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Thermal stability, crystallization and soft magnetic properties of Fe-P-C-based glassy alloys
Journal of Non-Crystalline Solids 454 (2016) 39–45
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Thermal stability, crystallization and soft magnetic properties of Fe-P-C-based glassy alloys Jianfeng Wang a,⁎, Yaxin Di a, Zhe Fang a, Shaokang Guan a, Tao Zhang b a b
School of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, PR China School of Materials Science and Engineering, Beihang University, Beijing 100191, PR China
a r t i c l e
i n f o
Article history: Received 24 July 2016 Received in revised form 11 September 2016 Accepted 16 October 2016 Available online xxxx Keywords: Iron alloys Metallic glasses Crystallization Magnetic properties
a b s t r a c t Thermal stability, crystallization and magnetic properties of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys were investigated. The relationships of glass-forming ability (GFA) with thermal properties and crystallization, magnetic properties with crystallization were also discussed. The variation trend of GFA for these glassy alloys is in agreement with the reduced glass transition temperature. The annealing experiment shows that the primary precipitation phase both for the Fe80P11C9 and Fe80P9B2C9 glassy alloys is α-Fe, and the addition of small amounts of Al is effective for suppression of the formation of α-Fe phase, resulting in the extension of supercooled liquid region and significant enhancement of GFA. The activation energy for primary crystallization calculated by Kissinger's equation indicates the addition of B in Fe–P–C alloy or Al in Fe–P–B–C alloy makes nucleation more difficult, which is responsible for the enhancement of GFA. In addition, it was found that the crystallized Fe80P9B2C9 specimen with ultrafine α-Fe grains exhibits high saturation magnetic polarization of up to 1.60 T and low coercive force of about 3.0 A/m. © 2016 Elsevier B.V. All rights reserved.
1. Introduction Fe-based amorphous alloys and the corresponding nanocrystalline alloys are of great value for commercial applications due to their excellent soft magnetic properties including high saturation magnetic polarization (Js), low coercive force (Hc), and high permeability (μ) [1,2]. In recent decades, Fe-based ferromagnetic amorphous foils have been widely used as key parts in magnetic devices, such as high-efficiency transformers and sensors. In order to further extend their application fields as magnetic materials, many efforts have been devoted to developing new Fe-based ferromagnetic glassy alloys with better glass-forming ability (GFA). For the last decades, a number of Fe-based ferromagnetic bulk metallic glasses (BMGs) have been synthesized, which can be broadly categorized into three groups. One is Fe–P–C-based BMG group [3–14], another is Fe–B–Si-based BMG group [15–20], and the other is Fe–Bbased BMG group [21–26]. Among these types, Fe–P–C-based BMGs have attracted increasing interests for their promising mechanical and magnetic properties, such as (Fe,Mo,Ni,Cr)80P12.5C5B2.5 BMGs with notch toughness values of 44.2–53.1 MPa·m1/2 [12], Fe74-xMoxP13C7 (x = 3 and 6) BMGs with plastic strain (εp) of above 5.0% [14], and Fe81Mo1P7.5C5.5B2Si3 and Fe82Mo1P6.5C5.5B2Si3 BMGs with saturation magnetization of about 1.6 T [13,14]. It is especially noted that some of these BMGs, such as Fe76Mo2P10C7.5B2.5Si2 [10], Fe78Mo1P9C6.5B3.5Si2 [13], and Fe77Mo3P13C7 [14] were reported to exhibit a unique combination of ⁎ Corresponding author. E-mail address: [email protected] (J. Wang).
http://dx.doi.org/10.1016/j.jnoncrysol.2016.10.014 0022-3093/© 2016 Elsevier B.V. All rights reserved.
high Js and appreciable εp. Such alloys could be more viable for use directly in bulk form as magnetic sensors, valves, and clutches. However, these alloys contain relatively expensive metallic element Mo, which leads to a significant increase in materials cost. Recently, considering that the raw material cost of Fe-based glassy alloys is very important for their practical applications, we have developed a Mo-free Fe80P11C9 BMG with a critical diameter of 1.5 mm by copper mold casting in ternary Fe–P–C alloy system [27]. The alloy with a lower materials cost (compared with Mo-bearing Fe–P–C-based BMGs) exhibits good overall properties, such as high Js of up to 1.37 T and significant εp of about 1.4%. Although Li et al. [28] also reported that a ternary Fe80P13C7 BMG with Js of 1.53 T and εp of about 1.1% can be synthesized by the combination of fluxing treatment and J-quenching technique at the same time, the preparation process of BMGs is more complicated. Subsequently, we designed new Fe80P9B2C9 [29] and Fe77Al3P9B2C9 [30] BMGs with a better combination of GFA and properties by adding B and Al into Fe–P–C and Fe–(P,B)–C alloys, respectively. In this paper, we intend to further investigate the substitution effects of B and Al elements on the thermal properties, crystallization and soft magnetic properties of Fe–P–C-based glassy alloys. In addition, the relationships of GFA with thermal properties and crystallization, soft magnetic properties with crystallization were also discussed. 2. Experimental Mother alloy ingots with nominal compositions of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 were prepared by induction melting the mixture of pure Fe (99.9 mass%), Al (99.9 mass%), C (99.99 mass%),
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J. Wang et al. / Journal of Non-Crystalline Solids 454 (2016) 39–45
and pre-alloyed Fe–P (22.41 mass% P) and Fe–B (17.24 mass% B) in a purified argon atmosphere. The weight loss due to melting was found to be less than 0.15% of the starting materials. Alloy ribbons with a dimension of about 20 μm in thickness and about 1.0 mm in width were prepared by single-roller melt spinning. Alloy rods with different diameters of 1.0–3.0 mm and a length of a couple of centimeters were prepared by copper mold casting. The glassy and crystallized structures were identified by X-ray diffraction (XRD) with Cu Kα radiation. Glass transition, crystallization, and melting behaviors were evaluated using differential scanning calorimetry (DSC) at a heating rate of 20 K/min. To reduce the influence of undercooling, the solidification behavior was investigated using DSC at a very low cooling rate of 5 K/min. The crystallization kinetics of each as-spun specimen was evaluated using DSC at different heating rates of 10, 20, 30, and 40 K/min. The annealing experiment was carried out by keeping the as-spun amorphous specimens in the tubular furnace preheated to annealing temperatures for 600 s under vacuum atmosphere followed by water quenching. The density was measured by Archimedes' method. The saturation magnetic polarization (Js) and coercive force (Hc) were measured at room temperature with a vibrating sample magnetometer (VSM) under an applied field of 150 kA/m. For alloys with a smaller Hc than 100 A/m, Hc was measured by a DC B-H loop tracer under a field of 400 A/m. 3. Results 3.1. Thermal stability Fig. 1(a)–(c) display DSC curves of the melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys at different heating rates from 10 to 40 K/min showing the glass transition and crystallization behaviors. Upon heating, each of the curves exhibits an obvious endothermic event characteristic of the glass transition and a supercooled liquid region, followed by exothermic reactions corresponding to the crystallization of the undercooled liquid. From Fig. 1(a) and (b), the Fe80P11C9 and Fe80P9B2C9 glassy alloys exhibit a weak exothermic shoulder peak before the main exothermic peak which is caused due to the overlapping of two crystallization peaks. The substitution of 3 at.% Fe by Al in Fe80P9B2C9 alloy leads to disappearing of the shoulder peak as shown in Fig. 1(c). The values of the glass transition temperature (Tg) and onset crystallization temperature (Tx) determined from the DSC traces of 20 K/min heating rate are summarized in Table 1. Both Tg and Tx increase with substitution of 2 at.% P by B element in Fe80P11C9 alloy as well as with further substitution of 3 at.% Fe by Al in Fe80P9B2C9 alloy. The melting and solidification behaviors of these alloys are shown in Fig. 2. The near invariance of Tm for these compositions indicates that they may have the same eutectic reaction. Upon cooling, the Fe80P11C9 alloy shows two exothermic peaks. The first peak at the high temperature corresponds to the solidification of the primary phases, while the second one at the relatively low temperature is ascribed to that of the eutectic phase. When compared with Fe80P11C9 alloy, the first peak for Fe80P9B2C9 alloy is shifted to a lower temperature, which leads to the overlapping of two peaks in the DSC curve. Consequently, Tl is also lowered from 1262 to 1240 K as shown in Table 1. The further decrease of Tl is caused by the addition of Al element. Therefore, it is suggested that Fe77Al3P9B2C9 alloy with the lowest Tl is the closest to the eutectic point among these three alloys. 3.2. Crystallization behavior Fig. 3(a)–(c) shows XRD patterns of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys subjected to annealing for 600 s at different temperatures. The XRD patterns of as-spun specimens are also shown for comparison. Obviously, after annealing at (Tg − 20) K, these alloys maintain glassy structure. As we have known, the isothermal heat treatment in the supercooled liquid region can induce phase separation,
Fig. 1. DSC curves of (a) Fe80P11C9, (b) Fe80P9B2C9, and (c) Fe77Al3P9B2C9 glassy alloys at different heating rates from 10 to 40 K/min, showing glass transition and crystallization behaviors. Arrows indicate the glass transition temperature (Tg) and onset crystallization temperature (Tx), respectively.
nucleation, and growth of nuclei. For Fe80P11C9 alloy, as shown in Fig. 3(a), α-Fe phase primary precipitates from the glassy matrix after annealed at 695 K and 700 K in the supercooled liquid region. The result
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Table 1 Critical diameters (dc), thermal parameters, and activation energy of crystallization (Ec) for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys. Tg, Tx, and Tl are the glass transition temperature, onset crystallization temperature, and liquidus temperature, respectively. ΔTx is equal to Tx − Tg, and Trg is equal to Tg / Tl. Alloys
dc (mm)
Tg (K)
Tx (K)
Tl (K)
ΔTx (K)
Trg
Ec (kJ·mol−1)
Fe80P11C9 Fe80P9B2C9 Fe77Al3P9B2C9
1.5 1.8 3.0
690 698 712
720 727 751
1262 1240 1234
30 29 39
0.547 0.563 0.577
486 ± 31 496 ± 32 537 ± 18
shows that the exothermic shoulder peak in the Fe80P11C9 alloy is attributed to the primary crystallization of α-Fe phase. Here, it should be pointed out that, since the overlapping of crystallization peaks, the annealing temperature range to obtain α-Fe phase is very narrow. When annealing temperature (Ta) goes up to 760 K beyond the main exothermic peak, the pattern shows Bragg peaks of α-Fe, Fe3P and Fe3C phases superimposed on the broad diffuse scattering peak from the amorphous phase. For the specimen annealed at 880 K corresponding to temperature higher than the second crystallization peak, the crystallization phases remain unchanged while the intensities of the crystalline components increase, indicating the fraction of crystallization phases increases. This infers that the second peak observed by DSC is most likely due to crystallization of the residual amorphous phase. Similar crystallization process was observed for Fe80P9B2C9 alloy as well, as shown in Fig. 3(b). That is, the primary α-Fe phase is precipitated at 703 K and 708 K, while Fe3P and Fe3C phases are formed when Ta is raised to 780 K above the main exothermic peak. For Fe77Al3P9B2C9 alloy, formation of α-Fe phase is suppressed, as evidenced by the disappearing of the shoulder peak. Fig. 3(c) confirms that the diffraction peak corresponding to α-Fe does not appear in the XRD pattern obtained after annealed at 720 K corresponding to the temperature just below the first exothermic peak, suggesting that the supercooled liquid of the Fe77Al3P9B2C9 alloy crystallizes in different path by a simultaneous formation of Fe3P, Fe3C and Fe23(C,B)6. The same phases can be found when Ta is up to 800 K above the first exothermic reaction. The XRD pattern of the specimen annealed at 930 K corresponding to the temperature above the second peak are identified as the mixed phases of Fe3P, Fe3C, α-Fe and Fe23(C,B)6. 3.3. Magnetic properties The magnetic hysteresis loops of the melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloy ribbons are shown in Fig. 4. All the
Fig. 3. XRD patterns of (a) Fe80P11C9, (b) Fe80P9B2C9, and (c) Fe77Al3P9B2C9 glassy alloys subjected to annealing for 600 s at different temperatures. The XRD patterns of as-spun specimens are also shown for comparison.
Fig. 2. DSC curves of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys, showing melting and solidification behaviors. Arrows indicate the melting temperature (Tm) and liquidus temperature (Tl), respectively.
alloy specimens exhibit typical soft magnetic hysteresis curves. The Js increases from 1.37 to 1.46 T as substituting 2 at.% P by B, but decreases from 1.46 to 1.35 T as substituting 3 at.% Fe by Al. Here, it should be pointed out that Js for Fe80P9B2C9 amorphous alloy in this study is slightly higher than that reported in Ref. 29. This difference may be caused by different purities of the raw materials used in the two experiments. The Hc values for melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 ribbons are 16.4, 8.8, and 12.1 A/m, respectively. Fig. 5(a)–(c) show magnetic hysteresis loops of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 specimens annealed at different temperatures. For comparison, the data of as-spun specimens are also plotted. The values of Js for the glassy and annealed specimens are listed in Table 2. As can be seen, the change of Js for all alloys is very small with increasing Ta to the temperatures lower than Tg. For the Fe80P11C9 and
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Fig. 4. Magnetic hysteresis loops of the as-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloy ribbons.
Fe80P9B2C9 specimens annealed at 695 K and 703 K, respectively, the Js values increase to 1.49 T and 1.60 T. Such a high Js for Fe80P9B2C9 alloy is rarely observed in the Fe-based bulk glassy materials. The appearance of a single precipitate of α-Fe phase in glassy matrix could be responsible for the high Js [31,32]. With further increasing Ta, Js values of the Alfree alloys decrease. It should be pointed out that, for the Fe77Al3P9B2C9 alloy, annealing at 720 K causes a slight decrease in Js and a dramatic increase in Hc as shown in Fig. 5(c). The result indicates that Js of this alloy could not be effectively improved by annealing experiment. The Hc values for the as-spun and annealed specimens are also listed in Table 2. The changes in Hc for Fe80P11C9 and Fe80P9B2C9 alloys as a function of Ta are shown in Fig. 6. There is no remarkable different for the trend of the curves between these two alloys. Compared with the as-spun specimens, the specimens annealed at (Tg − 20) K exhibit a lower Hc value of about 3.0 A/m, which can be attributed to the stress releasing. The further increase in Ta to 695 K for Fe80P11C9 and 703 K for Fe80P9B2C9 causes a slight increase in Hc to 3.8 A/m and 3.6 A/m. After that, Hc increases sharply to 880 A/m and 780 A/m as increasing Ta to 700 K and 708 K for Fe80P11C9 and Fe80P9B2C9, respectively. Consequently, we can speculate that enhanced Js as well as good softness for the Fe80P11C9 and Fe80P9B2C9 amorphous alloys can be obtained by annealing the specimens at 698 K and 703 K, respectively. 4. Discussion 4.1. Relations of GFA with thermal stability and crystallization behavior Previously, we have reported that the critical diameters of glassy alloy rods (dc) are up to 1.5, 1.8, and 3.0 mm for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloys [27,29,30], respectively. Although the purity of Fe used in this experiment for Fe80P9B2C9 is a little different from that reported in Ref. [29], glassy rods of this composition with a diameter of 1.8 mm were also synthesized (not shown here). Obviously, the GFA of Fe80P11C9 glassy alloy can be slightly enhanced by partial substituting 2 at.% B for P. Further enhancement in the size of glassy rods is observed when 3 at.% Al is substituted for Fe. The reason for the improvement of GFA in the present Fe–P–C-based alloy system can be explained from the thermodynamic point of view. As shown in Fig. 2 and Table 1, addition of the proper amount of B to ternary Fe80P11C9 and Al to quaternary Fe80P9B2C9 can stabilize the supercooled melts by lowering the liquidus temperature. This means that the undercooled melts can be more easily frozen to be glassy against the formation of the competing crystalline phases. As a result, the reduced glass transition temperature (Trg = Tg / Tl), which has been regarded as one of the important parameters for evaluation of glass formation [33], increases with the addition of B in Fe–P–C alloy or Al in Fe–P–B–
Fig. 5. Magnetic hysteresis loops of (a) Fe80P11C9, (b) Fe80P9B2C9, and Fe77Al3P9B2C9 specimens annealed at different temperatures. For comparison, the data of as-spun specimens are also plotted. The inset shows the enlarged partial hysteresis curves of the specimens.
C alloy. The result shows that Trg effectively reflects GFA in Fe–P–Cbased glassy alloys. The variation trend of GFA for Al-bearing Fe-based glassy alloy can be related with the crystallization behavior of glassy alloys. When compared with Fe80P9B2C9 glassy alloy, the crystallization of Fe77Al3P9B2C9 alloy upon continuous heating takes place by the simultaneous precipitation of three phases α-Fe, Fe3P and Fe23(C,B)6. It is considered that the formation of α-Fe phase upon devitrification is drastically impeded by alloying with small amount of Al, because the mixing enthalpies with a large negative value between Al and Fe could increase the difficulty of the formation of the α-Fe phase. The competitive formation of
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Table 2 Saturation magnetic polarization (Js), coercive force (Hc), and structure for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloy specimens annealed at different conditions. HT denotes heat treatment. Compositions
Annealing conditions
Structures
Js (T)
Hc (A/m)
Fe80P11C9
As-spun HT at 670 HT at 695 HT at 700 HT at 760 As-spun HT at 678 HT at 703 HT at 708 HT at 780 As-spun HT at 683 HT at 720
Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am.
1.36 1.37 1.49 1.44 1.34 1.46 1.46 1.60 1.56 1.43 1.37 1.37 1.30
16.4 3.3 3.8 2200 15,000 8.8 3.1 3.6 1800 12,700 12.1 2.0 11,300
Fe80P9C9B2
Fe77Al3P9B2C9
K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s
these crystallization phases results in the increase of stability of the supercooled liquid against crystallization. Meanwhile, the formation of Cr23B6-type Fe23(C,B)6 phase with complex face-centered cubic structure which requires long-range rearrangements of the constitute atoms also contributes to the high stability of the supercooled liquid [17,34,35]. These two effects cause the remarkable extension of the supercooled region (ΔTx = Tx − Tg) from 29 K to 39 K as shown in Table 1. The larger ΔTx revealing a higher resistance against crystallization could also be regarded as an indicator of higher GFA of alloys [36]. Therefore, we can conclude that higher GFA is obtained for the 3 at.% Al-bearing glassy alloy. However, the crystallization process cannot be changed by minor adding B in Fe–P–C alloy, which is responsible for the little change in ΔTx. At this point, it is suggested that the crystallization behavior and ΔTx cannot solely illustrate the GFA for the present Alfree glassy alloys. In addition, we further investigated the crystallization kinetics of amorphous alloys. The apparent activation energy (Ea) can be evaluated by the Kissinger equation [37]: ln β=T 2 ¼ −Ea =ðRT Þ þ C;
ð1Þ
where β is heating rate, T is the characteristic temperature and C is a constant. As observed from Fig. 1, Tx marked by arrows in the DSC traces are sensitive to the heating rate, and they shift to higher temperatures with increasing heating rate. This means that the crystallization have remarkable kinetic effects. By plotting ln(β/T2x) versus 1000/Tx, good linear fits can be obtained as shown in Fig. 7. From the slopes of these straight lines, the activation energy for primary crystallization (Ec) is determined as 486, 496, and 537 kJ/mol for Fe80P11C9, Fe80P9B2C9, and
Fig. 6. Changes in coercive force (Hc) for Fe80P11C9 and Fe80P9B2C9 alloys as a function of annealing temperature (Ta).
+ α-Fe (3–5) nm + α-Fe (~61 nm) + Fe3P + Fe3C
+ α-Fe (3–5) nm + α-Fe (~78 nm) + Fe3P + Fe3C
+ Fe3P + Fe3C + Fe23(C,B)6
Fe77Al3P9B2C9 glassy alloys, respectively. It is well known that the onset crystallization temperature is considered to be associated with the nucleation process. Thus, the activation energy deduced from the onset crystallization temperatures corresponds to nucleation [38]. The result of crystallization kinetics indicates that the addition of B in Fe– P–C alloy or Al in Fe–P–B–C alloy makes nucleation more difficult, which is responsible for the enhancement of GFA.
4.2. Relations of magnetic properties with alloy composition and crystallization behavior As well known, the Js is an intrinsic magnetic property that is mainly determined by alloy composition and phase constitution of alloy. For Fe–metalloid-type glassy alloys, the effect of metalloid elements on saturation magnetization has been extensively discussed. According to charge-transfer model [39], the valence electrons of metalloid elements transfer to the band of Fe which contributes to magnetic moment. When metallic elements are doped into a base alloy, the valence electronic structures and magnetic moment alignments of the added elements, as well as their electronic interactions with Fe, markedly affect the magnetic moments. Generally, the addition of metalloid elements would reduce the Js of amorphous alloys, and the metalloid element with more sp electrons results in the more decrease in Js. The number of sp electrons for the metalloid elements P, C, and B is 5, 4, and 3, respectively. Therefore, the Fe80P9B2C9 alloy with a lower proportion of P exhibit a higher Js compared with the Fe80P11C9 alloy. Further, the introduction of non-ferromagnetic Al element to Fe80P9B2C9 alloy reduces the number of magnetic atoms, and thus leads to a decrease in Js [40,41].
Fig. 7. Kissinger plots of ln(β/T2x) versus 1000/Tx for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys.
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Herzer [42] proposed a simple two-phase model to quantitatively predict the change of Js in annealed glassy alloy specimens. According to this model, the saturation magnetic polarization reflects the ratio of the volume fraction of the crystalline phase Vc/V to that of the amorphous phase Va/V with V = Vc + Va, and is expressed as Js = JscVc/ V + JsaVa/V, where Jsc and Jsa are the saturation magnetic polarization in the crystalline and amorphous phases, respectively. The increase in Js for the annealed Fe80P11C9 and Fe80P9B2C9 specimens with a single precipitation of α-Fe phase can be understood in terms of the contribution of α-Fe nanocrystallites with high saturation magnetic polarization of about 2.1 T [43]. With the further increasing of Ta, the grain size and volume fraction of α-Fe phase increase which leads to less Fe concentration leaving in amorphous matrix. The enrichment of nonmagnetic atoms including P, C and B in the matrix is responsible for a possibly rapid decrease in Jsa, and so as to a slight decrease of Js. On the other hand, it is well known that the magnetic softness for the partial nanocrystalline alloys strongly depends on the grain size of α-Fe phase [44–46]. The effective anisotropy (Keff), intimately related with Hc, is highly sensitive to the grain size (D) of nanocrystalline phase. A small increase in D may dramatically increase the Keff, which results in deterioration of the magnetic softness. As shown in Fig. 3(a) and (b), for the Fe80P11C9 and Fe80P9B2C9 specimens obtained after annealing at temperatures of 695 K and 703 K, respectively, the diffraction intensity of α-Fe is weak and the volume fraction of crystalline phase is small, so the grain size is impossible to estimate using Scherer's formula. Previous research [27] on the microstructure of crystallized Fe80P11C9 specimens has shown that ultrafine α-Fe crystallites with a diameter of about 3– 5 nm is embedded in the amorphous matrix. The SAED pattern shows only a few spots scatter in the diffraction rings, which indicates the volume fraction of α-Fe phase is small. The results are agreement with the XRD analysis. The low volume fraction of nanoscale α-Fe grains may have little influence on Keff of matrix [45,46]. Thus, the Hc of these specimens remains almost unchanged compared with the ones after relaxation. As for the Fe80P11C9 and Fe80P9B2C9 alloy specimens annealed at higher temperatures of 700 K and 708 K, respectively, obvious crystalline peak corresponding to (200) α-Fe phase is observed. The grain size estimated on the Scherer's formula from (200) plane is up to 61 nm and 78 nm, respectively. Accordingly, we could speculate that the dramatic increase in Hc is due to the increase in D. The specimens annealed at temperatures higher than the main exothermic peak exhibit a larger Hc because they compose of α-Fe phase and Fe-metalloid compounds with high magnetocrystalline anisotropy.
5. Conclusions In this work, the substitution effect of B and Al elements on the thermal properties, crystallization behavior, and magnetic properties of Fe– P–C-based amorphous alloys were investigated. The addition of the proper amount of B to ternary Fe80P11C9 and Al to quaternary Fe80P9B2C9 can stabilize the supercooled melts by lowering the liquidus temperature, which leads to the increase in Trg. The annealing experiments shows that α-Fe is the primary precipitation phase for the Fe80P11C9 and Fe80P9B2C9 alloys, while α-Fe, Fe3P and Fe23(C,B)6 are the competing crystalline phases for the Fe77Al9P9B2C9 alloy. Therefore, the suppression of the formation of α-Fe during crystallization increases the thermal stability of glass phase and thus improves the GFA significantly. The result of crystallization kinetics indicates that the addition of B in Fe80P11C9 alloy or Al in Fe80P9B2C9 alloy makes nucleation more difficult, which is in agreement with GFA. The relationship between crystallization and magnetic properties was also discussed. It is found that the resulting annealed Fe-based alloy specimen with nanoscale α-Fe grains exhibits high Js of up to 1.60 T and low Hc of about 3 A/m. However, since the optimum annealing temperature range for Al-free Fe-based amorphous alloys to obtain α-Fe phase is very narrow, further improvement in the soft magnetic properties is difficult.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant No. 51501166), Foundation of Henan Educational Committee (Grant Nos. 13A430668 and 15A430002), and Outstanding Young Talent Research Fund of Zhengzhou University (Grant No. 1421320046).
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Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/locate/jnoncrysol
Thermal stability, crystallization and soft magnetic properties of Fe-P-C-based glassy alloys Jianfeng Wang a,⁎, Yaxin Di a, Zhe Fang a, Shaokang Guan a, Tao Zhang b a b
School of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, PR China School of Materials Science and Engineering, Beihang University, Beijing 100191, PR China
a r t i c l e
i n f o
Article history: Received 24 July 2016 Received in revised form 11 September 2016 Accepted 16 October 2016 Available online xxxx Keywords: Iron alloys Metallic glasses Crystallization Magnetic properties
a b s t r a c t Thermal stability, crystallization and magnetic properties of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys were investigated. The relationships of glass-forming ability (GFA) with thermal properties and crystallization, magnetic properties with crystallization were also discussed. The variation trend of GFA for these glassy alloys is in agreement with the reduced glass transition temperature. The annealing experiment shows that the primary precipitation phase both for the Fe80P11C9 and Fe80P9B2C9 glassy alloys is α-Fe, and the addition of small amounts of Al is effective for suppression of the formation of α-Fe phase, resulting in the extension of supercooled liquid region and significant enhancement of GFA. The activation energy for primary crystallization calculated by Kissinger's equation indicates the addition of B in Fe–P–C alloy or Al in Fe–P–B–C alloy makes nucleation more difficult, which is responsible for the enhancement of GFA. In addition, it was found that the crystallized Fe80P9B2C9 specimen with ultrafine α-Fe grains exhibits high saturation magnetic polarization of up to 1.60 T and low coercive force of about 3.0 A/m. © 2016 Elsevier B.V. All rights reserved.
1. Introduction Fe-based amorphous alloys and the corresponding nanocrystalline alloys are of great value for commercial applications due to their excellent soft magnetic properties including high saturation magnetic polarization (Js), low coercive force (Hc), and high permeability (μ) [1,2]. In recent decades, Fe-based ferromagnetic amorphous foils have been widely used as key parts in magnetic devices, such as high-efficiency transformers and sensors. In order to further extend their application fields as magnetic materials, many efforts have been devoted to developing new Fe-based ferromagnetic glassy alloys with better glass-forming ability (GFA). For the last decades, a number of Fe-based ferromagnetic bulk metallic glasses (BMGs) have been synthesized, which can be broadly categorized into three groups. One is Fe–P–C-based BMG group [3–14], another is Fe–B–Si-based BMG group [15–20], and the other is Fe–Bbased BMG group [21–26]. Among these types, Fe–P–C-based BMGs have attracted increasing interests for their promising mechanical and magnetic properties, such as (Fe,Mo,Ni,Cr)80P12.5C5B2.5 BMGs with notch toughness values of 44.2–53.1 MPa·m1/2 [12], Fe74-xMoxP13C7 (x = 3 and 6) BMGs with plastic strain (εp) of above 5.0% [14], and Fe81Mo1P7.5C5.5B2Si3 and Fe82Mo1P6.5C5.5B2Si3 BMGs with saturation magnetization of about 1.6 T [13,14]. It is especially noted that some of these BMGs, such as Fe76Mo2P10C7.5B2.5Si2 [10], Fe78Mo1P9C6.5B3.5Si2 [13], and Fe77Mo3P13C7 [14] were reported to exhibit a unique combination of ⁎ Corresponding author. E-mail address: [email protected] (J. Wang).
http://dx.doi.org/10.1016/j.jnoncrysol.2016.10.014 0022-3093/© 2016 Elsevier B.V. All rights reserved.
high Js and appreciable εp. Such alloys could be more viable for use directly in bulk form as magnetic sensors, valves, and clutches. However, these alloys contain relatively expensive metallic element Mo, which leads to a significant increase in materials cost. Recently, considering that the raw material cost of Fe-based glassy alloys is very important for their practical applications, we have developed a Mo-free Fe80P11C9 BMG with a critical diameter of 1.5 mm by copper mold casting in ternary Fe–P–C alloy system [27]. The alloy with a lower materials cost (compared with Mo-bearing Fe–P–C-based BMGs) exhibits good overall properties, such as high Js of up to 1.37 T and significant εp of about 1.4%. Although Li et al. [28] also reported that a ternary Fe80P13C7 BMG with Js of 1.53 T and εp of about 1.1% can be synthesized by the combination of fluxing treatment and J-quenching technique at the same time, the preparation process of BMGs is more complicated. Subsequently, we designed new Fe80P9B2C9 [29] and Fe77Al3P9B2C9 [30] BMGs with a better combination of GFA and properties by adding B and Al into Fe–P–C and Fe–(P,B)–C alloys, respectively. In this paper, we intend to further investigate the substitution effects of B and Al elements on the thermal properties, crystallization and soft magnetic properties of Fe–P–C-based glassy alloys. In addition, the relationships of GFA with thermal properties and crystallization, soft magnetic properties with crystallization were also discussed. 2. Experimental Mother alloy ingots with nominal compositions of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 were prepared by induction melting the mixture of pure Fe (99.9 mass%), Al (99.9 mass%), C (99.99 mass%),
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J. Wang et al. / Journal of Non-Crystalline Solids 454 (2016) 39–45
and pre-alloyed Fe–P (22.41 mass% P) and Fe–B (17.24 mass% B) in a purified argon atmosphere. The weight loss due to melting was found to be less than 0.15% of the starting materials. Alloy ribbons with a dimension of about 20 μm in thickness and about 1.0 mm in width were prepared by single-roller melt spinning. Alloy rods with different diameters of 1.0–3.0 mm and a length of a couple of centimeters were prepared by copper mold casting. The glassy and crystallized structures were identified by X-ray diffraction (XRD) with Cu Kα radiation. Glass transition, crystallization, and melting behaviors were evaluated using differential scanning calorimetry (DSC) at a heating rate of 20 K/min. To reduce the influence of undercooling, the solidification behavior was investigated using DSC at a very low cooling rate of 5 K/min. The crystallization kinetics of each as-spun specimen was evaluated using DSC at different heating rates of 10, 20, 30, and 40 K/min. The annealing experiment was carried out by keeping the as-spun amorphous specimens in the tubular furnace preheated to annealing temperatures for 600 s under vacuum atmosphere followed by water quenching. The density was measured by Archimedes' method. The saturation magnetic polarization (Js) and coercive force (Hc) were measured at room temperature with a vibrating sample magnetometer (VSM) under an applied field of 150 kA/m. For alloys with a smaller Hc than 100 A/m, Hc was measured by a DC B-H loop tracer under a field of 400 A/m. 3. Results 3.1. Thermal stability Fig. 1(a)–(c) display DSC curves of the melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys at different heating rates from 10 to 40 K/min showing the glass transition and crystallization behaviors. Upon heating, each of the curves exhibits an obvious endothermic event characteristic of the glass transition and a supercooled liquid region, followed by exothermic reactions corresponding to the crystallization of the undercooled liquid. From Fig. 1(a) and (b), the Fe80P11C9 and Fe80P9B2C9 glassy alloys exhibit a weak exothermic shoulder peak before the main exothermic peak which is caused due to the overlapping of two crystallization peaks. The substitution of 3 at.% Fe by Al in Fe80P9B2C9 alloy leads to disappearing of the shoulder peak as shown in Fig. 1(c). The values of the glass transition temperature (Tg) and onset crystallization temperature (Tx) determined from the DSC traces of 20 K/min heating rate are summarized in Table 1. Both Tg and Tx increase with substitution of 2 at.% P by B element in Fe80P11C9 alloy as well as with further substitution of 3 at.% Fe by Al in Fe80P9B2C9 alloy. The melting and solidification behaviors of these alloys are shown in Fig. 2. The near invariance of Tm for these compositions indicates that they may have the same eutectic reaction. Upon cooling, the Fe80P11C9 alloy shows two exothermic peaks. The first peak at the high temperature corresponds to the solidification of the primary phases, while the second one at the relatively low temperature is ascribed to that of the eutectic phase. When compared with Fe80P11C9 alloy, the first peak for Fe80P9B2C9 alloy is shifted to a lower temperature, which leads to the overlapping of two peaks in the DSC curve. Consequently, Tl is also lowered from 1262 to 1240 K as shown in Table 1. The further decrease of Tl is caused by the addition of Al element. Therefore, it is suggested that Fe77Al3P9B2C9 alloy with the lowest Tl is the closest to the eutectic point among these three alloys. 3.2. Crystallization behavior Fig. 3(a)–(c) shows XRD patterns of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys subjected to annealing for 600 s at different temperatures. The XRD patterns of as-spun specimens are also shown for comparison. Obviously, after annealing at (Tg − 20) K, these alloys maintain glassy structure. As we have known, the isothermal heat treatment in the supercooled liquid region can induce phase separation,
Fig. 1. DSC curves of (a) Fe80P11C9, (b) Fe80P9B2C9, and (c) Fe77Al3P9B2C9 glassy alloys at different heating rates from 10 to 40 K/min, showing glass transition and crystallization behaviors. Arrows indicate the glass transition temperature (Tg) and onset crystallization temperature (Tx), respectively.
nucleation, and growth of nuclei. For Fe80P11C9 alloy, as shown in Fig. 3(a), α-Fe phase primary precipitates from the glassy matrix after annealed at 695 K and 700 K in the supercooled liquid region. The result
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Table 1 Critical diameters (dc), thermal parameters, and activation energy of crystallization (Ec) for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys. Tg, Tx, and Tl are the glass transition temperature, onset crystallization temperature, and liquidus temperature, respectively. ΔTx is equal to Tx − Tg, and Trg is equal to Tg / Tl. Alloys
dc (mm)
Tg (K)
Tx (K)
Tl (K)
ΔTx (K)
Trg
Ec (kJ·mol−1)
Fe80P11C9 Fe80P9B2C9 Fe77Al3P9B2C9
1.5 1.8 3.0
690 698 712
720 727 751
1262 1240 1234
30 29 39
0.547 0.563 0.577
486 ± 31 496 ± 32 537 ± 18
shows that the exothermic shoulder peak in the Fe80P11C9 alloy is attributed to the primary crystallization of α-Fe phase. Here, it should be pointed out that, since the overlapping of crystallization peaks, the annealing temperature range to obtain α-Fe phase is very narrow. When annealing temperature (Ta) goes up to 760 K beyond the main exothermic peak, the pattern shows Bragg peaks of α-Fe, Fe3P and Fe3C phases superimposed on the broad diffuse scattering peak from the amorphous phase. For the specimen annealed at 880 K corresponding to temperature higher than the second crystallization peak, the crystallization phases remain unchanged while the intensities of the crystalline components increase, indicating the fraction of crystallization phases increases. This infers that the second peak observed by DSC is most likely due to crystallization of the residual amorphous phase. Similar crystallization process was observed for Fe80P9B2C9 alloy as well, as shown in Fig. 3(b). That is, the primary α-Fe phase is precipitated at 703 K and 708 K, while Fe3P and Fe3C phases are formed when Ta is raised to 780 K above the main exothermic peak. For Fe77Al3P9B2C9 alloy, formation of α-Fe phase is suppressed, as evidenced by the disappearing of the shoulder peak. Fig. 3(c) confirms that the diffraction peak corresponding to α-Fe does not appear in the XRD pattern obtained after annealed at 720 K corresponding to the temperature just below the first exothermic peak, suggesting that the supercooled liquid of the Fe77Al3P9B2C9 alloy crystallizes in different path by a simultaneous formation of Fe3P, Fe3C and Fe23(C,B)6. The same phases can be found when Ta is up to 800 K above the first exothermic reaction. The XRD pattern of the specimen annealed at 930 K corresponding to the temperature above the second peak are identified as the mixed phases of Fe3P, Fe3C, α-Fe and Fe23(C,B)6. 3.3. Magnetic properties The magnetic hysteresis loops of the melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloy ribbons are shown in Fig. 4. All the
Fig. 3. XRD patterns of (a) Fe80P11C9, (b) Fe80P9B2C9, and (c) Fe77Al3P9B2C9 glassy alloys subjected to annealing for 600 s at different temperatures. The XRD patterns of as-spun specimens are also shown for comparison.
Fig. 2. DSC curves of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys, showing melting and solidification behaviors. Arrows indicate the melting temperature (Tm) and liquidus temperature (Tl), respectively.
alloy specimens exhibit typical soft magnetic hysteresis curves. The Js increases from 1.37 to 1.46 T as substituting 2 at.% P by B, but decreases from 1.46 to 1.35 T as substituting 3 at.% Fe by Al. Here, it should be pointed out that Js for Fe80P9B2C9 amorphous alloy in this study is slightly higher than that reported in Ref. 29. This difference may be caused by different purities of the raw materials used in the two experiments. The Hc values for melt-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 ribbons are 16.4, 8.8, and 12.1 A/m, respectively. Fig. 5(a)–(c) show magnetic hysteresis loops of Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 specimens annealed at different temperatures. For comparison, the data of as-spun specimens are also plotted. The values of Js for the glassy and annealed specimens are listed in Table 2. As can be seen, the change of Js for all alloys is very small with increasing Ta to the temperatures lower than Tg. For the Fe80P11C9 and
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Fig. 4. Magnetic hysteresis loops of the as-spun Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloy ribbons.
Fe80P9B2C9 specimens annealed at 695 K and 703 K, respectively, the Js values increase to 1.49 T and 1.60 T. Such a high Js for Fe80P9B2C9 alloy is rarely observed in the Fe-based bulk glassy materials. The appearance of a single precipitate of α-Fe phase in glassy matrix could be responsible for the high Js [31,32]. With further increasing Ta, Js values of the Alfree alloys decrease. It should be pointed out that, for the Fe77Al3P9B2C9 alloy, annealing at 720 K causes a slight decrease in Js and a dramatic increase in Hc as shown in Fig. 5(c). The result indicates that Js of this alloy could not be effectively improved by annealing experiment. The Hc values for the as-spun and annealed specimens are also listed in Table 2. The changes in Hc for Fe80P11C9 and Fe80P9B2C9 alloys as a function of Ta are shown in Fig. 6. There is no remarkable different for the trend of the curves between these two alloys. Compared with the as-spun specimens, the specimens annealed at (Tg − 20) K exhibit a lower Hc value of about 3.0 A/m, which can be attributed to the stress releasing. The further increase in Ta to 695 K for Fe80P11C9 and 703 K for Fe80P9B2C9 causes a slight increase in Hc to 3.8 A/m and 3.6 A/m. After that, Hc increases sharply to 880 A/m and 780 A/m as increasing Ta to 700 K and 708 K for Fe80P11C9 and Fe80P9B2C9, respectively. Consequently, we can speculate that enhanced Js as well as good softness for the Fe80P11C9 and Fe80P9B2C9 amorphous alloys can be obtained by annealing the specimens at 698 K and 703 K, respectively. 4. Discussion 4.1. Relations of GFA with thermal stability and crystallization behavior Previously, we have reported that the critical diameters of glassy alloy rods (dc) are up to 1.5, 1.8, and 3.0 mm for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloys [27,29,30], respectively. Although the purity of Fe used in this experiment for Fe80P9B2C9 is a little different from that reported in Ref. [29], glassy rods of this composition with a diameter of 1.8 mm were also synthesized (not shown here). Obviously, the GFA of Fe80P11C9 glassy alloy can be slightly enhanced by partial substituting 2 at.% B for P. Further enhancement in the size of glassy rods is observed when 3 at.% Al is substituted for Fe. The reason for the improvement of GFA in the present Fe–P–C-based alloy system can be explained from the thermodynamic point of view. As shown in Fig. 2 and Table 1, addition of the proper amount of B to ternary Fe80P11C9 and Al to quaternary Fe80P9B2C9 can stabilize the supercooled melts by lowering the liquidus temperature. This means that the undercooled melts can be more easily frozen to be glassy against the formation of the competing crystalline phases. As a result, the reduced glass transition temperature (Trg = Tg / Tl), which has been regarded as one of the important parameters for evaluation of glass formation [33], increases with the addition of B in Fe–P–C alloy or Al in Fe–P–B–
Fig. 5. Magnetic hysteresis loops of (a) Fe80P11C9, (b) Fe80P9B2C9, and Fe77Al3P9B2C9 specimens annealed at different temperatures. For comparison, the data of as-spun specimens are also plotted. The inset shows the enlarged partial hysteresis curves of the specimens.
C alloy. The result shows that Trg effectively reflects GFA in Fe–P–Cbased glassy alloys. The variation trend of GFA for Al-bearing Fe-based glassy alloy can be related with the crystallization behavior of glassy alloys. When compared with Fe80P9B2C9 glassy alloy, the crystallization of Fe77Al3P9B2C9 alloy upon continuous heating takes place by the simultaneous precipitation of three phases α-Fe, Fe3P and Fe23(C,B)6. It is considered that the formation of α-Fe phase upon devitrification is drastically impeded by alloying with small amount of Al, because the mixing enthalpies with a large negative value between Al and Fe could increase the difficulty of the formation of the α-Fe phase. The competitive formation of
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Table 2 Saturation magnetic polarization (Js), coercive force (Hc), and structure for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 alloy specimens annealed at different conditions. HT denotes heat treatment. Compositions
Annealing conditions
Structures
Js (T)
Hc (A/m)
Fe80P11C9
As-spun HT at 670 HT at 695 HT at 700 HT at 760 As-spun HT at 678 HT at 703 HT at 708 HT at 780 As-spun HT at 683 HT at 720
Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am. Am.
1.36 1.37 1.49 1.44 1.34 1.46 1.46 1.60 1.56 1.43 1.37 1.37 1.30
16.4 3.3 3.8 2200 15,000 8.8 3.1 3.6 1800 12,700 12.1 2.0 11,300
Fe80P9C9B2
Fe77Al3P9B2C9
K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s K for 600 s
these crystallization phases results in the increase of stability of the supercooled liquid against crystallization. Meanwhile, the formation of Cr23B6-type Fe23(C,B)6 phase with complex face-centered cubic structure which requires long-range rearrangements of the constitute atoms also contributes to the high stability of the supercooled liquid [17,34,35]. These two effects cause the remarkable extension of the supercooled region (ΔTx = Tx − Tg) from 29 K to 39 K as shown in Table 1. The larger ΔTx revealing a higher resistance against crystallization could also be regarded as an indicator of higher GFA of alloys [36]. Therefore, we can conclude that higher GFA is obtained for the 3 at.% Al-bearing glassy alloy. However, the crystallization process cannot be changed by minor adding B in Fe–P–C alloy, which is responsible for the little change in ΔTx. At this point, it is suggested that the crystallization behavior and ΔTx cannot solely illustrate the GFA for the present Alfree glassy alloys. In addition, we further investigated the crystallization kinetics of amorphous alloys. The apparent activation energy (Ea) can be evaluated by the Kissinger equation [37]: ln β=T 2 ¼ −Ea =ðRT Þ þ C;
ð1Þ
where β is heating rate, T is the characteristic temperature and C is a constant. As observed from Fig. 1, Tx marked by arrows in the DSC traces are sensitive to the heating rate, and they shift to higher temperatures with increasing heating rate. This means that the crystallization have remarkable kinetic effects. By plotting ln(β/T2x) versus 1000/Tx, good linear fits can be obtained as shown in Fig. 7. From the slopes of these straight lines, the activation energy for primary crystallization (Ec) is determined as 486, 496, and 537 kJ/mol for Fe80P11C9, Fe80P9B2C9, and
Fig. 6. Changes in coercive force (Hc) for Fe80P11C9 and Fe80P9B2C9 alloys as a function of annealing temperature (Ta).
+ α-Fe (3–5) nm + α-Fe (~61 nm) + Fe3P + Fe3C
+ α-Fe (3–5) nm + α-Fe (~78 nm) + Fe3P + Fe3C
+ Fe3P + Fe3C + Fe23(C,B)6
Fe77Al3P9B2C9 glassy alloys, respectively. It is well known that the onset crystallization temperature is considered to be associated with the nucleation process. Thus, the activation energy deduced from the onset crystallization temperatures corresponds to nucleation [38]. The result of crystallization kinetics indicates that the addition of B in Fe– P–C alloy or Al in Fe–P–B–C alloy makes nucleation more difficult, which is responsible for the enhancement of GFA.
4.2. Relations of magnetic properties with alloy composition and crystallization behavior As well known, the Js is an intrinsic magnetic property that is mainly determined by alloy composition and phase constitution of alloy. For Fe–metalloid-type glassy alloys, the effect of metalloid elements on saturation magnetization has been extensively discussed. According to charge-transfer model [39], the valence electrons of metalloid elements transfer to the band of Fe which contributes to magnetic moment. When metallic elements are doped into a base alloy, the valence electronic structures and magnetic moment alignments of the added elements, as well as their electronic interactions with Fe, markedly affect the magnetic moments. Generally, the addition of metalloid elements would reduce the Js of amorphous alloys, and the metalloid element with more sp electrons results in the more decrease in Js. The number of sp electrons for the metalloid elements P, C, and B is 5, 4, and 3, respectively. Therefore, the Fe80P9B2C9 alloy with a lower proportion of P exhibit a higher Js compared with the Fe80P11C9 alloy. Further, the introduction of non-ferromagnetic Al element to Fe80P9B2C9 alloy reduces the number of magnetic atoms, and thus leads to a decrease in Js [40,41].
Fig. 7. Kissinger plots of ln(β/T2x) versus 1000/Tx for Fe80P11C9, Fe80P9B2C9, and Fe77Al3P9B2C9 glassy alloys.
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J. Wang et al. / Journal of Non-Crystalline Solids 454 (2016) 39–45
Herzer [42] proposed a simple two-phase model to quantitatively predict the change of Js in annealed glassy alloy specimens. According to this model, the saturation magnetic polarization reflects the ratio of the volume fraction of the crystalline phase Vc/V to that of the amorphous phase Va/V with V = Vc + Va, and is expressed as Js = JscVc/ V + JsaVa/V, where Jsc and Jsa are the saturation magnetic polarization in the crystalline and amorphous phases, respectively. The increase in Js for the annealed Fe80P11C9 and Fe80P9B2C9 specimens with a single precipitation of α-Fe phase can be understood in terms of the contribution of α-Fe nanocrystallites with high saturation magnetic polarization of about 2.1 T [43]. With the further increasing of Ta, the grain size and volume fraction of α-Fe phase increase which leads to less Fe concentration leaving in amorphous matrix. The enrichment of nonmagnetic atoms including P, C and B in the matrix is responsible for a possibly rapid decrease in Jsa, and so as to a slight decrease of Js. On the other hand, it is well known that the magnetic softness for the partial nanocrystalline alloys strongly depends on the grain size of α-Fe phase [44–46]. The effective anisotropy (Keff), intimately related with Hc, is highly sensitive to the grain size (D) of nanocrystalline phase. A small increase in D may dramatically increase the Keff, which results in deterioration of the magnetic softness. As shown in Fig. 3(a) and (b), for the Fe80P11C9 and Fe80P9B2C9 specimens obtained after annealing at temperatures of 695 K and 703 K, respectively, the diffraction intensity of α-Fe is weak and the volume fraction of crystalline phase is small, so the grain size is impossible to estimate using Scherer's formula. Previous research [27] on the microstructure of crystallized Fe80P11C9 specimens has shown that ultrafine α-Fe crystallites with a diameter of about 3– 5 nm is embedded in the amorphous matrix. The SAED pattern shows only a few spots scatter in the diffraction rings, which indicates the volume fraction of α-Fe phase is small. The results are agreement with the XRD analysis. The low volume fraction of nanoscale α-Fe grains may have little influence on Keff of matrix [45,46]. Thus, the Hc of these specimens remains almost unchanged compared with the ones after relaxation. As for the Fe80P11C9 and Fe80P9B2C9 alloy specimens annealed at higher temperatures of 700 K and 708 K, respectively, obvious crystalline peak corresponding to (200) α-Fe phase is observed. The grain size estimated on the Scherer's formula from (200) plane is up to 61 nm and 78 nm, respectively. Accordingly, we could speculate that the dramatic increase in Hc is due to the increase in D. The specimens annealed at temperatures higher than the main exothermic peak exhibit a larger Hc because they compose of α-Fe phase and Fe-metalloid compounds with high magnetocrystalline anisotropy.
5. Conclusions In this work, the substitution effect of B and Al elements on the thermal properties, crystallization behavior, and magnetic properties of Fe– P–C-based amorphous alloys were investigated. The addition of the proper amount of B to ternary Fe80P11C9 and Al to quaternary Fe80P9B2C9 can stabilize the supercooled melts by lowering the liquidus temperature, which leads to the increase in Trg. The annealing experiments shows that α-Fe is the primary precipitation phase for the Fe80P11C9 and Fe80P9B2C9 alloys, while α-Fe, Fe3P and Fe23(C,B)6 are the competing crystalline phases for the Fe77Al9P9B2C9 alloy. Therefore, the suppression of the formation of α-Fe during crystallization increases the thermal stability of glass phase and thus improves the GFA significantly. The result of crystallization kinetics indicates that the addition of B in Fe80P11C9 alloy or Al in Fe80P9B2C9 alloy makes nucleation more difficult, which is in agreement with GFA. The relationship between crystallization and magnetic properties was also discussed. It is found that the resulting annealed Fe-based alloy specimen with nanoscale α-Fe grains exhibits high Js of up to 1.60 T and low Hc of about 3 A/m. However, since the optimum annealing temperature range for Al-free Fe-based amorphous alloys to obtain α-Fe phase is very narrow, further improvement in the soft magnetic properties is difficult.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Grant No. 51501166), Foundation of Henan Educational Committee (Grant Nos. 13A430668 and 15A430002), and Outstanding Young Talent Research Fund of Zhengzhou University (Grant No. 1421320046).
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