and thermal stability of the Fe66 Nb4 B30 was already studied and published in a previous work [14]. The 2 mm diameter rod samples were proved to be fully amorphous, upon X-ray diffraction in transmission configuration. Using the same apparatus, the in situ crystallization behaviour during heating with a constant rate of 20 K/min was recorded. Fig. 1(a) shows the diffraction spectra, from room temperature up to a temperature close to the melting point. There it is possible to observe that as the temperature increases the amorphous sample crystallizes and the crystalline product(s) are stable up to a given temperature, when they transform again. The observations are consistent with the DSC data (not shown here), which revealed a glass transition event with an onset temperature of 845 K, followed by a first crystallization event with onset at 876 K and a second crystallization event at 1067 K. The melting takes place through a peritectic reaction. Thus, the liquidus temperature, Tliq , measured as the onset of the last melting peak, is 1530 K. In order to study the phase formation, Fig. 1(b) shows in detail three diffraction spectra: one corresponding to the sample at room temperature, one recorded at 1000 K (above the first crystallization event) and one at 1300 K (above the second crystallization event). As can be observed, after the first crystallization some crystalline peaks appear, which are superimposed on the broad maxima characteristic of the amorphous phase. At this moment, the sample contains nanocrystals of a metastable Fe23 B6 -type phase and, eventually, some residual amorphous matrix. The phase has in fact the composition (Fe,Nb)23 B6 , with a fcc structure and a lattice constant a = 1.076 nm. As the temperature increases further, this metastable phase transforms completely into ␣-Fe, Fe2 B and FeNbB phases. Also, no residual amorphous matrix can be observed. It is interesting to follow the changes in magnetic properties as the structure changes. Fig. 2(a) shows the variation of the saturation magnetization MS with temperature, measured at a constant heating and cooling rate of 20 K/min. At a certain temperature, the saturation drops to zero. The Curie temperature of the amorphous phase is TC = 550 K. As the temperature increases and approaches the point where (Fe,Nb)23 B6

Fig. 1. Crystallization behaviour of Fe66 Nb4 B30 BMG, recorded in situ during heating with a constant rate of 20 K/min: (a) all patterns ranging from room temperature up to 1400 K and (b) details at three given temperatures (room temperature, 1000 K and 1300 K).

starts to form, the saturation increases again (inset in Fig. 2(a)), indicating that this phase is ferromagnetic. The cooling curve does not follow the heating curve, because of the presence of the crystalline magnetic phase. Due to experimental limitation, the heating did not reach 1035 K, the point where the (Fe,Nb)23 B6 phase transforms completely. A Curie temperature of 720 K can be determined from the point of inflection in the cooling curve of Fig. 2(a). This Curie temperature is attributed to the residual amorphous matrix (which has another composition than the starting one). (Fe,Nb)23 B6 has a Curie temperature larger than 900 K, as can be seen from the inset in Fig. 2(a). Thus, the Curie temperature of the emerging nanocrystalline metastable (Fe,Nb)23 B6 is much higher than the Curie temperature of amorphous Fe66 NB4 B30 and the Curie temperature of the amorphous phase increases considerably upon precipitating the crystalline (Fe,Nb)23 B6 phase. Taking into account that the saturation magnetization usually follows the trend of the Curie temperature [16], one also expects to observe an increase in saturation upon nanocrystallization. Fig. 2(b) compares the hysteresis loops recorded for two slices cut from a 2 mm diameter rod, fully amorphous at room temperature, and annealed at 900 K and 1300 K, respectively. The annealing temperatures were chosen to match the maximum temperature attained in the case of the Curie temperature and in situ crystallization measurements, respectively. For the amorphous sample, the coercivity is 1.5 A/m and the

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Fig. 3. The XRD patterns of Fe66 Nb4 B30 ribbons annealed for 10 min at 973 K and 1223 K.

Fig. 2. Fe66 Nb4 B30 BMG: (a) variation of the saturation magnetization with temperature during heating and subsequent cooling with 20 K/min, (b) hysteresis loops recorded for samples annealed 10 min at 900 K and 1300 K.

saturation magnetization 105 Am2 /kg. The sample annealed at 900 K shows a larger coercivity, but still very low (85 A/m), and a higher saturation (130 Am2 /kg). A further increase of the annealing temperature leads to macroscopic crystallization of the amorphous sample. Even if the saturation further increases to 145 Am2 /kg, the soft magnetic properties almost disappear: the coercivity is three orders of magnitude higher (5000 A/m) and the hysteresis loop is no more rectangular. For a better understanding of the phase formation in the case of this ternary BMG, amorphous ribbons of the same master alloy were cast. Ribbons with the same composition but from a notfluxed master alloy were also cast. The XRD and DSC investigations shown that the ribbons made from the fluxed alloy and the ribbons made from the not-fluxed alloy are completely amorphous and behave identical upon heating in the DSC. They behave similar also with the BMG (a distinct glass transition event, followed by two separate crystallizations and finally melting through a peritectic reaction), but the main temperatures and phases formed upon heating are somehow different. Fig. 3 shows the XRD spectra (recorded this time using the Bragg–Brentano configuration and a Co-K␣ radiation) of ribbons annealed at temperatures above the first and the second crystallization events, respectively, and Table 1 summarizes the main temperatures and phases formed upon heating for both BMG and ribbons. If at the first crystallization event only the Fe23 B6 -type phase forms in the case of the BMG, in the case of ribbons a mixture of Fe23 B6 -type, Fe2 B and ␣Fe was observed. Judging from the diffraction spectrum, the Fe2 B

phase seems to be the one with the highest volume ratio. Above the second crystallization event, both BMG and ribbon display a mixture of ␣-Fe, Fe2 B and FeNbB phases. As in the case of the BMG, the magnetic behaviour of the ribbons is strongly influenced by the structure. Fig. 4(a) shows the variation of the saturation magnetization with temperature, measured at a constant heating and cooling rate of 20 K/min. At a moment, the magnetization drops to zero, indicating that the Curie temperature of the amorphous phase was attained. Its value is 587 K, higher than in the case of BMG. As the temperature increases and approaches the crystallization temperature, the saturation increases again, in two distinct steps, which indicates the formation of different magnetic phases. As in the case of bulk sample, the cooling curve does not follow the heating curve, because of the presence of the crystalline magnetic phases. Due to experimental limitation, the heating did not reach the second crystallization event. The Curie temperature of the mixture Fe23 B6 -type, Fe2 B and ␣-Fe crystalline phases (formed above the first crystallization which take place at 925 K—see Table 1) is 911 K, a temperature which is higher with more than 200 K than the one measured for BMG heated above its first crystallization temperature. This is simply due by the differences in crystallization behaviour between rod and ribbon. The formed crystalline phases behave different from magnetic point of view. Fig. 4(b) compares the hysteresis loops recorded for ribbon at room temperature, and annealed at 973 K and 1223 K, respectively. The annealing temperatures were chosen to be above each crystallization event. For the amorphous sample, the coercivity is 2 A/m and the saturation magnetization 120 Am2 /kg. The sample annealed at 973 K shows a much larger coercivity, 2900 A/m, and a 15% higher saturation (138 Am2 /kg). A further increase of the annealing temperature leads Table 1 The main temperatures and the crystalline phases formed upon heating for Fe66 Nb4 B30 BMG and ribbon. Heating rate of 20 K/min

Rod

Ribbon

Tg (K) Tx1 (K) Tx2 (K) Tliq (K) Tx (K) Trg = Tg /Tx Melting Phases between Tx1 and Tx2 Phases above Tx2

845 876 1067 1530 31 0.55 Peritectic Fe23 B6 -type ␣-Fe, Fe2 B, FeNbB

897 925 1082 1530 28 0.58 Peritectic Fe23 B6 , Fe2 B, a-Fe ␣-Fe, Fe2 B, FeNbB

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Fig. 4. Fe66 Nb4 B30 ribbons: (a) variation of the saturation magnetization with temperature during heating and subsequent cooling with 20 K/min, (b) hysteresis loops recorded for as-cast samples and annealed 10 min at 973 K and 1223 K, respectively.

to complete crystallization of the amorphous sample. The saturation increases to 145 Am2 /kg but the soft magnetic properties disappear: the coercivity reaches 11170 A/m and the hysteresis loop is no more rectangular. Interesting is to compare the magnetic data measured for rod and ribbon and to discuss the differences between them. First of all, in the as-cast amorphous state both shows a very small coercivity, 1.5 A/m for BMG and 2 A/m for ribbon. The slightly higher value measured for ribbon can be explained taking in account that the Hc depend mostly on surface and volume pinning of magnetic domain walls. It has been reported [17] that due to surface irregularities the Hc is proportional to the ratio of the surface roughness amplitude to the specimen thickness. This contribution to Hc should be rather low for bulk samples because (i) the surface upon casting is very smooth and without scratches and (ii) the thickness is significantly larger than that of rapidly quenched ribbons. The saturation magnetization is also a bit higher in the case of ribbon than in the case of BMG, and can be explained by considering the contribution of the demagnetizing field: the ribbon has a more proper geometry for VSM measurements because its thickness is orders of magnitude smaller than the wide. Upon first crystallization event, the crystalline products are different in ribbon than in bulk, so the saturation is also different. Upon second crystallization, when the samples are supposed to reach their crystalline equilibrium phases, the saturation magnetization takes the same values, while the coercivity values are different (5 kA/m for bulk and 11.17 kA/m

Fig. 5. Fe66 Nb4 B30 ribbons: (a) isochronal DSC curves and (b) isothermal DSC curves.

for ribbon). A higher coercivity value may be a sign of a higher amount of non-magnetic phase in the sample, which give rise to a stray field [16]. With the help of amorphous Fe66 Nb4 B30 melt-spun ribbons, the crystallization kinetic was studied. Fig. 5(a) shows the continuous heating DSC curves of the sample with the heating rate varying from 5 K/min to 60 K/min. The peak temperatures Tp were further used to calculate the activation energy for crystallization using the Kissinger approach [18]: ln(/T2 ) = −E/RT + const., where  is the heating rate, R is the gas constant and T stands for the crystallization peak temperature. By plotting ln(/T2 ) versus 1/Tp2 , an approximately straight line was obtained, and from the slope the value of effective activation energy was calculated: Ec = 752 kJ/mol (±48 kJ/mol), which corresponds to 7.79 ± 0.49 eV. The isothermal crystallization kinetics of the ribbon prepared from composition Fe66 Nb4 B30 was further studied at different temperatures in the supercooled liquid region between 898 K (Tg + 1 K) and 909 K (Tg + 12 K). The corresponding DSC curves are shown in Fig. 5(b). Similar to continuous heating, all isotherms exhibit a single exothermic peak after passing a certain incubation period. It is assumed that the volume fraction transformed x, up to any time t, is taken as proportional to the fractional area of the exothermic peak.

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M. Stoica et al. / Journal of Alloys and Compounds 483 (2009) 632–637 Table 2 The kinetic parameters during isothermal annealing at different temperatures for an amorphous Fe66 Nb4 B30 ribbon. Kinetic parameters

Incubation time,  (min) Avrami exponent, n Reaction constant, k

Fig. 6. Fe66 Nb4 B30 ribbons: (a) volume fraction of the crystalline phases as a function of time and (b) the corresponding JMA plots.

So it is possible to measure the volume fraction of crystallization accurately by measuring the partial area of the exothermic peak. The measurement results for different temperatures are shown in Fig. 6(a), which are of typically sigmoidal-type curves for the crystallized volume fraction as a function of annealing time [19]. The sigmoidal-type curve mean that initially the crystalline volume fraction evolution with time is slow and then increases rapidly and finally slow down again or reaches saturation. During the initial part of the transformation, the nuclei are formed at different incubation times and this process continues with evolution of time depending on the nucleation rate during the overall transformation process. The saturation could be understood in terms of site saturation, i.e. no sites are available for further nucleation. This is because the matrix changes its composition too as the crystalline precipitates are formed. The shape of the sigmoidal curve at different temperatures is also different. The curve for the temperature close to glass transition shows a relatively slower rate of nucleation and growth than the curve for the temperature far away from Tg . This is only due to the higher viscosity at lower temperature, which leads to a slow diffusion of atoms. The isothermal crystallization kinetics of amorphous alloys is normally studied by the Johnson–Mehl–Avrami (JMA) equation [20]: x = 1 − exp{[k(t − )]n }, where x is the transformed volume fraction (%),  is the incubation time for the transformation defined as the time interval between the sample reaching the isothermal temperature and the initiation of the transformation, n is the Avrami exponent, and k is the reaction rate constant. The reaction rate constant k depends on temperature and can be described by Arrhenius equation k = k0 exp(−Ec /RT), where k0 is a constant and Ec is the apparent activation energy for crystalliza-

Annealing temperature (K) 898

901

903

905

909

2.350 3.384 4.850

1.717 3.560 3.698

1.450 3.549 3.095

1.500 3.365 2.332

0.783 3.589 1.457

tion. The values of k and n can also be determined by using the relationship [20]: ln{ln[1/(1 − x)]} = n ln k + n ln(t − ). By plotting ln{ln[1/(1 − x)]} versus ln(t − ) for different temperatures, the JMA plots can be obtained (Fig. 6(b)). The data for 10 < x < 85 (%) crystallized fraction shows two slopes for each temperature. The values of Avrami exponent n calculated from the higher slope of the JMA plots (for 40 < x < 85 (%) crystallized fraction) varies from 4.6 to 5.6 for this alloy. Normally, values of n > 4 are not considered in theories of phase transformation kinetics. Transformations for which 3 < n < 4 imply that the process is diffusion controlled with a nucleation rate which decreases with time [21]. Turnbull [22] suggested that if the nucleation rate increases with time, n > 4 might be obtained. The increasing nucleation rate can be explained by the phenomenon of self-heating of the specimen during transformation or a gradual short-range ordering of the glass, which changes the nucleation energetics [23]. The Avrami exponent n and the reaction constant k calculated from Fig. 6(b) only for the range 10 < x < 40 (%), on the basis of the slope and intercept of the straight line, are listed in Table 2. The plot ln k as a function of 1/T yields also a straight line. According to equation k = k0 exp(−Ec /RT), the activation energy for crystallization is Ec = 758 kJ/mol (±46 kJ/mol). This value is equal (within the errors) with the value determined from the Kissinger analysis for continuous heating of the sample. These results indicate that both the Kissinger analysis and the JMA approach can be successfully employed to determine the activation energy for the crystallization of this alloy. The value of activation energy for crystallization for our metallic glass composition is very high. For instance, the activation energies measured for Fe62 Co9.5 Gd3.5 Si10 B15 [24], Fe67 Co9.5 Nd3 Dy0.5 B20 [25] and Fe70 Nb2 Al5 Ga2 P11 C6 B4 [26] were 525 kJ/mol, 557 kJ/mol and 422 kJ/mol, respectively. Thus, even among ferrous alloys our glass shows a higher stability. In comparison with non-ferrous alloys, the activation energy obtained for the Fe-glasses is indeed extremely high since the activation energy was 245 kJ/mol, as measured for Zr55 Cu30 Al10 Ni5 [27] and 356 kJ/mol, as measured for gas atomized powder of Cu47 Ti33 Zr11 Ni8 Si1 [28]. In the present studied case, the high value of activation energy may be due by the combination of two distinct phenomena: the crystallization itself takes place very quick and results in the formation of a crystalline metastable phase which immediately transforms into another crystalline metastable phase or a mixture of different crystalline phases. In support of this affirmation come the DSC traces. Some of the crystallization peaks measured in the case of isochronal experiments (Fig. 5(a)), and more clearly seen in the case of isothermal experiments (Fig. 5(b)), have the tendency to split. This can be an indication that several transformations may take place in a very short time interval and, as a result, the calculated activation energy in fact contains both energies, the one necessary to form the primary crystallization phase(s) and the one corresponding to a further transformation. The splitting of the crystallization peak of Fe-based BMGs is not an “exotic” event, it was also found in other several cases [7,29], but up to now the scientific literature lacks in data concerning the crystallization kinetics for ferromagnetic BMGs. The incubation time is also related to these phenomena and, as was seen (Table 2) it is the order of only few

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minutes, in fact less than 150 s. Taking in account the high values of Avrami exponent, one can conclude that for this metallic glass the crystallization is diffusion controlled until the crystalline fraction reaches almost 40% of the entire volume. After that, the nucleation process becomes the driving force for crystallization. During the nucleation some new crystalline phases forms, which may explain the splitting of the peaks in the DSC traces. 4. Conclusions Bulk amorphous Fe66 Nb4 B30 soft magnetic alloys may be successfully prepared using liquid metallurgy route. The combination of a high thermal stability against crystallization with a low coercivity and a high saturation magnetization, make this glass an excellent material for applications as soft magnetic components in electro-magnetic conversion devices. The amorphous structure of as-cast samples and the nanocrystalline structure of devitrified samples were investigated by means of X-ray diffraction in transmission configuration using the synchrotron radiation and in reflection mode using Co-K␣ radiation. Thermal stability and nanocrystallization of samples were studied and it was found that they require very high activation energy for crystallization (almost 760 kJ/mol). Regarding the glass forming ability, the Fe66 Nb4 B30 shows a supercooled liquid region of around 31 K and a reduced glass transition temperature Trg = 0.55. Comparing the bulk samples with ribbon samples of the same composition, they have a different crystallization behaviour. Both shows very high glass transition temperatures and Curie temperatures and this make them suitable for applications which require a continuous operation at high temperatures. Acknowledgements The authors thank S. Venkataraman and K. Biswas for helpful discussions and to B. Bartusch, S. Donath and H. Schulze for technical help. The financial support of the EU trough the European

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