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Fabrication of FINEMET bulk alloy from amorphous powders by spark plasma sintering
Powder Technology 289 (2016) 163–168
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Fabrication of FINEMET bulk alloy from amorphous powders by spark plasma sintering T. Gheiratmand a,⁎, H.R. Madaah Hosseini a, P. Davami a, C. Sarafidis b a b
Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11155-9466, Azadi Ave., Tehran, Iran Department of Physics, Aristotle University of Thessaloniki, Thessaloniki, Greece
a r t i c l e
i n f o
Article history: Received 26 March 2015 Received in revised form 27 November 2015 Accepted 28 November 2015 Available online 2 December 2015 Keywords: FINEMET bulk alloy Amorphous materials Spark plasma sintering Viscous flow
a b s t r a c t Finemet bulk soft magnetic alloy was fabricated by spark plasma sintering of the milled ribbons. The amorphous melt-spun ribbons were milled for 36 min by high energy vibrational mill and then sieved to separate particles smaller than 125 μm. The size distribution of particles was determined by a laser diffraction particle size analyzer. Spark plasma sintering was carried out at super-cooled liquid region for short times of 7 and 21 min. The structure of bulk samples was characterized using X-ray diffraction, scanning electron microscopy, differential scanning calorimetry and transmission electron microscopy techniques. The magnetization and coercivity of samples were measured using SQUID magnetometry. The results indicated that during sintering the Fe(Si) phase with grain size of 9 nm forms in the amorphous matrix. The amount of Fe(Si) phase was calculated as 84 wt.% in the sample consolidated from fully amorphous powder for 7 min. The magnetization and coercivity of this sample was measured as 122.29 emu g−1 and 123 Am−1, respectively. The relative density of consolidated samples for 7 and 21 min was reached to 96 and ~97% of theoretical density, respectively. Recording the punch displacement with respect to the time showed that the particles rearrangement and viscous flow with activation energy of 23.4 kJ mol−1 were the main mechanisms of densification at the first and second stages of densification, respectively. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Over the past decades, Finemet soft magnetic alloy has attracted considerable technological attention due to the excellent soft magnetic properties such as high magnetization, low coercivity and high Curie temperature. These properties originate from the particular microstructure characterized by randomly distributed Fe(Si) nanograins in the amorphous matrix [1]. These alloys are generally synthesized by controlled annealing of an amorphous precursor [2,3]. The production of amorphous precursor requires very high cooling rates, thus only powders, wires and ribbons with maximum thickness of ~ 30 μm could be achieved which are not suitable in industrial applications where a large volume of magnetic materials is required [4]. Producing bulk amorphous alloys by direct casting methods such as suction casting is dependent on the glass forming ability of the alloy. In addition; the maximum size of the magnetic alloys produced by this technique is seriously restricted to ~ 10 mm [5]; while powder metallurgy techniques could produce bulk amorphous alloys in the larger sizes. The consolidation of the amorphous powders is carried out in the super-cooled liquid region between Tg (glass transition temperature) and Tx (onset of
crystallization temperature). In this region, the amorphous materials soften and the viscous flow allows the compaction of amorphous powders [6–8]. Spark plasma sintering (SPS) is a unique technique for fabrication of dense and near net shape bulk amorphous alloys where a DC on–off pulse is used to produce spark plasma for instantaneous heating of powders. The current passes through the die as well as the sample enabling the sample to be heated homogenously from both inside and outside [9]. The very high heating rates in the order of 100 °C min− 1 as well as the short holding times in the range of 3–10 min outstand SPS among other conventional vacuum hot pressing techniques in the prevention of crystallization and grain growth [5,10–14]. Despite too many studies have been done on the magnetic properties of Finemet ribbons and powders, very few investigations has been reported about consolidation of these materials. In this research, Fe73.5Si13.5B9Nb3Cu1 bulk alloy was fabricated by SPS of amorphous melt-spun powders. The microstructure and magnetic properties as well as the mechanisms responsible for the densification and the attendant structural changes during SPS were also investigated. 2. Experimental procedure
⁎ Corresponding author. E-mail addresses: [email protected], [email protected] (T. Gheiratmand).
http://dx.doi.org/10.1016/j.powtec.2015.11.060 0032-5910/© 2015 Elsevier B.V. All rights reserved.
Finemet powder for consolidation was produced by mechanical milling of melt-spun amorphous ribbons. The detail of processing
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procedure has been given in our previous publications [3,15–17]. Milled powders with different milling times of 36 and 45 min were sieved to obtain different particle size ranges. Particles smaller than 125 μm were used for consolidation. The milled powder for 36 min (SPS1 and SPS3) was fully amorphous while that milled for 45 min (SPS2) was partially crystallized during milling. The degassing of powders was conducted in a vacuum furnace up to 10−5 Torr at 310 °C for 50 min. The amorphous powders were consolidated using SPS equipment (model: Easy fashion) at temperature between Tg (glass transition temperature) and Tx (onset of crystallization). The powders were charged in a graphite die with 13 mm inner diameter lined by graphite paper. It was then inserted in a vacuum chamber evacuated up to 6 × 10−2 Torr at room temperature. The temperature rose up to 350 °C, then the powders were pressed under the pressure of 40 MPa. Afterwards, the temperature of samples (SPS1 and SPS2) increased to the consolidation temperature by the rate of 30 k min−1 which took 7 min. In order to prolong the consolidation time, the temperature of another sample (SPS3) was increased from 350 °C to 560 °C by the rate of 10 k min−1. The duration of consolidation in this sample was taken as 21 min. The amperage and voltage for consolidation were selected in the range of 1500–2500 A and 2–4 V, respectively. Consolidated alloys were sanded and polished in order to remove graphite paper and carburized layer from the surface. The size distribution of particles was determined by a laser diffraction particle size analyzer (PSA, NanoTecANALYSETTE22 FRITSCH). The particle size at 50% of cumulative curve, D50, was found out. In addition, D10 and D90 were reported. The density of bulk samples was measured by the Archimedean method using ethanol. The microstructure of both degassed powders and consolidated samples was investigated by X-ray diffractometer, XRD (STOE D-64,295) using Co-Kα radiation with a step size of 0.02°. Thermal stability of milled and compacted samples was conducted by DSC analysis performed on a METTLER TOLEDO TGA/DSC1 device at heating rate of 20 °C min−1 under N2 atmosphere. Tg and Tx were measured as 465 °C and 565 °C, respectively. The microstructure of sintered samples was characterized using a Scanning Electron Microscope, SEM (VEGA-TESCAN) and an optical microscope. Transmission electron microscopy observations were carried out on TEM (Tecnai G2 20) operating at 200 kV. The TEM sample was thinned by ion milling. Magnetic properties were determined at Room temperature using a superconducting quantum interference device (SQUID, Quantum Design MPMS-5) at external magnetic field up to 20 kOe. To determine the coercivity in the unit of Am− 1 Helmholtz coil device was applied.
Fig. 1. Cumulative passing volume percent of amorphous powder and its particle size distribution.
3. Result and discussion Fig. 1 shows the cumulative passing volume percent of amorphous powder together with the particle size distribution. The particles have a bimodal distribution with a D50 of 7 μm. Moreover, D10 and D90 have been measured as 2 and 17 μm, respectively. This implies that the size of 10% of particles is below 2 μm and 10% is larger than 17 μm. The DSC curves of consolidated samples together with those of milled ribbon for 45 min and melt-spun ribbon are presented in Fig. 2. The melt spun ribbon data has two exothermic peaks. The first peak at the lower temperature is ascribed to the formation of the Fe(Si) phase while the second one corresponds to the formation of Fe23B6/Fe3B boride phases. In the plot of ribbon milled for 45 min, three exothermic peaks could be observed at 560, 712 and 900 °C which are attributed to the formation of Fe(Si), Fe3B and Fe2B phases, respectively. The area under the first peak of sintered samples decreased indicating the formation of Fe(Si) phase during sintering. Actually, the crystallization enthalpy of Fe(Si) phase decreases from 60.06 Jg−1 in the melt-spun ribbon to 9.21 and 4.68 Jg−1 for SPS1 and SPS2 samples, respectively. The Fe(Si) crystallization peak of SPS3 sample is so broad that its enthalpy could not be measured. This implies that the high fraction of Fe(Si) phase has been formed during sintering. In addition, Fe3B/Fe23B6 crystallization enthalpies are 38.28, 38.21, 13.02 and 38.27 Jg−1 for melt-spun ribbon, SPS1, SPS2 and SPS3 samples, respectively. Since the structure of SPS1 and SPS3 samples before sintering was similar to the melt-spun ribbon, their crystallization enthalpies are almost identical. The decomposition of Fe23B6 phase into the mixture of Fe3B and Fe2B phases occurred during milling for 45 min, which causes the crystallization enthalpy of SPS2 sample to become lower. The fraction of crystalline phase formed in the sintered samples was calculated using the heat change of crystallization peaks in DSC curves. These values are 84 wt.% and 90 wt.% for SPS1 and SPS2 samples, respectively. Fig. 3 shows the XRD pattern of the sintered samples by SPS as well as degassed powders. As could be seen, the patterns of both degassed samples show typical broad peak, characteristic of amorphous materials, indicating that no significant phase change occurs during degassing. The structure of milled sample for 45 min was partially crystalline before degassing, i.e. small amounts of Fe(Si) nanocrystals together with Fe2B and Fe3B phases developed in the structure during milling. The structure of all three sintered samples shows the formation of Fe(Si) phase. The higher intensity of Fe(Si) characteristic peaks in the SPS3 than SPS1 sample confirms the formation of more Fe(Si) crystalline phase. This result is the consequence of longer time of keeping SPS3 sample at high temperature of consolidation. Thus, based on the DSC and XRD analyses, it could be concluded that the fraction of Fe(Si) phase in SPS3 sample is higher than the amount of 84 wt.% obtained
Fig. 2. DSC curves of as milled ribbons and spark plasma sintered samples.
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Fig. 3. XRD patterns of degassed powders and spark plasma sintered samples.
in SPS1 sample. It also seems that the amount of boride phases does not change significantly during sintering. The Fe(Si) crystallite size in SPS1 sample was calculated by Scherrer's relationship after eliminating the instrumental errors as ~8 nm. The density of bulk samples was determined using Archimedean method which is summarized in Table 1. The densities were an average of three measurements. It is obvious that relative density of fully amorphous samples is higher than the partially crystallized one. Despite the high density of SPS3 sample obtained due to the longer time of consolidation, the high fraction of Fe(Si) crystalline phase formed in this sample could be detrimental for soft magnetic properties. This will be more discussed later where the magnetic properties are given. The microstructure of bulk samples, characterized by SEM and optical microscope, are illustrated in Fig. 4. In the SPS1 and SPS3 samples, which were initially fully amorphous, the density of porosities was measured by image analyzer software as ~3 and 2.5 vol.%, respectively. In the microstructure of partially crystalline SPS2 sample more percentage of porosities could be observed. The detachment of some fine particles during polishing, marked by arrowed lines, is indicative of weak bonding between them. Achieving a more densified bulk alloy in the SPS3 sample could be attributed to the viscous flow of the amorphous powder above glass transition temperature under high pressures. Indeed, heating in the supercooled liquid region causes the homogenous flow of the amorphous material which results in filling the gaps between particles. Since the SPS2 sample was not fully amorphous, the viscosity of the amorphous portion in the glass transition temperature range was not sufficiently low to be completely densified by viscous flow mechanism. It has been reported that the viscosity of amorphous alloy decreases by hundredfold above Tg in the supercooled liquid region [5,18]. The wide distribution of powder particles could also be distinguished from optical and SEM images. It seems that the smaller particles fulfill the porosities between larger particles causing the density to increase. Thus, as microstructural features illustrated, the particle size distribution has significant influence on the sintering behavior of Finemet amorphous powders. Moreover, using PSA data, D50, D10 and D90 have been measured as 7, 2 and 17 μm, respectively. Therefore, several larger
Table 1 Density of the consolidated powders. Sample
Density (g/cm3)
Relative density
SPS 1 SPS 2 SPS 3
7.01 6.77 7.08
96.02 92.73 96.98
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particles observed on the SEM and optical microscope images could also be the result of connecting the smaller particles by sintering. The other important parameter affecting on magnetic properties is the location of porosities. In general, for high permeability, highfrequency and low-loss applications, the most dense structure with the uniform grain size is desirable. Any porosity that exists should be inter-granular so that the domain walls cross them without any difficulty. The permeability of a soft magnetic material is mainly determined by domain wall displacement as well as the spin rotational magnetization. The spin contribution is specified by sintering density while domain wall displacement is mostly determined by grain size. Actually, porosities are like a second paramagnetic phase which reduce the saturation magnetization but increase the electrical resistivity of soft magnetic material. At high frequencies, porosities act as an insulating barrier for particle to particle eddy current paths [19–21]. In nanocrystalline soft magnetic alloys the average crystallite size is smaller than the single domain size (35 nm for Fe base alloys). Hence, each domain comprises from several crystallites [1]. The sintering is performed by connecting of particles which contain too many magnetic domains. Existence of fewer pores within the particles and their aggregation along the boundaries means that there are fewer obstacles for pinning the domain walls during magnetization cycle. As could be seen in Fig. 4(b) and (f) the porosities are mostly aggregate between powder particles marked by circles. Therefore, it could be concluded that the amount of pores should be minimized to keep the magnetization at the highest level and demagnetizing field at the lowest level. The structure of SPS1 sample has also been investigated by TEM as depicted in Fig. 5. As could be observed, the microstructure of this sample consists of Fe(Si) crystalline phase embedded in the amorphous matrix. Mean crystallite size is about 9 nm which is in good agreement with XRD results. According to diffraction pattern, the structure of nano crystallites is BCC. It could be obviously seen that there are amorphous regions around nano crystallites which have a different composition from primary amorphous matrix as a result of chemical partitioning. The densification of amorphous powders during SPS may be sometimes the result of surface melting of powder particles. This densification mechanism is associated with devitrification of the amorphous powder and formation of both Fe(Si) and boride phases. Since no crystalline peak corresponding to the boride phases was detected in the XRD pattern after degassing, temperature never reached 600 °C. The magnetization and coercivity of bulk samples are listed in Table 2. The higher magnetization of SPS3 sample could be related to its higher density and consequently, the lower magnetostatic (shape) anisotropy. The high amount of Fe(Si) crystalline phase formed during consolidation in this sample should be also regarded in increasing the magnetization. The surface poles, which are distributed on the surface of porosities, produce internal fields in addition to the external fields. The internal field known as the demagnetizing field (Hd) is proportional to the magnetization of the body with direction opposite to it. Hd is sensitive to the shape of sample. Thus, the high density of porosities in a bulk sample leads to the higher demagnetizing field and the lower total magnetization. The coercivity of sintered samples was measured as 123, 189 and 147 Am−1for SPS1, SPS2 and SPS3 samples, respectively. Porosities act as paramagnetic phases which lock the domain walls and prevent their easy rotation during magnetization cycle. Besides, an enhancement of coercivity occurs. Moreover, the demagnetization fields produced by porosities are effective in increasing the coercivity. Furthermore, the high volume fraction of Fe(Si) crystalline phase could increase the magnetostriction and thus the coercivity. It should be noted that about 75–80% [22] of Fe(Si) crystalline phase embedded in amorphous matrix is needed to have the total near zero magnetostriction. It should be said that although the value of coercivity in bulk sample is higher than that reported for Finemet ribbons (1–16 Am−1 [23–26]), it is much smaller than that of toroidally winded ribbons which are not suitable for industrial applications where a large volume of magnetic materials with a complex shape is required.
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Fig. 4. Optical micrographs and SEM images of spark plasma sintered samples consolidated from powders milled for a) and b) 36 min and sintered for 7 min, c) and d) 45 min, sintered for 7 min, e) and f) 36 min and sintered for 21 min.
where B is a term that collects materials and geometric constants, R is the particles radius, n is the representative of dominant densification mechanism and m relates to the particles size. These parameters for different densification mechanisms are listed in Table 3 [27]. According to Table 3, if the shrinkage and time be linearly related to each other (as Eq. (2)) the dominant densification mechanism will be viscous flow [28].
The values of microhardness of bulk samples are given in Table 2 as well. It could be seen that SPS3 sample has the highest value of microhardness which can be explained on the base of the high density and high volume fraction of crystalline phase formed during consolidation. In addition, the microhardness of SPS2 sample is slightly higher than that of SPS1. This result has been obtained despite the improved density of SPS1 sample and could be explained based on the higher fraction of the crystalline phase formed during sintering in the sample sintered from the partial crystalline powder. It has been shown that the hardness of the sample consisting of the crystalline phase is much higher than that composed of fully amorphous phase [3]. In order to perceive the densification mechanism of Finemet powders by SPS, the shrinkage of SPS1 sample with respect to sintering time has been plotted in Fig. 6. It was proposed that relationship between the shrinkage (ΔL/L0) and sintering time (t) is stated as follow [27]:
η ¼ η0 expðQ v =RT Þ
n=2 ΔL B ¼ n mt L0 2 R
As stated earlier, in order to sinter the powders the temperature increased up to 350 °C before applying the pressure and then it rose up by the heating rate of 30 k min−1 during applying the pressure up to the
ð1Þ
ΔL 3γ t ¼− L0 4Dη
ð2Þ
Here γ is the surface energy, D is the particle size and η is the viscosity of amorphous materials. In addition, the viscosity of amorphous materials is temperature dependent and could be stated by Arrhenius equation as follow [28]: ð3Þ
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seems that at the first stage of densification, where the temperature of powder is below the glass transition temperature and not high enough for viscous flow, the rearrangement of particles is the main mechanism of densification. Since the size of milled amorphous powder particles was selected below 125 μm, the smaller size particles get redistributed at some of the interparticle spaces between larger amorphous particles. However, a significant number of smaller interparticle spaces still remain empty. The full densification then accomplished by more rising the temperature as well as applying the pressure above Tg, where the viscosity of the amorphous powder becomes smaller and the viscous flow of the amorphous particles is the main reason for increasing the sample density. As could be seen in Fig. 6(b), the relationship between the shrinkage and time at the second stage of sintering could be distinguished almost linear. Furthermore, the slope of the straight line 0Þ LnðdðΔL=L Þ versus 1/(623 + 0.5t) allows the calculation of the activation dt energy for viscous flow as 22.9 kJ mol−1. This value is in good agreement with the literature reports [29–32].
4. Conclusion Fig. 5. TEM bright field image and SAED pattern of spark plasma sintered consolidated from powder milled for 36 min.
Table 2 Saturation magnetization, coercivity and hardness of different sintered samples. Sample
Magnetization (emug−1) ± 1.0
Coercivity (Am−1) ± 1
Hardness (VHN) ± 10
SPS1 SPS2 SPS3
122.29 118.85 124.01
123 189 147
1097 1110 1118
maximum value of 560 °C. Thus, the relationship between time and temperature can be declared as T = 623 + 0.5 t. Thus, Eq. (2) could be written as follow: dðΔL=L0 Þ −3γ Qv ¼ Ln − Ln dt 8Dη0 Rð0:5t þ 623Þ
ð4Þ
Fig. 6(a) presents the shrinkage of Finemet amorphous powder during SPS. The densification of amorphous powder begins immediately upon applying the pressure by movement of punch. As could be seen in Fig. 6(a), the plot of shrinkage versus time is not a single straight line indicating the act of different mechanisms during sintering. It
This work sought to examine the possibility of production of Finemet magnetic alloy from amorphous powders by SPS method. The properties of produced Finemet bulk sample has been also investigated, which according to the authors' best knowledge has not been discussed elsewhere. Finemet bulk alloy with nearly ~ 97% relative density was successfully fabricated from amorphous powders by SPS in the supercooled liquid region for 21 min. The density of the sample consolidated from fully amorphous powders was higher than that obtained from partially crystallized one due to the lower viscosity and consequently, more homogenous flow of the fully amorphous material in the glass transition range. The wide distribution of powder particles was observed in the microstructure of sintered samples which is the result of particle size distribution and also smaller particles connection during sintering. The aggregation of porosities along the particle boundaries increases the resistivity and provides an insulating barrier for particle to particle eddy current paths at high frequencies. The microstructure in the bulk sample consisted of Fe(Si) nanocrystals with average size of 9 nm and a BCC structure embedded in the amorphous matrix. The fraction of crystalline phase was estimated ~ 84 wt.% from DSC measurements. The magnetization and coercivity in the sintered sample was measured as 122.9 emu g−1 and 123 Am−1, respectively. By applying the pressure, two distinct consolidation mechanisms were distinguished during the SPS process based on the relationship between shrinkage (ΔL/L0) and sintering time (t). At initial period of pressing up to 90 s, the densification was due to the particles rearrangement.
Fig. 6. (a) The variation of shrinkage with respect to the time, (b) plot of Ln(d(ΔL/L0)/dt) versus 1/(0.5 t + 623) corresponding to the second stage of densification by viscous flow.
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Table 3 The values of n, m and B for different densification mechanisms [27]. Mechanism
n
m
B
Viscous flow Plastic deformation Volumetric diffusion Grain boundary diffusion Surface diffusion
2 2 5 6 7
1 1 3 4 4
3γ/2η 9πγbDv/RT 80DvγΩ/RT 20δDBγΩ/RT 56DsγΩ4/3/RT
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Powder Technology journal homepage: www.elsevier.com/locate/powtec
Fabrication of FINEMET bulk alloy from amorphous powders by spark plasma sintering T. Gheiratmand a,⁎, H.R. Madaah Hosseini a, P. Davami a, C. Sarafidis b a b
Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11155-9466, Azadi Ave., Tehran, Iran Department of Physics, Aristotle University of Thessaloniki, Thessaloniki, Greece
a r t i c l e
i n f o
Article history: Received 26 March 2015 Received in revised form 27 November 2015 Accepted 28 November 2015 Available online 2 December 2015 Keywords: FINEMET bulk alloy Amorphous materials Spark plasma sintering Viscous flow
a b s t r a c t Finemet bulk soft magnetic alloy was fabricated by spark plasma sintering of the milled ribbons. The amorphous melt-spun ribbons were milled for 36 min by high energy vibrational mill and then sieved to separate particles smaller than 125 μm. The size distribution of particles was determined by a laser diffraction particle size analyzer. Spark plasma sintering was carried out at super-cooled liquid region for short times of 7 and 21 min. The structure of bulk samples was characterized using X-ray diffraction, scanning electron microscopy, differential scanning calorimetry and transmission electron microscopy techniques. The magnetization and coercivity of samples were measured using SQUID magnetometry. The results indicated that during sintering the Fe(Si) phase with grain size of 9 nm forms in the amorphous matrix. The amount of Fe(Si) phase was calculated as 84 wt.% in the sample consolidated from fully amorphous powder for 7 min. The magnetization and coercivity of this sample was measured as 122.29 emu g−1 and 123 Am−1, respectively. The relative density of consolidated samples for 7 and 21 min was reached to 96 and ~97% of theoretical density, respectively. Recording the punch displacement with respect to the time showed that the particles rearrangement and viscous flow with activation energy of 23.4 kJ mol−1 were the main mechanisms of densification at the first and second stages of densification, respectively. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Over the past decades, Finemet soft magnetic alloy has attracted considerable technological attention due to the excellent soft magnetic properties such as high magnetization, low coercivity and high Curie temperature. These properties originate from the particular microstructure characterized by randomly distributed Fe(Si) nanograins in the amorphous matrix [1]. These alloys are generally synthesized by controlled annealing of an amorphous precursor [2,3]. The production of amorphous precursor requires very high cooling rates, thus only powders, wires and ribbons with maximum thickness of ~ 30 μm could be achieved which are not suitable in industrial applications where a large volume of magnetic materials is required [4]. Producing bulk amorphous alloys by direct casting methods such as suction casting is dependent on the glass forming ability of the alloy. In addition; the maximum size of the magnetic alloys produced by this technique is seriously restricted to ~ 10 mm [5]; while powder metallurgy techniques could produce bulk amorphous alloys in the larger sizes. The consolidation of the amorphous powders is carried out in the super-cooled liquid region between Tg (glass transition temperature) and Tx (onset of
crystallization temperature). In this region, the amorphous materials soften and the viscous flow allows the compaction of amorphous powders [6–8]. Spark plasma sintering (SPS) is a unique technique for fabrication of dense and near net shape bulk amorphous alloys where a DC on–off pulse is used to produce spark plasma for instantaneous heating of powders. The current passes through the die as well as the sample enabling the sample to be heated homogenously from both inside and outside [9]. The very high heating rates in the order of 100 °C min− 1 as well as the short holding times in the range of 3–10 min outstand SPS among other conventional vacuum hot pressing techniques in the prevention of crystallization and grain growth [5,10–14]. Despite too many studies have been done on the magnetic properties of Finemet ribbons and powders, very few investigations has been reported about consolidation of these materials. In this research, Fe73.5Si13.5B9Nb3Cu1 bulk alloy was fabricated by SPS of amorphous melt-spun powders. The microstructure and magnetic properties as well as the mechanisms responsible for the densification and the attendant structural changes during SPS were also investigated. 2. Experimental procedure
⁎ Corresponding author. E-mail addresses: [email protected], [email protected] (T. Gheiratmand).
http://dx.doi.org/10.1016/j.powtec.2015.11.060 0032-5910/© 2015 Elsevier B.V. All rights reserved.
Finemet powder for consolidation was produced by mechanical milling of melt-spun amorphous ribbons. The detail of processing
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procedure has been given in our previous publications [3,15–17]. Milled powders with different milling times of 36 and 45 min were sieved to obtain different particle size ranges. Particles smaller than 125 μm were used for consolidation. The milled powder for 36 min (SPS1 and SPS3) was fully amorphous while that milled for 45 min (SPS2) was partially crystallized during milling. The degassing of powders was conducted in a vacuum furnace up to 10−5 Torr at 310 °C for 50 min. The amorphous powders were consolidated using SPS equipment (model: Easy fashion) at temperature between Tg (glass transition temperature) and Tx (onset of crystallization). The powders were charged in a graphite die with 13 mm inner diameter lined by graphite paper. It was then inserted in a vacuum chamber evacuated up to 6 × 10−2 Torr at room temperature. The temperature rose up to 350 °C, then the powders were pressed under the pressure of 40 MPa. Afterwards, the temperature of samples (SPS1 and SPS2) increased to the consolidation temperature by the rate of 30 k min−1 which took 7 min. In order to prolong the consolidation time, the temperature of another sample (SPS3) was increased from 350 °C to 560 °C by the rate of 10 k min−1. The duration of consolidation in this sample was taken as 21 min. The amperage and voltage for consolidation were selected in the range of 1500–2500 A and 2–4 V, respectively. Consolidated alloys were sanded and polished in order to remove graphite paper and carburized layer from the surface. The size distribution of particles was determined by a laser diffraction particle size analyzer (PSA, NanoTecANALYSETTE22 FRITSCH). The particle size at 50% of cumulative curve, D50, was found out. In addition, D10 and D90 were reported. The density of bulk samples was measured by the Archimedean method using ethanol. The microstructure of both degassed powders and consolidated samples was investigated by X-ray diffractometer, XRD (STOE D-64,295) using Co-Kα radiation with a step size of 0.02°. Thermal stability of milled and compacted samples was conducted by DSC analysis performed on a METTLER TOLEDO TGA/DSC1 device at heating rate of 20 °C min−1 under N2 atmosphere. Tg and Tx were measured as 465 °C and 565 °C, respectively. The microstructure of sintered samples was characterized using a Scanning Electron Microscope, SEM (VEGA-TESCAN) and an optical microscope. Transmission electron microscopy observations were carried out on TEM (Tecnai G2 20) operating at 200 kV. The TEM sample was thinned by ion milling. Magnetic properties were determined at Room temperature using a superconducting quantum interference device (SQUID, Quantum Design MPMS-5) at external magnetic field up to 20 kOe. To determine the coercivity in the unit of Am− 1 Helmholtz coil device was applied.
Fig. 1. Cumulative passing volume percent of amorphous powder and its particle size distribution.
3. Result and discussion Fig. 1 shows the cumulative passing volume percent of amorphous powder together with the particle size distribution. The particles have a bimodal distribution with a D50 of 7 μm. Moreover, D10 and D90 have been measured as 2 and 17 μm, respectively. This implies that the size of 10% of particles is below 2 μm and 10% is larger than 17 μm. The DSC curves of consolidated samples together with those of milled ribbon for 45 min and melt-spun ribbon are presented in Fig. 2. The melt spun ribbon data has two exothermic peaks. The first peak at the lower temperature is ascribed to the formation of the Fe(Si) phase while the second one corresponds to the formation of Fe23B6/Fe3B boride phases. In the plot of ribbon milled for 45 min, three exothermic peaks could be observed at 560, 712 and 900 °C which are attributed to the formation of Fe(Si), Fe3B and Fe2B phases, respectively. The area under the first peak of sintered samples decreased indicating the formation of Fe(Si) phase during sintering. Actually, the crystallization enthalpy of Fe(Si) phase decreases from 60.06 Jg−1 in the melt-spun ribbon to 9.21 and 4.68 Jg−1 for SPS1 and SPS2 samples, respectively. The Fe(Si) crystallization peak of SPS3 sample is so broad that its enthalpy could not be measured. This implies that the high fraction of Fe(Si) phase has been formed during sintering. In addition, Fe3B/Fe23B6 crystallization enthalpies are 38.28, 38.21, 13.02 and 38.27 Jg−1 for melt-spun ribbon, SPS1, SPS2 and SPS3 samples, respectively. Since the structure of SPS1 and SPS3 samples before sintering was similar to the melt-spun ribbon, their crystallization enthalpies are almost identical. The decomposition of Fe23B6 phase into the mixture of Fe3B and Fe2B phases occurred during milling for 45 min, which causes the crystallization enthalpy of SPS2 sample to become lower. The fraction of crystalline phase formed in the sintered samples was calculated using the heat change of crystallization peaks in DSC curves. These values are 84 wt.% and 90 wt.% for SPS1 and SPS2 samples, respectively. Fig. 3 shows the XRD pattern of the sintered samples by SPS as well as degassed powders. As could be seen, the patterns of both degassed samples show typical broad peak, characteristic of amorphous materials, indicating that no significant phase change occurs during degassing. The structure of milled sample for 45 min was partially crystalline before degassing, i.e. small amounts of Fe(Si) nanocrystals together with Fe2B and Fe3B phases developed in the structure during milling. The structure of all three sintered samples shows the formation of Fe(Si) phase. The higher intensity of Fe(Si) characteristic peaks in the SPS3 than SPS1 sample confirms the formation of more Fe(Si) crystalline phase. This result is the consequence of longer time of keeping SPS3 sample at high temperature of consolidation. Thus, based on the DSC and XRD analyses, it could be concluded that the fraction of Fe(Si) phase in SPS3 sample is higher than the amount of 84 wt.% obtained
Fig. 2. DSC curves of as milled ribbons and spark plasma sintered samples.
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Fig. 3. XRD patterns of degassed powders and spark plasma sintered samples.
in SPS1 sample. It also seems that the amount of boride phases does not change significantly during sintering. The Fe(Si) crystallite size in SPS1 sample was calculated by Scherrer's relationship after eliminating the instrumental errors as ~8 nm. The density of bulk samples was determined using Archimedean method which is summarized in Table 1. The densities were an average of three measurements. It is obvious that relative density of fully amorphous samples is higher than the partially crystallized one. Despite the high density of SPS3 sample obtained due to the longer time of consolidation, the high fraction of Fe(Si) crystalline phase formed in this sample could be detrimental for soft magnetic properties. This will be more discussed later where the magnetic properties are given. The microstructure of bulk samples, characterized by SEM and optical microscope, are illustrated in Fig. 4. In the SPS1 and SPS3 samples, which were initially fully amorphous, the density of porosities was measured by image analyzer software as ~3 and 2.5 vol.%, respectively. In the microstructure of partially crystalline SPS2 sample more percentage of porosities could be observed. The detachment of some fine particles during polishing, marked by arrowed lines, is indicative of weak bonding between them. Achieving a more densified bulk alloy in the SPS3 sample could be attributed to the viscous flow of the amorphous powder above glass transition temperature under high pressures. Indeed, heating in the supercooled liquid region causes the homogenous flow of the amorphous material which results in filling the gaps between particles. Since the SPS2 sample was not fully amorphous, the viscosity of the amorphous portion in the glass transition temperature range was not sufficiently low to be completely densified by viscous flow mechanism. It has been reported that the viscosity of amorphous alloy decreases by hundredfold above Tg in the supercooled liquid region [5,18]. The wide distribution of powder particles could also be distinguished from optical and SEM images. It seems that the smaller particles fulfill the porosities between larger particles causing the density to increase. Thus, as microstructural features illustrated, the particle size distribution has significant influence on the sintering behavior of Finemet amorphous powders. Moreover, using PSA data, D50, D10 and D90 have been measured as 7, 2 and 17 μm, respectively. Therefore, several larger
Table 1 Density of the consolidated powders. Sample
Density (g/cm3)
Relative density
SPS 1 SPS 2 SPS 3
7.01 6.77 7.08
96.02 92.73 96.98
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particles observed on the SEM and optical microscope images could also be the result of connecting the smaller particles by sintering. The other important parameter affecting on magnetic properties is the location of porosities. In general, for high permeability, highfrequency and low-loss applications, the most dense structure with the uniform grain size is desirable. Any porosity that exists should be inter-granular so that the domain walls cross them without any difficulty. The permeability of a soft magnetic material is mainly determined by domain wall displacement as well as the spin rotational magnetization. The spin contribution is specified by sintering density while domain wall displacement is mostly determined by grain size. Actually, porosities are like a second paramagnetic phase which reduce the saturation magnetization but increase the electrical resistivity of soft magnetic material. At high frequencies, porosities act as an insulating barrier for particle to particle eddy current paths [19–21]. In nanocrystalline soft magnetic alloys the average crystallite size is smaller than the single domain size (35 nm for Fe base alloys). Hence, each domain comprises from several crystallites [1]. The sintering is performed by connecting of particles which contain too many magnetic domains. Existence of fewer pores within the particles and their aggregation along the boundaries means that there are fewer obstacles for pinning the domain walls during magnetization cycle. As could be seen in Fig. 4(b) and (f) the porosities are mostly aggregate between powder particles marked by circles. Therefore, it could be concluded that the amount of pores should be minimized to keep the magnetization at the highest level and demagnetizing field at the lowest level. The structure of SPS1 sample has also been investigated by TEM as depicted in Fig. 5. As could be observed, the microstructure of this sample consists of Fe(Si) crystalline phase embedded in the amorphous matrix. Mean crystallite size is about 9 nm which is in good agreement with XRD results. According to diffraction pattern, the structure of nano crystallites is BCC. It could be obviously seen that there are amorphous regions around nano crystallites which have a different composition from primary amorphous matrix as a result of chemical partitioning. The densification of amorphous powders during SPS may be sometimes the result of surface melting of powder particles. This densification mechanism is associated with devitrification of the amorphous powder and formation of both Fe(Si) and boride phases. Since no crystalline peak corresponding to the boride phases was detected in the XRD pattern after degassing, temperature never reached 600 °C. The magnetization and coercivity of bulk samples are listed in Table 2. The higher magnetization of SPS3 sample could be related to its higher density and consequently, the lower magnetostatic (shape) anisotropy. The high amount of Fe(Si) crystalline phase formed during consolidation in this sample should be also regarded in increasing the magnetization. The surface poles, which are distributed on the surface of porosities, produce internal fields in addition to the external fields. The internal field known as the demagnetizing field (Hd) is proportional to the magnetization of the body with direction opposite to it. Hd is sensitive to the shape of sample. Thus, the high density of porosities in a bulk sample leads to the higher demagnetizing field and the lower total magnetization. The coercivity of sintered samples was measured as 123, 189 and 147 Am−1for SPS1, SPS2 and SPS3 samples, respectively. Porosities act as paramagnetic phases which lock the domain walls and prevent their easy rotation during magnetization cycle. Besides, an enhancement of coercivity occurs. Moreover, the demagnetization fields produced by porosities are effective in increasing the coercivity. Furthermore, the high volume fraction of Fe(Si) crystalline phase could increase the magnetostriction and thus the coercivity. It should be noted that about 75–80% [22] of Fe(Si) crystalline phase embedded in amorphous matrix is needed to have the total near zero magnetostriction. It should be said that although the value of coercivity in bulk sample is higher than that reported for Finemet ribbons (1–16 Am−1 [23–26]), it is much smaller than that of toroidally winded ribbons which are not suitable for industrial applications where a large volume of magnetic materials with a complex shape is required.
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Fig. 4. Optical micrographs and SEM images of spark plasma sintered samples consolidated from powders milled for a) and b) 36 min and sintered for 7 min, c) and d) 45 min, sintered for 7 min, e) and f) 36 min and sintered for 21 min.
where B is a term that collects materials and geometric constants, R is the particles radius, n is the representative of dominant densification mechanism and m relates to the particles size. These parameters for different densification mechanisms are listed in Table 3 [27]. According to Table 3, if the shrinkage and time be linearly related to each other (as Eq. (2)) the dominant densification mechanism will be viscous flow [28].
The values of microhardness of bulk samples are given in Table 2 as well. It could be seen that SPS3 sample has the highest value of microhardness which can be explained on the base of the high density and high volume fraction of crystalline phase formed during consolidation. In addition, the microhardness of SPS2 sample is slightly higher than that of SPS1. This result has been obtained despite the improved density of SPS1 sample and could be explained based on the higher fraction of the crystalline phase formed during sintering in the sample sintered from the partial crystalline powder. It has been shown that the hardness of the sample consisting of the crystalline phase is much higher than that composed of fully amorphous phase [3]. In order to perceive the densification mechanism of Finemet powders by SPS, the shrinkage of SPS1 sample with respect to sintering time has been plotted in Fig. 6. It was proposed that relationship between the shrinkage (ΔL/L0) and sintering time (t) is stated as follow [27]:
η ¼ η0 expðQ v =RT Þ
n=2 ΔL B ¼ n mt L0 2 R
As stated earlier, in order to sinter the powders the temperature increased up to 350 °C before applying the pressure and then it rose up by the heating rate of 30 k min−1 during applying the pressure up to the
ð1Þ
ΔL 3γ t ¼− L0 4Dη
ð2Þ
Here γ is the surface energy, D is the particle size and η is the viscosity of amorphous materials. In addition, the viscosity of amorphous materials is temperature dependent and could be stated by Arrhenius equation as follow [28]: ð3Þ
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seems that at the first stage of densification, where the temperature of powder is below the glass transition temperature and not high enough for viscous flow, the rearrangement of particles is the main mechanism of densification. Since the size of milled amorphous powder particles was selected below 125 μm, the smaller size particles get redistributed at some of the interparticle spaces between larger amorphous particles. However, a significant number of smaller interparticle spaces still remain empty. The full densification then accomplished by more rising the temperature as well as applying the pressure above Tg, where the viscosity of the amorphous powder becomes smaller and the viscous flow of the amorphous particles is the main reason for increasing the sample density. As could be seen in Fig. 6(b), the relationship between the shrinkage and time at the second stage of sintering could be distinguished almost linear. Furthermore, the slope of the straight line 0Þ LnðdðΔL=L Þ versus 1/(623 + 0.5t) allows the calculation of the activation dt energy for viscous flow as 22.9 kJ mol−1. This value is in good agreement with the literature reports [29–32].
4. Conclusion Fig. 5. TEM bright field image and SAED pattern of spark plasma sintered consolidated from powder milled for 36 min.
Table 2 Saturation magnetization, coercivity and hardness of different sintered samples. Sample
Magnetization (emug−1) ± 1.0
Coercivity (Am−1) ± 1
Hardness (VHN) ± 10
SPS1 SPS2 SPS3
122.29 118.85 124.01
123 189 147
1097 1110 1118
maximum value of 560 °C. Thus, the relationship between time and temperature can be declared as T = 623 + 0.5 t. Thus, Eq. (2) could be written as follow: dðΔL=L0 Þ −3γ Qv ¼ Ln − Ln dt 8Dη0 Rð0:5t þ 623Þ
ð4Þ
Fig. 6(a) presents the shrinkage of Finemet amorphous powder during SPS. The densification of amorphous powder begins immediately upon applying the pressure by movement of punch. As could be seen in Fig. 6(a), the plot of shrinkage versus time is not a single straight line indicating the act of different mechanisms during sintering. It
This work sought to examine the possibility of production of Finemet magnetic alloy from amorphous powders by SPS method. The properties of produced Finemet bulk sample has been also investigated, which according to the authors' best knowledge has not been discussed elsewhere. Finemet bulk alloy with nearly ~ 97% relative density was successfully fabricated from amorphous powders by SPS in the supercooled liquid region for 21 min. The density of the sample consolidated from fully amorphous powders was higher than that obtained from partially crystallized one due to the lower viscosity and consequently, more homogenous flow of the fully amorphous material in the glass transition range. The wide distribution of powder particles was observed in the microstructure of sintered samples which is the result of particle size distribution and also smaller particles connection during sintering. The aggregation of porosities along the particle boundaries increases the resistivity and provides an insulating barrier for particle to particle eddy current paths at high frequencies. The microstructure in the bulk sample consisted of Fe(Si) nanocrystals with average size of 9 nm and a BCC structure embedded in the amorphous matrix. The fraction of crystalline phase was estimated ~ 84 wt.% from DSC measurements. The magnetization and coercivity in the sintered sample was measured as 122.9 emu g−1 and 123 Am−1, respectively. By applying the pressure, two distinct consolidation mechanisms were distinguished during the SPS process based on the relationship between shrinkage (ΔL/L0) and sintering time (t). At initial period of pressing up to 90 s, the densification was due to the particles rearrangement.
Fig. 6. (a) The variation of shrinkage with respect to the time, (b) plot of Ln(d(ΔL/L0)/dt) versus 1/(0.5 t + 623) corresponding to the second stage of densification by viscous flow.
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Table 3 The values of n, m and B for different densification mechanisms [27]. Mechanism
n
m
B
Viscous flow Plastic deformation Volumetric diffusion Grain boundary diffusion Surface diffusion
2 2 5 6 7
1 1 3 4 4
3γ/2η 9πγbDv/RT 80DvγΩ/RT 20δDBγΩ/RT 56DsγΩ4/3/RT
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