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Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys
Accepted Manuscript Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys A. Grabias, M. Kopcewicz, J. Latuch, D. Oleszak, M. Pękała, M. Kowalczyk PII: DOI: Reference:
S0304-8853(16)33437-0 http://dx.doi.org/10.1016/j.jmmm.2017.03.047 MAGMA 62570
To appear in:
Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
22 December 2016 18 March 2017 24 March 2017
Please cite this article as: A. Grabias, M. Kopcewicz, J. Latuch, D. Oleszak, M. Pękała, M. Kowalczyk, Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys, Journal of Magnetism and Magnetic Materials (2017), doi: http://dx.doi.org/10.1016/j.jmmm.2017.03.047
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Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys A. Grabiasa,*, M. Kopcewicza, J. Latuchb, D. Oleszakb, M. Pękałac, M. Kowalczykb
a
Institute of Electronic Materials Technology, Wólczyńska 133, 01-919 Warsaw, Poland
b
Faculty of Materials Science and Engineering, Warsaw University of Technology, Wołoska
141, 02-507 Warsaw, Poland c
Department of Chemistry, University of Warsaw, Al. Żwirki i Wigury 101, 02-089 Warsaw,
Poland
*Corresponding author: Agnieszka Grabias Institute of Electronic Materials Technology Wólczyńska 133 01-919 Warsaw, Poland Tel.: +48 22 6395547 Fax: +48 22 8645496 E-mail: [email protected]
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Abstract
The influence of Co content on the structural and hard magnetic properties of two sets of nanocrystalline Fe52-xCoxPt28B20 (x = 0 − 26) and Fe60-yCoyPt25B15 (y = 0 − 40) alloys was studied. The alloys were prepared as ribbons by the rapid quenching technique. The nanocomposite structure in the alloys was obtained by annealing at 840-880 K for 30 min. Structural characterization of the samples was performed using the Mössbauer spectroscopy and X-ray diffraction. Magnetic properties of the samples were studied by the measurements of the hysteresis loops and of the magnetization at increasing temperatures. An amorphous phase prevailed in the as-quenched Fe52-xCoxPt28B20 alloys while a disordered solid solution of fcc-(Fe,Co)Pt was a dominating phase in the Fe60-yCo yPt25B15 ribbons. Differential scanning calorimetry measurements revealed one or two exothermic peaks at temperatures up to 993 K, depending on the composition of the alloys. Thermal treatment of the samples led to the formation of the magnetically hard ordered L10 tetragonal (Fe,Co)Pt nanocrystallites and magnetically softer phases of (Fe,Co)B (for Fe52-xCoxPt28B20) or (Fe,Co)2B (for Fe60-yCoyPt25B15). Detailed Mössbauer spectroscopy studies revealed that cobalt substituted for iron in both the L1 0 phase and in iron borides. The nanocomposite Fe60-yCo yPt25B15 alloys exhibited significantly larger magnetic remanence and maximum energy products but a smaller coercivity than those observed for the Fe52-xCoxPt28B20 alloys. Co addition caused a reduction of the magnetization and the energy product in both series of the alloys. The largest magnetic remanence of 0.87 T and the highest energy product (BH)max = 80 kJ/m3 were obtained for the Co-free Fe52Pt28B20 alloy while the largest coercivity (HC > 950 kA/m) was observed for the Fe50Co10Pt25B15 and Fe30Co30Pt25B15 alloys. Differences in the hard magnetic properties of the nanocomposite alloys were related to different phase compositions influencing the strength of inter-phase exchange coupling interactions.
Keywords: nanocomposite alloy; Fe-Pt-B alloy; hard magnetic properties; Mössbauer spectroscopy
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1. INTRODUCTION
The ordered L1 0 face-centered-tetragonal FePt phase with a high magnetocrystalline anisotropy energy (K = 7 MJ/m3) has been widely used as a hard magnetic component in designing various structures in the form of thin films, multilayers, nanoparticles or nanocomposite alloys for applications in ultrahigh-density magnetic storage media or as permanent magnets [1-4]. It was shown that in Fe-Pt-B alloys a nanocomposite structure, consisting of magnetically hard, ordered L10 FePt phase and soft magnetic, disordered facecentered-cubic FePt and Fe2B phases, can be used for the preparation of exchange-coupled nanocrystalline magnets with a high energy product up to about 120 kJ/m3 [4-7]. Such nanocomposite structures can be formed directly by rapid quenching [7] or by annealing disordered alloys [4-6]. It is also known that the ordered L1 0 CoPt phase reveals a high magnetocrystalline anisotropy similar to that of the L1 0 FePt one. Therefore, in search of enhanced hard magnetic properties the preparation and magnetic properties of ternary (FeCo)Pt thin films have been studied in, e.g., [8-11]. Only several works devoted to Co-containing Fe-PtB alloys have been published [12,13]. In all cases it was found that Co substitutes for Fe in the ordered L10 FePt phase. In the Fe49-xCo xPt51 thin films the substitution of Fe with Co resulted in a higher magnetization and a lower coercivity as compared with the FePt film [9]. A refined grain size and a reduced long-range order parameter were observed in ternary L10 FeCoPt films [9]. In the nanocrystalline (Fe,Co)-Pt-B alloys with a fairly low Pt content the increase of Co addition caused a gradual increase of the coercivity, however, the magnetic remanence and the maximum energy product became smaller [12]. In the nanocomposite (Fe,Co)-Pt-Zr-B melt-spun alloys a monotonous decrease of the coercivity was observed with the increase of Co content, which was attributed to a reduction of the magnetocrystalline anisotropy originating from decreasing the long-range order parameter of the L1 0 phase [13]. The differences in hard magnetic properties of the Fe-Pt-B-based nanocomposite alloys strongly depend on their multi-phase nanostructure features, such as the chemical composition and ordering of the phases, their relative fraction, and the size of the nanograins. The aim of the present work was to investigate the influence of Co substitution for Fe on the formation of nanocomposite structures and on hard magnetic properties of two sets of (Fe,Co)-Pt-B alloys with different Pt and B contents. Therefore, two series of Fe52-xCoxPt28B20 (x = 0 − 26) and Fe60-yCoyPt25B15 (y = 0 − 40) disordered alloys were prepared by the melt spinning technique. The rapidly quenched alloys were annealed in order to obtain a nanocomposite structure with L10 FePt nanograins. The phase composition of the as-quenched and annealed samples was
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probed by 57Fe Mössbauer spectroscopy measurements, which provided valuable information regarding local atomic arrangements in iron-containing phases as well as the magnetic behavior of iron atoms. This experimental method is sensitive enough to detect even small composition fluctuations in Fe environments. The Mössbauer spectroscopy has been successfully applied for studying magnetic systems of Fe-Pt [14-19] and Fe-Pt-B [20-23]. However, according to our best knowledge, the Mössbauer spectroscopy has not yet been used in studying Co-containing (Fe,Co)-Pt-B alloys, except for our recent publication [24]. In the present work, the Mössbauer spectroscopy proved to be very useful in identification of the phases and determination of their chemical compositions and relative fractions, thus enabling the study of the relations between the nanostructure and magnetic properties of the (Fe,Co)-Pt-B alloys. Understanding these relations is important for the application of the alloys as exchange-coupled permanent magnets.
2. EXPERIMENTAL DETAILS
The Fe52-xCo xPt28B20 (x = 0, 10, 18, 26 at.%) and Fe60-yCo yPt25B15 (y = 0, 10, 30, 40 at.%) alloys were prepared by melt spinning under Ar protective atmosphere. The linear velocity of 45 m/s at the spinning wheel was applied. The ribbons were about 1 mm wide and 15-20 µm thick. Differential scanning calorimetry (DSC) measurements were performed for the rapidly quenched alloys in the range of 300−993 K with the heating rate of 40 K/min. The samples were annealed at 840−880 K for 30 min. in vacuum. Structural properties of the as-quenched and annealed ribbons were investigated by the Mössbauer spectroscopy and X-ray diffraction (XRD) using CuKα radiation. All Mössbauer spectra were measured in transmission geometry at room temperature. A constant-acceleration spectrometer with a 57Co-in-Rh source with activity of about 25 mCi was used. Mössbauer spectra were fitted using the NORMOS program [25]. The spectra of the as-quenched alloys were fitted with a hyperfine field distribution P(Bhf) method. The spectra of the crystalline samples were analyzed using the least squares method, assuming the Lorentzian profile of lines. Determination of the hyperfine parameters, such as magnetic hyperfine field, Bhf, isomer shift, δ, and quadrupole shift, ∆, characteristic for magnetically split six-line spectral components or quadrupole splitting and isomer shift for doublets, enabled the identification of iron-containing magnetic and non-magnetic phases. Relative fractions of the spectral components related to the identified phases were calculated as a ratio of the area of the relevant subspectrum to the total spectral area, assuming similar Debye-Waller factors for each phase. Isomer shifts are related to α-Fe standard at room temperature. Temperature dependence of magnetization of the as-quenched samples was measured in the temperature range of 290-980
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K under applied magnetic field of 1.5 T using a Faraday balance. A vibrating sample magnetometer (VSM) with maximum applied magnetic field of 2 T was used to investigate magnetic properties of the samples at room temperature.
3. RESULTS AND DISCUSSION 3.1. As-quenched alloys
Detailed XRD and Mössbauer results obtained for the as-quenched Fe52-xCoxPt28B20 (x = 0 - 26) and Fe60-yCo yPt25B1 5 (y = 0 - 40) alloys were presented in our previous work [24]. It was found that the rapidly quenched Fe52-xCoxPt28B20 alloys were predominantly amorphous with a small contribution of the disordered fcc-FePt phase. In the as-quenched Fe60-yCo yPt25B15 alloys the fcc-(Fe,Co)Pt disordered solid solution coexisted with amorphous (Fe,Co)-B regions. Thermal stability of the rapidly quenched (Fe,Co)-Pt-B alloys was studied by DSC measurements. Examples of DSC curves are shown in Fig. 1 for the extreme compositions studied: without Co and with the highest Co content. The characteristic temperatures, determined from DSC curves, are collected in Table 1 for all alloys studied. For the Fe52-xCoxPt28B20 alloys two separated exothermic peaks are observed whereas for the Fe60-yCo yPt25B15 samples a broad exothermic effect is recorded. In the case of predominantly amorphous Fe52-xCoxPt28B20 alloys a weak peak at lower temperatures is attributed to thermally induced atomic rearrangements in the amorphous phase, leading to a formation of the fcc-(Fe,Co)Pt disordered solid solution. This exothermic effect is more pronounced for larger Co content, suggesting larger amount of the amorphous phase in Co-containing alloys. As seen from Table 1, the onset temperature of this peak, Tx1, increases with increasing Co content (from 648 to 686 K for x = 0 and 26, respectively), indicating that Co addition increases thermal stability of the amorphous phase. Similar finding has been reported in [12]. Such an exothermic peak is not seen for the Fe60-yCoyPt25B15 alloys, in which the fcc-(Fe,Co)Pt disordered solid solution is already formed in the as-quenched state. The main exothermic effect, observed for all alloys studied, is attributed to the formation of iron and cobalt borides and the transformation of the fcc-(Fe,Co)Pt disordered solid solution to the ordered tetragonal phase as discussed in the next section. Onset temperatures of these phase transformations are in the range of 788−856 K and are significantly higher for the set of the Fe60-yCoyPt25B15 alloys than for the Fe52-xCoxPt28B20 ones.
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3.2. Annealed alloys
In order to form a nanocomposite structure consisting of the hard and soft magnetic phases, the FeCoPtB alloys were annealed for 30 min. at temperatures corresponding to the main exothermic peak in the DSC curves (Fig. 1). In each case the selected annealing temperatures were slightly larger than Tp2 (Table 1), ensuring required phase transformations. The XRD patterns of the Fe52-xCo xPt28B20 (x = 0 and 26) and Fe60-yCo yPt25B15 (y = 0 and 40) alloys annealed for 30 min. at 840 K and 880 K, respectively, are presented in Fig.2. In each pattern two sets of diffraction lines related to two phases are identified. All XRD patterns contain diffraction lines characteristic of the ordered L10 fct-(Fe,Co)Pt phase. The presence of cobalt in the L1 0 phase is evidenced for Co-containing alloys by a slight shift towards larger 2θ angles of the diffraction lines characteristic of the ordered L10 FePt phase, which is best seen in the pattern obtained for Fe20Co40Pt25B15 (Fig. 2d, y = 40). It is worth noting that a formation of the ordered, ternary (Fe,Co)Pt phase instead of a mixture of the L10 CoPt and L10 FePt phases was evidenced in annealed (Co1-xFex)Pt thin films [8]. The average crystallite size of the L1 0 phase was determined by the Williamson-Hall method (Table 2). It increases from 20 to 24 nm with increasing Co content from x = 0 to x = 26. In the case of the Fe60-yCoyPt25B15 samples the average crystallite size is contained in the narrow range of 30−32 nm. The second set of less intense diffraction lines is attributed to the formation of borides: FeB (Fig. 2a) or (Fe,Co)B (Fig. 2b) for the Fe52-xCo xPt28B20 alloys, and Fe2B (Fig. 2c) or (Fe,Co)2B (Fig. 2d) for the Fe60-yCoyPt25B15 alloys. In the case of iron borides Co substitution for Fe also causes a shift of diffraction lines towards larger 2θ angles. Mössbauer spectra recorded for the Fe52-xCoxPt28B20 alloys annealed at 840 K are shown in Fig. 3. The Mössbauer spectrum of the Co-free alloy is fitted with two magnetic components (Fig. 3a). The S1 sextet with hyperfine field Bhf = 27.8 T, isomer shift δ = +0.27 mm/s and quadrupole shift ∆ = 0.30 mm/s corresponds to the tetragonal ordered L10 FePt phase [14,22]. The S3a sextet with Bhf = 10.8 T, δ = +0.26 mm/s and ∆ = 0.12 mm/s is assigned to the crystalline FeB phase [26]. The relative fractions of these two sextets (57% and 43% for S1 and S3a, respectively), calculated from the Mössbauer spectrum, are directly related to two structurally different positions of iron nuclei in the two-phase sample, and indicate that both phases have nearly equiatomic compositions with a slightly larger Fe content as compared with the stoichiometric L1 0 FePt and FeB phases. Incorporation of 10% of cobalt causes the appearance of two additional subspectra (Fig. 3b). Beside the S1 sextet related to the L1 0 FePt phase the S2 sextet with Bhf = 28.6 T, δ = +0.27 mm/s and ∆ = 0.17 mm/s is observed. In the
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central part of the spectrum a clear quadrupole doublet D1 with quadrupole splitting of 0.60 mm/s and isomer shift δ = +0.26 mm/s appears. With increasing Co content in the alloy the relative spectral fraction of the S2 sextet increases at the expense of the S1 one (Fig. 3b-3d). The hyperfine parameters of the S2 magnetic component strongly suggest a formation of the ternary tetragonal (Fe,Co)Pt phase. The value of the quadrupole shift obtained for the S2 sextet (∆S2) is large but significantly smaller as compared with ∆S1 = 0.30 mm/s observed for the ordered L1 0 FePt phase, which provides an evidence of somewhat less distorted local atomic symmetry in the tetragonal (Fe,Co)Pt phase than in the L1 0 FePt one. Furthermore, the ∆S2 parameter decreases with increasing Co content from 0.17 to 0.11 mm/s for x = 10 and 26, respectively. It was shown experimentally that Co substitution for Fe in the L10 FePt phase causes a decrease of the ordering and the lattice parameters ratio c/a becomes closer to 1 as compared with the data of the ordered binary FePt phase with the same Pt content [8,9,27]. For larger Co content the S3a sextet, assigned to the magnetic FeB phase, is not observed in the spectra (Fig. 3c, 3d). The appearance of the quadrupole doublet D1 in the place of the S3a sextet is related to the incorporation of Co into the FeB lattice. A presence of Co in this iron boride is already evidenced for x = 10 by a reduction of the hyperfine field of the S3a sextet to about 8 T (Fig. 3b). Further substitution of Fe with Co in the (Fe,Co)B compound leads to the formation of the nonmagnetic phase as seen by the appearance of the quadrupole doublet. It was reported for the orthorhombic Fe1-xCoxB system that the magnetic moment decreases almost linearly with the increase of Co content and then it drops to zero for x ≥ 0.85 [28]. Changes in the relative fractions of all spectral components with increasing Co content in the alloys studied are discussed in detail below (Fig. 6). Mössbauer spectra recorded for the Fe60-yCoyPt25B15 alloys annealed at 880 K are shown in Fig. 4. The spectrum of the Fe60Pt25B15 alloy (Fig. 4a) is more complex than that of the Fe52Pt28B20 alloy (Fig. 3a). The spectrum in Fig. 4a is fitted with 4 magnetic components. The sextet S3b with Bhf = 23.7 T, δ = +0.11 mm/s and ∆ = 0.00 mm/s constitutes the prevailing component, which is characteristic of the tetragonal Fe2B compound [26]. It is worth noting that for the tetragonal Fe2B phase two slightly different Fe positions are observed, therefore, two sextets with Bhf of 23.2 and 24.2 T are usually fitted in the Mössbauer spectra [26]. Both sextets reveal, however, almost identical values of the isomer shift and the quadrupole shift (δ = +0.12 mm/s, ∆ = 0.02-0.04 mm/s) [26]. Therefore, in the case of the complex spectra in Fig. 4 one sextet with slightly broadened lines is fitted for Fe2B. The observed hyperfine field of the S3b component is the mean value. According to the fitting of the Mössbauer spectrum about 54% of Fe atoms are located in iron boride, which indicates that the Fe2B phase is nearly stoichiometric.
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The S1 sextet with Bhf = 28.1 T, δ = +0.27 mm/s and ∆ = 0.30 mm/s corresponds, as discussed above (Fig. 3), to the tetragonal ordered L10 FePt phase. Beside these two components with major spectral fractions, two minor sextets are fitted. The S4 sextet with Bhf = 29.1 T, δ = +0.20 mm/s and ∆ = 0.00 mm/s is attributed to the remaining fcc-FePt phase [15,22]. The relative spectral fraction of this component is below 8%. The amount of the fcc phase, detected in the Mössbauer spectrum, is too small to be observed in the XRD pattern. A small spectral contribution (~4%) of the S5 sextet with a large hyperfine field of about 36 T, δ = +0.25 mm/s and ∆ = -0.31 mm/s is also detected in the spectrum in Fig. 4a. A very similar additional spectral component was observed for the binary Fe61Pt39 L1 0 ordered alloy [17]. It was attributed to Fe atoms on Pt-sites in the L10 structure. Therefore, the S5 sextet is assigned to iron-rich FePt environments in the L10 FePt phase. The origin of this component results most probably from the excess of iron content in the alloy. After incorporation of 10% of Co to the Fe-Pt-B alloy the shape of the Mössbauer spectrum changes significantly (Fig. 4b). It is fitted with 4 subspectra: the previously observed sextets S3b (related to Fe2B) and S1 and S5 (both related to the L10 phase), and a new S2 component. The S3b sextet reveals a slightly reduced mean hyperfine field of 23.1 T as compared with 23.7 T obtained for the Co-free alloy. This reduction is related to Co presence in the iron boride. The S2 sextet with Bhf = 28.8 T, δ = +0.25 mm/s and ∆ = 0.16 mm/s is assigned to the formation of ternary (Fe,Co)Pt phase with the ordered L1 0 structure. A similar spectral component (S2) with very close hyperfine parameters appears in the Mössbauer spectra of the annealed Fe52-xCo xPt28B20 alloys for x≥10 (Fig. 3). As can be seen from Fig. 4b-4d the intensity of the S2 sextet increases with increasing Co content. For cobalt concentrations exceeding 10% the S5 sextet, originating from the excess of iron atoms, disappears (Fig. 4c, 4d). Such observation in the Mössbauer spectra is obviously related to the reduction of Fe content in the alloys and the substitution of Fe atoms with Co atoms in the L10 phase. Therefore, the spectra of the samples with y = 30 and 40 are fitted with only 3 spectral components, originating from the ordered L10 FePt and (Fe,Co)Pt (S1 and S2, respectively) and (Fe,Co)2B (S3b) phases (Fig. 4c, 4d). The influence of Co addition on the Mössbauer spectra of the annealed Fe52-xCo xPt28B20 and Fe60-yCoyPt25B15 alloys is demonstrated in Fig. 5 and Fig. 6, in which hyperfine field values and relative fractions of each spectral component are plotted vs. Co content, respectively. As can be seen from Fig. 5 the hyperfine fields of the L10 FePt (S1) and (Fe,Co)Pt (S2) phases do not change substantially with increasing Co content and are in the narrow ranges of 27.3−28.1 T and 28.4−28.8 T, respectively. In the case of the Fe52-xCo xPt28B20 samples a significant reduction by
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about 3 T is observed for the hyperfine field related to the FeB phase (S3a) already for the addition of 10% of Co to the alloy (Fig. 5a). Further increase of cobalt content causes further substitution of Co for Fe in the (Fe,Co)B phase and the hyperfine field vanishes for x>10 (Fig. 3). A continuous decrease of the hyperfine field with increasing Co content is also observed for the (Fe,Co)2B phase (S3b, Fig. 5b). This reduction is directly related to a gradual substitution of Co for Fe in this phase. This is confirmed by thermodynamic calculations of ternary Fe-Co-B system revealing that Fe2B and Co2B form continuous solid solutions of (Fe,Co)2B [29]. The formation of continuous solid solutions of (Fe,Co)B by FeB and CoB was also demonstrated in the same work [29]. It is worth noting that we measured the Mössbauer spectrum of a crystalline sample with the composition of (Fe0.7Co0.3)2B, which exhibited the mean hyperfine field of 22.8 T (about 1 T lower than that of Fe2B). This value is only slightly smaller than Bhf = 23.1 T obtained for the sample with y = 10 (Fig. 5b), which indicates that the composition of the (Fe,Co)2B phase is close to that of about (Fe0.7Co0.3)2B. This is confirmed by the estimation based on the relative fraction of iron atoms contributing to the S3b spectral component, which gives the composition of (Fe0.74Co0.26)2B. For the alloys with y = 30 and 40 Co concentration in this phase prevails over Fe content as evidenced by a significant reduction of the hyperfine field to 21.6 and 19.7 T, respectively (Fig. 5b). The estimated composition is (Fe0.1Co0.9)2B for y = 40. Dependence of the relative fractions of all spectral components on the Co content in the alloys is presented in Fig. 6. In the case of the Fe52-xCoxPt28B20 alloys (Fig. 6a) the increase of Co content causes a significant decrease of the relative fractions of the L10 FePt (S1) and (Fe,Co)B (S3a+D1) components whereas the fraction of L1 0 (Fe,Co)Pt (S2) increases substantially. The decrease of the fraction of the (Fe,Co)B component (S3a+D1) is directly related to the decrease of iron content in this phase when Fe is substituted with Co. It is worth noting that Fe-free CoB or Co 2B phases would not be seen in Mössbauer spectra. The relative spectral fractions of the phases, calculated for the Fe60-yCoyPt25B15 alloys, are plotted in Fig. 6b. Comparing the results obtained for both series of the alloys clear differences in the relative fractions between the L10 FePt (S1) and Fe-B (S3) components are seen for both Co-free alloys (x = 0, y = 0). The relative fraction of L1 0 FePt (S1) is about 1.5 times larger for the Fe52Pt28B20 alloy than for Fe60Pt25B15. The relative fraction of FeB (S3a, Fig. 6a) is by about 10% smaller than that of Fe2B (S3b, Fig. 6b). Furthermore, unlike in the case of the Fe52Pt28B20 alloy a small abundance of about 7% of the fcc-FePt phase (S4) remained in the Fe60Pt25B15 sample after annealing (Fig. 6b), which also weighed on a reduced fraction of the ordered L10 phase. It is worth noting that the relative fractions of the S1, S2 and S3 components are quite similar for both alloys with the highest Co content (x = 26, y = 40). Thermally induced formation of the L10 (Fe,Co)Pt phase (S2) reached
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the relative fraction of about 61% and 57% for the Fe26Co26Pt28B20 and Fe20Co40Pt25B15 alloys, respectively.
3.3. Magnetic measurements
Temperature variations of magnetization in the range from room temperature up to 980 K, depicted in Fig. 7, show that the (Fe,Co)-Pt-B systems studied are ferromagnetically ordered up to high temperatures. The as-quenched Fe52-xCoxPt28B20 alloys with the prevailing amorphous phase reveal Curie temperatures, TC, in the range of 655 − 665 K (Fig. 7a). Subsequent decrease of the magnetization up to about 700 -720 K results from the paramagnetic amorphous phase and from the formed fcc disordered solid solution (the first DSC peak in Fig. 1). Then the crystallization process of a ferromagnetic phase with higher TC is initiated as indicated by the increase of the magnetization. This magnetization increase is markedly larger for x = 18 and 26 than for x = 0 and 10. This enhancement of the magnetization is related to the formation of the ordered L10 FePt phase with TC of about 750 K for x = 0 and the L1 0 (Fe,Co)Pt phase for x ≥ 10. The Curie temperature of the latter phase increases with the increase of Co content in the alloy (TC ≈ 825 K for x = 18). It is worth noting that the orthorhombic FeB phase has TC of about 600 K. Since the crystallization of this iron boride starts at temperatures exceeding TC, the formation of FeB is not seen in the M(T) dependence. The temperature variation of magnetization observed for the Fe60-yCoyPt25B15 alloys (Fig. 7b) indicates a multicomponent behavior, originating from the inhomogeneity of the as-quenched samples as revealed by Mössbauer spectroscopy results [24]. The Curie temperature of the as-quenched samples with y = 0, 10 and 30 increases with increasing Co content from about 740 K to about 760 K. The magnetization enhancement observed for these alloys at about 770 K is related to the crystallization of the (Fe,Co)2B phase. Somewhat different magnetization behavior is observed for the as-quenched Fe20Co 40Pt25B15 alloy (y = 40) with a quite large content of the fcc solid solution, for which TC reaches 860 K. In this case only a small magnetization increase is seen at about 780 K, which corresponds to the crystallization of the boride phase as observed in other samples of this series. The M(T) plots above 770 K enable to locate TC of the high temperature phase of (Fe,Co)2B between 960 and 980 K, depending on the chemical composition. These values are smaller than the nominal Curie temperature of the Fe2B phase (1015 K). The calculated room temperature magnetic moment per formula unit is plotted in Fig. 8 for both series of the as-quenched alloys. The dependence of the magnetic moment on Co content in the alloy represents a monotonically decreasing function with increasing cobalt
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content for both series. In the case of the as-quenched Fe52-xCo xPt28B20 alloys the magnetic moment decreases from 1.09 to 0.84 µB with the increase of Co content from x = 0 to x = 26. For the Fe60-yCo yPt25B15 series the values of magnetic moment are larger when comparing the alloys with the same Co content in both series, and decrease from 1.29 to 0.99 µB with the increase of Co content from y = 0 to y = 40. Therefore, in the (Fe,Co)-Pt-B alloys studied the substitution of iron atoms with cobalt atoms causes a decrease of the value of the magnetic moment per formula unit at room temperature. Room temperature hysteresis loops of the annealed alloys are presented in Fig. 9. The determined magnetic properties are collected in Table 2. All loops are smooth like in the case of a single-phase material, indicating inter-phase exchange-coupled interactions. The shape of the hysteresis loops shows that they approach saturation in magnetic fields somewhat larger than the applied field. The hysteresis loops recorded for the Fe52-xCoxPt28B20 samples exhibit good squareness and good hard magnetic properties (Fig. 9a). The coercivity exceeds 800 kA/m for all compositions studied (x) and reaches the largest values of 970 and 953 kA/m for the alloys with x = 10 and x = 18, respectively (Table 2). For these alloys quite similar values of 51-52 kJ/m3 of the maximum energy product, (BH)max, were obtained. The influence of Co substitution for Fe is seen in a reduction of the magnetization and of the energy product. Both the magnetic remanence (Mr) and the magnetization at 1.6 MA/m decrease with increasing Co content in the alloy (Fig. 9a, Table 2). The reduced remanence Mr/Ms is quite large of about 0.75-0.76 for all alloy compositions. The best hard magnetic properties in this series of the alloys were, however, obtained for the Co-free sample, for which (BH)max of about 65 kJ/m3 was observed. The hysteresis loops recorded for the Fe60-yCoyPt25B15 series reveal larger magnetization values but significantly smaller coercivity (up to 439 kA/m) and the reduced remanence (0.67-0.74) as compared with the Fe52-xCoxPt28B20 alloys (Fig. 9b, Table 2). The coercivity seems to be independent of the alloy composition for y ≤ 30. Also in this case the magnetic remanence, the saturation magnetization and the energy product decrease with increasing Co content. As in the previous case (Fig. 9a) the best hard magnetic properties were obtained for the Co-free Fe60Pt25B15 alloy with HC = 437 kA/m, Mr = 0.87 T, Mr/Ms = 0.74 and (BH)max of about 80 kJ/m3. The coercivity of the annealed Fe52-xCoxPt28B20 alloys is more than twice larger than that observed for the annealed Fe60-yCoyPt25B15 samples (Table 2). In the literature the largest coercivity was reported for the Fe49Pt33B18 alloy (HC > 955 kA/m, Mr = 0.66 T, (BH)max of about 71 kJ/m3) [5]. For the Fe52Pt30B18 alloy the values of HC = 783 kA/m, Mr = 0.70 T and (BH)max = 88 kJ/m3 were given [7]. The structure of both alloys consisted of the nanosized L10 FePt and
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Fe2B phases. The large coercivity is attributed to a quite large Pt content in the Fe-Pt-B alloys, which takes part in the formation of the L1 0 FePt phase revealing a high magnetocrystalline anisotropy. As shown by the Mössbauer measurements in the present work (Fig. 6) a significantly larger relative fraction (A) of the spectral component assigned to the hard magnetic L10 FePt phase and a larger coercivity were observed for the Fe52Pt28B20 alloy with larger Pt content (A~57%, HC = 823 kA/m) than for the Fe60Pt25B15 one (A~38%, HC = 437 kA/m). The increase of HC with increasing Pt content was reported in the literature for the L10 FePt/fcc-FePt nanomagnets [19] and for ternary L10 (Fe,Co)Pt thin films [10]. Furthermore, the weakly magnetic FeB or non-magnetic (Fe,Co)B phases may have an influence on increasing the coercivity by acting as pinning centers. It was suggested for L10 (Fe,Co)Pt + fcc-(Fe,Co)Pt thin film structures that a separation of the grains with a very thin non-magnetic material could result in sufficient inter-granular exchange decoupling, sustaining the large coercivity [10]. The larger magnetization observed for the Fe60-yCoyPt25B15 alloys than for the Fe52-xCoxPt28B20 samples is attributed to strong exchange coupling interactions between the hard magnetic L10 FePt and (Fe,Co)Pt and the soft magnetic Fe2B or (Fe,Co)2B nanocrystalline phases. The Fe2B phase is characterized by a significantly larger magnetization than the FeB phase. Cobalt substitution for Fe in the (Fe,Co)2B phase causes a reduction of the magnetization, resulting in decreasing magnetic remanence of the Fe60-yCoyPt25B15 alloys with increasing Co content. Comparing the HC, Mr and (BH)max values of both series of the alloys studied it can be concluded that the magnetization has larger impact on the magnitude of the effective energy product than the coercivity.
4. CONCLUSIONS
Annealing of the rapidly quenched Fe52-xCo xPt28B20 and Fe60-yCoyPt25B15 alloys at temperatures exceeding the main exothermic peak in the DSC curves (840 and 880 K, respectively) resulted in the formation of different nanocomposite structures, depending on the alloy composition. In the Co-free alloys their nanostructures consisted of the magnetically hard ordered L10 tetragonal FePt and magnetically soft FeB (for x = 0) or Fe2B (for y = 0) nanocrystalline phases. A significantly larger relative fraction of the L1 0 FePt component was observed in the Mössbauer spectrum of the Fe52Pt28B20 alloy than that for the Fe60Pt25B15 one. In the case of the (Fe,Co)-Pt-B samples cobalt substituted for iron in the L1 0 FePt phase as well as in both borides. The L10 FePt and L1 0 (Fe,Co)Pt atomic arrangements were distinguished in the Mössbauer spectra. Iron-rich L1 0 FePt environments with a hyperfine field as large as 36 T were
12
observed for the Fe60-yCo yPt25B15 alloys with y = 0 and 10. Cobalt addition was found to decrease the hyperfine field of the (Fe,Co)2B phase as compared with the Fe2B phase. A significant reduction and eventually vanishing of the hyperfine field with increasing Co content was observed for the (Fe,Co)B phase, which became non-magnetic for the Fe52-xCoxPt28B20 alloys with x ≥ 30. The formation of phases with Curie temperatures exceeding 700 K was observed in the temperature dependence of magnetization of the as-quenched samples. Cobalt addition increased the Curie temperature of the L10 (Fe,Co)Pt phase. The room temperature values of the magnetic moment per formula unit calculated for both series of the as-quenched alloys decreased with increasing Co content. The Fe60-yCoyPt25B15 alloys exhibited significantly larger magnetic remanence and maximum energy products but a smaller coercivity (Mr = 0.63-0.87 T, Hc up to 439 kA/m, (BH)max up to 80 kJ/m3) than those observed for the Fe52-xCo xPt28B20 alloys (Mr = 0.53-0.66 T, Hc > 820 kA/m, (BH)max up to 65 kJ/m3). The best hard magnetic properties were obtained for both Co-free alloys. The largest magnetic remanence of 0.87 T and the highest maximum energy product (BH)max = 80 kJ/m3 were recorded for the Co-free Fe52Pt28B20 alloy. Co addition caused a reduction of the magnetization and the energy product in both series of the alloys. Differences in the hard magnetic properties of the nanocomposite alloys are related to different phase compositions in their nanostructures, which may enhance or hinder the interphase exchange coupling interactions. In the Fe52-xCoxPt28B20 alloys the hard magnetic properties origin mainly from the interactions between the ordered L10 FePt and (Fe,Co)Pt nanocrystallites, revealing a large magnetocrystalline anisotropy regardless of Co content. The weakly magnetic FeB or non-magnetic (Fe,Co)B phases may have an influence on increasing coercivity by acting as pinning centers. In the Fe60-yCoyPt25B15 alloys strong exchange coupling interactions between the hard magnetic L1 0 FePt and (Fe,Co)Pt phases and the soft magnetic Fe2B or (Fe,Co)2B phases were observed, which resulted in higher energy products than for the Fe52-xCoxPt28B20 alloys. It is concluded that in the (Fe,Co)-Pt-B alloys the magnetization of the soft phase plays a more important role for the energy product than the coercivity related to the magnetic anisotropy of the hard phase. It was shown that cobalt addition was not effective in the enhancement of the energy product. On the other hand, the relative contents of Pt and B in the alloys had a substantial influence on the effective hard magnetic properties of the nanocomposite (Fe,Co)-PtB alloys. The present results contribute to the knowledge related to tailoring the alloy chemical composition in order to improve hard magnetic properties of the Fe-Pt-B alloys, which are the rare-earth-free alternative for applications as permanent magnets.
13
References
[1] G. Varvaro, S. Laureti, D. Fiorani, J. Magn. Magn. Mater. 368 (2014) 415. [2] L. Suber, G. Marchegiani, E.S. Olivetti, F. Celegato, M. Coïsson, P. Tiberto, P. Allia, G. Barrera, L. Pilloni, L. Barba, F. Padella, P. Cossari, A. Chiolerio, Mater. Chem. Phys. 144 (2014) 186. [3] M. Pousthomis, C. Garnero, C.G. Marcelot, T. Blon, S. Cayez, C. Cassignol, V.A. Du, M. Krispin, R. Arenal, K. Soulantica, G. Viau, L.-M. Lacroix, J. Magn. Magn. Mater. 424 (2017) 304. [4] W. Zhang, D. V. Louzguine, A. Inoue, Appl. Phys. Lett. 85 (2004) 4998. [5] C. W. Chang, H. W. Chang, C. H. Chiu, W. C. Chang, S. K. Chen, A. C. Sun, J. Magn. Magn. Mater. 292 (2005) 120. [6] W. Zhang, P. Sharma, K. Shin, D. V. Louzguine, A. Inoue, Scr. Mater. 54 (2006) 431. [7] W. Zhang, K. Yubuta, P. Sharma, A. Inoue, J. Appl. Phys. 99 (2006) 08E914. [8] P.W. Jang, D.W. Kim, C.H. Park, J.G. Na, S.R. Lee, J. Appl. Phys. 83 (1998) 6614. [9] F.T. Yuan, S.N. Hsiao, W.M. Liao, S.K. Chen, Y.D. Yao, J. Appl. Phys. 99 (2006) 08E915. [10] Y. Liu, T.A. George, R. Skomski, D.J. Sellmyer, J. Appl. Phys. 111 (2012) 07B537. [11] R. Goyal, N. Sehdev, S. Lamba, S. Annapoorni, Solid State Commun. 226 (2016) 44. [12] A. Inoue, W. Zhang, J. Appl. Phys. 97 (2005) 10H308. [13] A. Makino, T. Bitoh, M. Nakagawa, J. Non-Cryst. Solids 353 (2007) 3655. [14] T. Goto, H. Utsugi, A. Kashiwakura, J. Magn. Magn. Mater. 104-107 (1992) 2051. [15] F. E. Spada, F. T. Parker, C. L. Platt, J. K. Howard, J. Appl. Phys. 94 (2003) 5123. [16] V. Raghavendra Reddy, S. Kavita, A. Gupta, J. Appl. Phys. 99 (2006) 113906. [17] S. Koyama, T. Goto, J. Magn. Magn. Mater. 321 (2009) 2407. [18] M. Rennhofer, M. Kozlowski, B. Laenens, B. Sepiol, R. Kozubski, D. Smeets, A. Vantomme, Intermetallics 18 (2010) 2069. [19] S. Srivastava, N.S. Gajbhiye, J. Magn. Magn. Mater. 401 (2016) 969. [20] A. Grabias, M. Kopcewicz, D. Oleszak, J. Latuch, M. Kowalczyk, M. Pękała, J. Phys.: Conf. Series 144 (2009) 012077. [21] A. Grabias, M. Kopcewicz, D. Oleszak, J. Latuch, J. Phys.: Conf. Series 217 (2010) 012075. [22] A. Grabias, M. Kopcewicz, D. Oleszak, J. Latuch, M. Kowalczyk, M. Pękała, J. Magn. Magn. Mater. 322 (2010) 3137. [23] N. Randrianantoandro, A.D. Crisan, O. Crisan et al., J. Appl. Phys. 108 (2010) 093910. [24] A. Grabias, J. Latuch, D. Oleszak, M. Kopcewicz, Acta Phys. Polonica A 119 (2011) 68.
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[25] R. A. Brand, J. Lauer, D. M. Herlach, J. Phys. F 13 (1983) 675. [26] F. H. Sanchez, J. I. Budnick, Y. D. Zhang, W. A. Hines, M. Choi, R. Hasegawa, Phys. Rev. B 34 (1986) 4738. [27] K. Barmak, J. Kim, L.H. Lewis, K.R. Coffey, M.F. Toney, A.J. Kellock, J.-U. Thiele, J. Appl. Phys. 98 (2005) 033904. [28] P.H. Lee, Z.R. Xiao, K.L. Chen, Y. Chen, S.W. Kao, T.S. Chin, Physica B 404 (2009) 1989. [29] Y.Q. Liu, X.S. Zhao, J. Yang, J.Y. Shen, J. Alloys Compd. 509 (2011) 4805.
15
Figure captions
Fig. 1. DSC curves recorded for Fe52-xCo xPt28B20 (x = 0 and 26) and Fe60-yCo yPt25B15 (y = 0 and 40) ribbons at 40 K/min. Fig. 2. XRD patterns of the alloys annealed at 840 K for 30 min.: (a-b) Fe52-xCo xPt28B20 (x = 0 and 26), and (c-d) Fe60-yCo yPt25B15 (y = 0 and 40). The straight vertical lines, marking some positions of the nominal diffraction peaks of the ordered fct-FePt phase, are for eye-guide only. Fig. 3. Mössbauer spectra of Fe52-xCoxPt28B20 (x = 0, 10, 18, 26) samples annealed at 840 K for 30 min. Positions of the subspectral lines related to the L1 0 FePt, L10 (Fe,Co)Pt, FeB and (Fe,Co)B phases are marked by S1, S2, S3a, and D1, respectively. Fig. 4. Mössbauer spectra of Fe60-yCoyPt25B15 (y = 0, 10, 30, 40) samples annealed at 880 K for 30 min. Positions of the subspectral lines related to the L1 0 FePt, L10 (Fe,Co)Pt, Fe2B, fcc-FePt and iron-rich FePt phases are marked by S1, S2, S3b, S4 and S5, respectively. Fig. 5. Hyperfine field values of the L10 FePt (S1), (Fe,Co)Pt (S2), (Fe,Co)B (S3a) and (Fe,Co)2B (S3b) phases as a function of Co content in the alloy. Fig. 6. Relative spectral fractions of the phases present in the annealed (a) Fe52-xCo xPt28B20 and (b) Fe60-yCo yPt25B15 alloys, determined from the Mössbauer spectra in Fig. 5 and Fig. 6, respectively, as a function of Co content in the alloy: S1 - L10 FePt, S2 - L10 (Fe,Co)Pt, S3a (Fe,Co)B, S3b - (Fe,Co)2B, S4 - fcc-FePt, S5 - iron-rich FePt. Fig. 7. Magnetization vs. temperature dependence of the as-quenched (a) Fe52-xCo xPt28B20 and (b) Fe60-yCo yPt25B15 alloys measured at 1.5 T. Fig. 8. Magnetic moment per formula unit, µB, as a function of Fe content in the Fe52-xCoxPt28B20 and Fe60-yCoyPt25B15 alloys. Fig. 9. Room temperature magnetic hysteresis loops recorded for the annealed (a) Fe52-xCoxPt28B20 and (b) Fe60-yCoyPt25B15 alloys.
16
Table 1.
Onset (Tx1, Tx2) and peak (Tp1, Tp2) crystallization temperatures of the as-quenched alloys determined from DSC measurements.
Sample
Tx1 (K)
Tp1 (K)
Tx2 (K)
Tp2 (K)
Fe52Pt28B20
648
688
788
801
Fe42Co10Pt28B20
653
700
808
823
Fe34Co18Pt28B20
679
705
811
825
Fe26Co26Pt28B20
686
705
798
810
Fe60Pt25B15
847
865
Fe50Co10Pt25B15
856
877
Fe30Co30Pt25B15
829
873
Fe20Co40Pt25B15
834
869
17
Table 2. Comparison of the average crystallite size of the L10 phase,, coercivity, HC, magnetic remanence, Mr, reduced remanence, Mr/Ms, and maximum energy product, (BH)max, of the annealed alloys.
Sample Fe52-xCoxPt28B20
Fe60-yCoyPt25B15
< D > (nm) HC (kA/m)
Mr (T)
Mr/Ms* (BH)max (kJ/m3)
x=0
20
823
0.66
0.76
65
x = 10
21
970
0.56
0.76
51
x = 18
24
953
0.54
0.76
52
x = 26
24
829
0.53
0.75
46
y=0
30
437
0.87
0.74
80
y = 10
31
438
0.80
0.71
69
y = 30
30
439
0.68
0.68
51
y = 40
32
360
0.63
0.67
47
*
Ms at 1.6 MA/m
18
19
20
21
22
23
24
25
26
27
•
Nanocomposite alloys were formed by annealing of the rapidly quenched alloys
•
Magnetically hard L1 0 (Fe,Co)Pt and soft (Fe,Co)2B or (Fe,Co)B were formed
•
Mössbauer spectra revealed Co substitution for Fe in L10 FePt, FeB and Fe2B phases
•
Annealed alloys exhibit hard magnetic properties which depend on phase compositions
•
Co addition was found to decrease the magnetization and the energy product
28
S0304-8853(16)33437-0 http://dx.doi.org/10.1016/j.jmmm.2017.03.047 MAGMA 62570
To appear in:
Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
22 December 2016 18 March 2017 24 March 2017
Please cite this article as: A. Grabias, M. Kopcewicz, J. Latuch, D. Oleszak, M. Pękała, M. Kowalczyk, Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys, Journal of Magnetism and Magnetic Materials (2017), doi: http://dx.doi.org/10.1016/j.jmmm.2017.03.047
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Influence of cobalt content on the structure and hard magnetic properties of nanocomposite (Fe,Co)-Pt-B alloys A. Grabiasa,*, M. Kopcewicza, J. Latuchb, D. Oleszakb, M. Pękałac, M. Kowalczykb
a
Institute of Electronic Materials Technology, Wólczyńska 133, 01-919 Warsaw, Poland
b
Faculty of Materials Science and Engineering, Warsaw University of Technology, Wołoska
141, 02-507 Warsaw, Poland c
Department of Chemistry, University of Warsaw, Al. Żwirki i Wigury 101, 02-089 Warsaw,
Poland
*Corresponding author: Agnieszka Grabias Institute of Electronic Materials Technology Wólczyńska 133 01-919 Warsaw, Poland Tel.: +48 22 6395547 Fax: +48 22 8645496 E-mail: [email protected]
1
Abstract
The influence of Co content on the structural and hard magnetic properties of two sets of nanocrystalline Fe52-xCoxPt28B20 (x = 0 − 26) and Fe60-yCoyPt25B15 (y = 0 − 40) alloys was studied. The alloys were prepared as ribbons by the rapid quenching technique. The nanocomposite structure in the alloys was obtained by annealing at 840-880 K for 30 min. Structural characterization of the samples was performed using the Mössbauer spectroscopy and X-ray diffraction. Magnetic properties of the samples were studied by the measurements of the hysteresis loops and of the magnetization at increasing temperatures. An amorphous phase prevailed in the as-quenched Fe52-xCoxPt28B20 alloys while a disordered solid solution of fcc-(Fe,Co)Pt was a dominating phase in the Fe60-yCo yPt25B15 ribbons. Differential scanning calorimetry measurements revealed one or two exothermic peaks at temperatures up to 993 K, depending on the composition of the alloys. Thermal treatment of the samples led to the formation of the magnetically hard ordered L10 tetragonal (Fe,Co)Pt nanocrystallites and magnetically softer phases of (Fe,Co)B (for Fe52-xCoxPt28B20) or (Fe,Co)2B (for Fe60-yCoyPt25B15). Detailed Mössbauer spectroscopy studies revealed that cobalt substituted for iron in both the L1 0 phase and in iron borides. The nanocomposite Fe60-yCo yPt25B15 alloys exhibited significantly larger magnetic remanence and maximum energy products but a smaller coercivity than those observed for the Fe52-xCoxPt28B20 alloys. Co addition caused a reduction of the magnetization and the energy product in both series of the alloys. The largest magnetic remanence of 0.87 T and the highest energy product (BH)max = 80 kJ/m3 were obtained for the Co-free Fe52Pt28B20 alloy while the largest coercivity (HC > 950 kA/m) was observed for the Fe50Co10Pt25B15 and Fe30Co30Pt25B15 alloys. Differences in the hard magnetic properties of the nanocomposite alloys were related to different phase compositions influencing the strength of inter-phase exchange coupling interactions.
Keywords: nanocomposite alloy; Fe-Pt-B alloy; hard magnetic properties; Mössbauer spectroscopy
2
1. INTRODUCTION
The ordered L1 0 face-centered-tetragonal FePt phase with a high magnetocrystalline anisotropy energy (K = 7 MJ/m3) has been widely used as a hard magnetic component in designing various structures in the form of thin films, multilayers, nanoparticles or nanocomposite alloys for applications in ultrahigh-density magnetic storage media or as permanent magnets [1-4]. It was shown that in Fe-Pt-B alloys a nanocomposite structure, consisting of magnetically hard, ordered L10 FePt phase and soft magnetic, disordered facecentered-cubic FePt and Fe2B phases, can be used for the preparation of exchange-coupled nanocrystalline magnets with a high energy product up to about 120 kJ/m3 [4-7]. Such nanocomposite structures can be formed directly by rapid quenching [7] or by annealing disordered alloys [4-6]. It is also known that the ordered L1 0 CoPt phase reveals a high magnetocrystalline anisotropy similar to that of the L1 0 FePt one. Therefore, in search of enhanced hard magnetic properties the preparation and magnetic properties of ternary (FeCo)Pt thin films have been studied in, e.g., [8-11]. Only several works devoted to Co-containing Fe-PtB alloys have been published [12,13]. In all cases it was found that Co substitutes for Fe in the ordered L10 FePt phase. In the Fe49-xCo xPt51 thin films the substitution of Fe with Co resulted in a higher magnetization and a lower coercivity as compared with the FePt film [9]. A refined grain size and a reduced long-range order parameter were observed in ternary L10 FeCoPt films [9]. In the nanocrystalline (Fe,Co)-Pt-B alloys with a fairly low Pt content the increase of Co addition caused a gradual increase of the coercivity, however, the magnetic remanence and the maximum energy product became smaller [12]. In the nanocomposite (Fe,Co)-Pt-Zr-B melt-spun alloys a monotonous decrease of the coercivity was observed with the increase of Co content, which was attributed to a reduction of the magnetocrystalline anisotropy originating from decreasing the long-range order parameter of the L1 0 phase [13]. The differences in hard magnetic properties of the Fe-Pt-B-based nanocomposite alloys strongly depend on their multi-phase nanostructure features, such as the chemical composition and ordering of the phases, their relative fraction, and the size of the nanograins. The aim of the present work was to investigate the influence of Co substitution for Fe on the formation of nanocomposite structures and on hard magnetic properties of two sets of (Fe,Co)-Pt-B alloys with different Pt and B contents. Therefore, two series of Fe52-xCoxPt28B20 (x = 0 − 26) and Fe60-yCoyPt25B15 (y = 0 − 40) disordered alloys were prepared by the melt spinning technique. The rapidly quenched alloys were annealed in order to obtain a nanocomposite structure with L10 FePt nanograins. The phase composition of the as-quenched and annealed samples was
3
probed by 57Fe Mössbauer spectroscopy measurements, which provided valuable information regarding local atomic arrangements in iron-containing phases as well as the magnetic behavior of iron atoms. This experimental method is sensitive enough to detect even small composition fluctuations in Fe environments. The Mössbauer spectroscopy has been successfully applied for studying magnetic systems of Fe-Pt [14-19] and Fe-Pt-B [20-23]. However, according to our best knowledge, the Mössbauer spectroscopy has not yet been used in studying Co-containing (Fe,Co)-Pt-B alloys, except for our recent publication [24]. In the present work, the Mössbauer spectroscopy proved to be very useful in identification of the phases and determination of their chemical compositions and relative fractions, thus enabling the study of the relations between the nanostructure and magnetic properties of the (Fe,Co)-Pt-B alloys. Understanding these relations is important for the application of the alloys as exchange-coupled permanent magnets.
2. EXPERIMENTAL DETAILS
The Fe52-xCo xPt28B20 (x = 0, 10, 18, 26 at.%) and Fe60-yCo yPt25B15 (y = 0, 10, 30, 40 at.%) alloys were prepared by melt spinning under Ar protective atmosphere. The linear velocity of 45 m/s at the spinning wheel was applied. The ribbons were about 1 mm wide and 15-20 µm thick. Differential scanning calorimetry (DSC) measurements were performed for the rapidly quenched alloys in the range of 300−993 K with the heating rate of 40 K/min. The samples were annealed at 840−880 K for 30 min. in vacuum. Structural properties of the as-quenched and annealed ribbons were investigated by the Mössbauer spectroscopy and X-ray diffraction (XRD) using CuKα radiation. All Mössbauer spectra were measured in transmission geometry at room temperature. A constant-acceleration spectrometer with a 57Co-in-Rh source with activity of about 25 mCi was used. Mössbauer spectra were fitted using the NORMOS program [25]. The spectra of the as-quenched alloys were fitted with a hyperfine field distribution P(Bhf) method. The spectra of the crystalline samples were analyzed using the least squares method, assuming the Lorentzian profile of lines. Determination of the hyperfine parameters, such as magnetic hyperfine field, Bhf, isomer shift, δ, and quadrupole shift, ∆, characteristic for magnetically split six-line spectral components or quadrupole splitting and isomer shift for doublets, enabled the identification of iron-containing magnetic and non-magnetic phases. Relative fractions of the spectral components related to the identified phases were calculated as a ratio of the area of the relevant subspectrum to the total spectral area, assuming similar Debye-Waller factors for each phase. Isomer shifts are related to α-Fe standard at room temperature. Temperature dependence of magnetization of the as-quenched samples was measured in the temperature range of 290-980
4
K under applied magnetic field of 1.5 T using a Faraday balance. A vibrating sample magnetometer (VSM) with maximum applied magnetic field of 2 T was used to investigate magnetic properties of the samples at room temperature.
3. RESULTS AND DISCUSSION 3.1. As-quenched alloys
Detailed XRD and Mössbauer results obtained for the as-quenched Fe52-xCoxPt28B20 (x = 0 - 26) and Fe60-yCo yPt25B1 5 (y = 0 - 40) alloys were presented in our previous work [24]. It was found that the rapidly quenched Fe52-xCoxPt28B20 alloys were predominantly amorphous with a small contribution of the disordered fcc-FePt phase. In the as-quenched Fe60-yCo yPt25B15 alloys the fcc-(Fe,Co)Pt disordered solid solution coexisted with amorphous (Fe,Co)-B regions. Thermal stability of the rapidly quenched (Fe,Co)-Pt-B alloys was studied by DSC measurements. Examples of DSC curves are shown in Fig. 1 for the extreme compositions studied: without Co and with the highest Co content. The characteristic temperatures, determined from DSC curves, are collected in Table 1 for all alloys studied. For the Fe52-xCoxPt28B20 alloys two separated exothermic peaks are observed whereas for the Fe60-yCo yPt25B15 samples a broad exothermic effect is recorded. In the case of predominantly amorphous Fe52-xCoxPt28B20 alloys a weak peak at lower temperatures is attributed to thermally induced atomic rearrangements in the amorphous phase, leading to a formation of the fcc-(Fe,Co)Pt disordered solid solution. This exothermic effect is more pronounced for larger Co content, suggesting larger amount of the amorphous phase in Co-containing alloys. As seen from Table 1, the onset temperature of this peak, Tx1, increases with increasing Co content (from 648 to 686 K for x = 0 and 26, respectively), indicating that Co addition increases thermal stability of the amorphous phase. Similar finding has been reported in [12]. Such an exothermic peak is not seen for the Fe60-yCoyPt25B15 alloys, in which the fcc-(Fe,Co)Pt disordered solid solution is already formed in the as-quenched state. The main exothermic effect, observed for all alloys studied, is attributed to the formation of iron and cobalt borides and the transformation of the fcc-(Fe,Co)Pt disordered solid solution to the ordered tetragonal phase as discussed in the next section. Onset temperatures of these phase transformations are in the range of 788−856 K and are significantly higher for the set of the Fe60-yCoyPt25B15 alloys than for the Fe52-xCoxPt28B20 ones.
5
3.2. Annealed alloys
In order to form a nanocomposite structure consisting of the hard and soft magnetic phases, the FeCoPtB alloys were annealed for 30 min. at temperatures corresponding to the main exothermic peak in the DSC curves (Fig. 1). In each case the selected annealing temperatures were slightly larger than Tp2 (Table 1), ensuring required phase transformations. The XRD patterns of the Fe52-xCo xPt28B20 (x = 0 and 26) and Fe60-yCo yPt25B15 (y = 0 and 40) alloys annealed for 30 min. at 840 K and 880 K, respectively, are presented in Fig.2. In each pattern two sets of diffraction lines related to two phases are identified. All XRD patterns contain diffraction lines characteristic of the ordered L10 fct-(Fe,Co)Pt phase. The presence of cobalt in the L1 0 phase is evidenced for Co-containing alloys by a slight shift towards larger 2θ angles of the diffraction lines characteristic of the ordered L10 FePt phase, which is best seen in the pattern obtained for Fe20Co40Pt25B15 (Fig. 2d, y = 40). It is worth noting that a formation of the ordered, ternary (Fe,Co)Pt phase instead of a mixture of the L10 CoPt and L10 FePt phases was evidenced in annealed (Co1-xFex)Pt thin films [8]. The average crystallite size of the L1 0 phase was determined by the Williamson-Hall method (Table 2). It increases from 20 to 24 nm with increasing Co content from x = 0 to x = 26. In the case of the Fe60-yCoyPt25B15 samples the average crystallite size is contained in the narrow range of 30−32 nm. The second set of less intense diffraction lines is attributed to the formation of borides: FeB (Fig. 2a) or (Fe,Co)B (Fig. 2b) for the Fe52-xCo xPt28B20 alloys, and Fe2B (Fig. 2c) or (Fe,Co)2B (Fig. 2d) for the Fe60-yCoyPt25B15 alloys. In the case of iron borides Co substitution for Fe also causes a shift of diffraction lines towards larger 2θ angles. Mössbauer spectra recorded for the Fe52-xCoxPt28B20 alloys annealed at 840 K are shown in Fig. 3. The Mössbauer spectrum of the Co-free alloy is fitted with two magnetic components (Fig. 3a). The S1 sextet with hyperfine field Bhf = 27.8 T, isomer shift δ = +0.27 mm/s and quadrupole shift ∆ = 0.30 mm/s corresponds to the tetragonal ordered L10 FePt phase [14,22]. The S3a sextet with Bhf = 10.8 T, δ = +0.26 mm/s and ∆ = 0.12 mm/s is assigned to the crystalline FeB phase [26]. The relative fractions of these two sextets (57% and 43% for S1 and S3a, respectively), calculated from the Mössbauer spectrum, are directly related to two structurally different positions of iron nuclei in the two-phase sample, and indicate that both phases have nearly equiatomic compositions with a slightly larger Fe content as compared with the stoichiometric L1 0 FePt and FeB phases. Incorporation of 10% of cobalt causes the appearance of two additional subspectra (Fig. 3b). Beside the S1 sextet related to the L1 0 FePt phase the S2 sextet with Bhf = 28.6 T, δ = +0.27 mm/s and ∆ = 0.17 mm/s is observed. In the
6
central part of the spectrum a clear quadrupole doublet D1 with quadrupole splitting of 0.60 mm/s and isomer shift δ = +0.26 mm/s appears. With increasing Co content in the alloy the relative spectral fraction of the S2 sextet increases at the expense of the S1 one (Fig. 3b-3d). The hyperfine parameters of the S2 magnetic component strongly suggest a formation of the ternary tetragonal (Fe,Co)Pt phase. The value of the quadrupole shift obtained for the S2 sextet (∆S2) is large but significantly smaller as compared with ∆S1 = 0.30 mm/s observed for the ordered L1 0 FePt phase, which provides an evidence of somewhat less distorted local atomic symmetry in the tetragonal (Fe,Co)Pt phase than in the L1 0 FePt one. Furthermore, the ∆S2 parameter decreases with increasing Co content from 0.17 to 0.11 mm/s for x = 10 and 26, respectively. It was shown experimentally that Co substitution for Fe in the L10 FePt phase causes a decrease of the ordering and the lattice parameters ratio c/a becomes closer to 1 as compared with the data of the ordered binary FePt phase with the same Pt content [8,9,27]. For larger Co content the S3a sextet, assigned to the magnetic FeB phase, is not observed in the spectra (Fig. 3c, 3d). The appearance of the quadrupole doublet D1 in the place of the S3a sextet is related to the incorporation of Co into the FeB lattice. A presence of Co in this iron boride is already evidenced for x = 10 by a reduction of the hyperfine field of the S3a sextet to about 8 T (Fig. 3b). Further substitution of Fe with Co in the (Fe,Co)B compound leads to the formation of the nonmagnetic phase as seen by the appearance of the quadrupole doublet. It was reported for the orthorhombic Fe1-xCoxB system that the magnetic moment decreases almost linearly with the increase of Co content and then it drops to zero for x ≥ 0.85 [28]. Changes in the relative fractions of all spectral components with increasing Co content in the alloys studied are discussed in detail below (Fig. 6). Mössbauer spectra recorded for the Fe60-yCoyPt25B15 alloys annealed at 880 K are shown in Fig. 4. The spectrum of the Fe60Pt25B15 alloy (Fig. 4a) is more complex than that of the Fe52Pt28B20 alloy (Fig. 3a). The spectrum in Fig. 4a is fitted with 4 magnetic components. The sextet S3b with Bhf = 23.7 T, δ = +0.11 mm/s and ∆ = 0.00 mm/s constitutes the prevailing component, which is characteristic of the tetragonal Fe2B compound [26]. It is worth noting that for the tetragonal Fe2B phase two slightly different Fe positions are observed, therefore, two sextets with Bhf of 23.2 and 24.2 T are usually fitted in the Mössbauer spectra [26]. Both sextets reveal, however, almost identical values of the isomer shift and the quadrupole shift (δ = +0.12 mm/s, ∆ = 0.02-0.04 mm/s) [26]. Therefore, in the case of the complex spectra in Fig. 4 one sextet with slightly broadened lines is fitted for Fe2B. The observed hyperfine field of the S3b component is the mean value. According to the fitting of the Mössbauer spectrum about 54% of Fe atoms are located in iron boride, which indicates that the Fe2B phase is nearly stoichiometric.
7
The S1 sextet with Bhf = 28.1 T, δ = +0.27 mm/s and ∆ = 0.30 mm/s corresponds, as discussed above (Fig. 3), to the tetragonal ordered L10 FePt phase. Beside these two components with major spectral fractions, two minor sextets are fitted. The S4 sextet with Bhf = 29.1 T, δ = +0.20 mm/s and ∆ = 0.00 mm/s is attributed to the remaining fcc-FePt phase [15,22]. The relative spectral fraction of this component is below 8%. The amount of the fcc phase, detected in the Mössbauer spectrum, is too small to be observed in the XRD pattern. A small spectral contribution (~4%) of the S5 sextet with a large hyperfine field of about 36 T, δ = +0.25 mm/s and ∆ = -0.31 mm/s is also detected in the spectrum in Fig. 4a. A very similar additional spectral component was observed for the binary Fe61Pt39 L1 0 ordered alloy [17]. It was attributed to Fe atoms on Pt-sites in the L10 structure. Therefore, the S5 sextet is assigned to iron-rich FePt environments in the L10 FePt phase. The origin of this component results most probably from the excess of iron content in the alloy. After incorporation of 10% of Co to the Fe-Pt-B alloy the shape of the Mössbauer spectrum changes significantly (Fig. 4b). It is fitted with 4 subspectra: the previously observed sextets S3b (related to Fe2B) and S1 and S5 (both related to the L10 phase), and a new S2 component. The S3b sextet reveals a slightly reduced mean hyperfine field of 23.1 T as compared with 23.7 T obtained for the Co-free alloy. This reduction is related to Co presence in the iron boride. The S2 sextet with Bhf = 28.8 T, δ = +0.25 mm/s and ∆ = 0.16 mm/s is assigned to the formation of ternary (Fe,Co)Pt phase with the ordered L1 0 structure. A similar spectral component (S2) with very close hyperfine parameters appears in the Mössbauer spectra of the annealed Fe52-xCo xPt28B20 alloys for x≥10 (Fig. 3). As can be seen from Fig. 4b-4d the intensity of the S2 sextet increases with increasing Co content. For cobalt concentrations exceeding 10% the S5 sextet, originating from the excess of iron atoms, disappears (Fig. 4c, 4d). Such observation in the Mössbauer spectra is obviously related to the reduction of Fe content in the alloys and the substitution of Fe atoms with Co atoms in the L10 phase. Therefore, the spectra of the samples with y = 30 and 40 are fitted with only 3 spectral components, originating from the ordered L10 FePt and (Fe,Co)Pt (S1 and S2, respectively) and (Fe,Co)2B (S3b) phases (Fig. 4c, 4d). The influence of Co addition on the Mössbauer spectra of the annealed Fe52-xCo xPt28B20 and Fe60-yCoyPt25B15 alloys is demonstrated in Fig. 5 and Fig. 6, in which hyperfine field values and relative fractions of each spectral component are plotted vs. Co content, respectively. As can be seen from Fig. 5 the hyperfine fields of the L10 FePt (S1) and (Fe,Co)Pt (S2) phases do not change substantially with increasing Co content and are in the narrow ranges of 27.3−28.1 T and 28.4−28.8 T, respectively. In the case of the Fe52-xCo xPt28B20 samples a significant reduction by
8
about 3 T is observed for the hyperfine field related to the FeB phase (S3a) already for the addition of 10% of Co to the alloy (Fig. 5a). Further increase of cobalt content causes further substitution of Co for Fe in the (Fe,Co)B phase and the hyperfine field vanishes for x>10 (Fig. 3). A continuous decrease of the hyperfine field with increasing Co content is also observed for the (Fe,Co)2B phase (S3b, Fig. 5b). This reduction is directly related to a gradual substitution of Co for Fe in this phase. This is confirmed by thermodynamic calculations of ternary Fe-Co-B system revealing that Fe2B and Co2B form continuous solid solutions of (Fe,Co)2B [29]. The formation of continuous solid solutions of (Fe,Co)B by FeB and CoB was also demonstrated in the same work [29]. It is worth noting that we measured the Mössbauer spectrum of a crystalline sample with the composition of (Fe0.7Co0.3)2B, which exhibited the mean hyperfine field of 22.8 T (about 1 T lower than that of Fe2B). This value is only slightly smaller than Bhf = 23.1 T obtained for the sample with y = 10 (Fig. 5b), which indicates that the composition of the (Fe,Co)2B phase is close to that of about (Fe0.7Co0.3)2B. This is confirmed by the estimation based on the relative fraction of iron atoms contributing to the S3b spectral component, which gives the composition of (Fe0.74Co0.26)2B. For the alloys with y = 30 and 40 Co concentration in this phase prevails over Fe content as evidenced by a significant reduction of the hyperfine field to 21.6 and 19.7 T, respectively (Fig. 5b). The estimated composition is (Fe0.1Co0.9)2B for y = 40. Dependence of the relative fractions of all spectral components on the Co content in the alloys is presented in Fig. 6. In the case of the Fe52-xCoxPt28B20 alloys (Fig. 6a) the increase of Co content causes a significant decrease of the relative fractions of the L10 FePt (S1) and (Fe,Co)B (S3a+D1) components whereas the fraction of L1 0 (Fe,Co)Pt (S2) increases substantially. The decrease of the fraction of the (Fe,Co)B component (S3a+D1) is directly related to the decrease of iron content in this phase when Fe is substituted with Co. It is worth noting that Fe-free CoB or Co 2B phases would not be seen in Mössbauer spectra. The relative spectral fractions of the phases, calculated for the Fe60-yCoyPt25B15 alloys, are plotted in Fig. 6b. Comparing the results obtained for both series of the alloys clear differences in the relative fractions between the L10 FePt (S1) and Fe-B (S3) components are seen for both Co-free alloys (x = 0, y = 0). The relative fraction of L1 0 FePt (S1) is about 1.5 times larger for the Fe52Pt28B20 alloy than for Fe60Pt25B15. The relative fraction of FeB (S3a, Fig. 6a) is by about 10% smaller than that of Fe2B (S3b, Fig. 6b). Furthermore, unlike in the case of the Fe52Pt28B20 alloy a small abundance of about 7% of the fcc-FePt phase (S4) remained in the Fe60Pt25B15 sample after annealing (Fig. 6b), which also weighed on a reduced fraction of the ordered L10 phase. It is worth noting that the relative fractions of the S1, S2 and S3 components are quite similar for both alloys with the highest Co content (x = 26, y = 40). Thermally induced formation of the L10 (Fe,Co)Pt phase (S2) reached
9
the relative fraction of about 61% and 57% for the Fe26Co26Pt28B20 and Fe20Co40Pt25B15 alloys, respectively.
3.3. Magnetic measurements
Temperature variations of magnetization in the range from room temperature up to 980 K, depicted in Fig. 7, show that the (Fe,Co)-Pt-B systems studied are ferromagnetically ordered up to high temperatures. The as-quenched Fe52-xCoxPt28B20 alloys with the prevailing amorphous phase reveal Curie temperatures, TC, in the range of 655 − 665 K (Fig. 7a). Subsequent decrease of the magnetization up to about 700 -720 K results from the paramagnetic amorphous phase and from the formed fcc disordered solid solution (the first DSC peak in Fig. 1). Then the crystallization process of a ferromagnetic phase with higher TC is initiated as indicated by the increase of the magnetization. This magnetization increase is markedly larger for x = 18 and 26 than for x = 0 and 10. This enhancement of the magnetization is related to the formation of the ordered L10 FePt phase with TC of about 750 K for x = 0 and the L1 0 (Fe,Co)Pt phase for x ≥ 10. The Curie temperature of the latter phase increases with the increase of Co content in the alloy (TC ≈ 825 K for x = 18). It is worth noting that the orthorhombic FeB phase has TC of about 600 K. Since the crystallization of this iron boride starts at temperatures exceeding TC, the formation of FeB is not seen in the M(T) dependence. The temperature variation of magnetization observed for the Fe60-yCoyPt25B15 alloys (Fig. 7b) indicates a multicomponent behavior, originating from the inhomogeneity of the as-quenched samples as revealed by Mössbauer spectroscopy results [24]. The Curie temperature of the as-quenched samples with y = 0, 10 and 30 increases with increasing Co content from about 740 K to about 760 K. The magnetization enhancement observed for these alloys at about 770 K is related to the crystallization of the (Fe,Co)2B phase. Somewhat different magnetization behavior is observed for the as-quenched Fe20Co 40Pt25B15 alloy (y = 40) with a quite large content of the fcc solid solution, for which TC reaches 860 K. In this case only a small magnetization increase is seen at about 780 K, which corresponds to the crystallization of the boride phase as observed in other samples of this series. The M(T) plots above 770 K enable to locate TC of the high temperature phase of (Fe,Co)2B between 960 and 980 K, depending on the chemical composition. These values are smaller than the nominal Curie temperature of the Fe2B phase (1015 K). The calculated room temperature magnetic moment per formula unit is plotted in Fig. 8 for both series of the as-quenched alloys. The dependence of the magnetic moment on Co content in the alloy represents a monotonically decreasing function with increasing cobalt
10
content for both series. In the case of the as-quenched Fe52-xCo xPt28B20 alloys the magnetic moment decreases from 1.09 to 0.84 µB with the increase of Co content from x = 0 to x = 26. For the Fe60-yCo yPt25B15 series the values of magnetic moment are larger when comparing the alloys with the same Co content in both series, and decrease from 1.29 to 0.99 µB with the increase of Co content from y = 0 to y = 40. Therefore, in the (Fe,Co)-Pt-B alloys studied the substitution of iron atoms with cobalt atoms causes a decrease of the value of the magnetic moment per formula unit at room temperature. Room temperature hysteresis loops of the annealed alloys are presented in Fig. 9. The determined magnetic properties are collected in Table 2. All loops are smooth like in the case of a single-phase material, indicating inter-phase exchange-coupled interactions. The shape of the hysteresis loops shows that they approach saturation in magnetic fields somewhat larger than the applied field. The hysteresis loops recorded for the Fe52-xCoxPt28B20 samples exhibit good squareness and good hard magnetic properties (Fig. 9a). The coercivity exceeds 800 kA/m for all compositions studied (x) and reaches the largest values of 970 and 953 kA/m for the alloys with x = 10 and x = 18, respectively (Table 2). For these alloys quite similar values of 51-52 kJ/m3 of the maximum energy product, (BH)max, were obtained. The influence of Co substitution for Fe is seen in a reduction of the magnetization and of the energy product. Both the magnetic remanence (Mr) and the magnetization at 1.6 MA/m decrease with increasing Co content in the alloy (Fig. 9a, Table 2). The reduced remanence Mr/Ms is quite large of about 0.75-0.76 for all alloy compositions. The best hard magnetic properties in this series of the alloys were, however, obtained for the Co-free sample, for which (BH)max of about 65 kJ/m3 was observed. The hysteresis loops recorded for the Fe60-yCoyPt25B15 series reveal larger magnetization values but significantly smaller coercivity (up to 439 kA/m) and the reduced remanence (0.67-0.74) as compared with the Fe52-xCoxPt28B20 alloys (Fig. 9b, Table 2). The coercivity seems to be independent of the alloy composition for y ≤ 30. Also in this case the magnetic remanence, the saturation magnetization and the energy product decrease with increasing Co content. As in the previous case (Fig. 9a) the best hard magnetic properties were obtained for the Co-free Fe60Pt25B15 alloy with HC = 437 kA/m, Mr = 0.87 T, Mr/Ms = 0.74 and (BH)max of about 80 kJ/m3. The coercivity of the annealed Fe52-xCoxPt28B20 alloys is more than twice larger than that observed for the annealed Fe60-yCoyPt25B15 samples (Table 2). In the literature the largest coercivity was reported for the Fe49Pt33B18 alloy (HC > 955 kA/m, Mr = 0.66 T, (BH)max of about 71 kJ/m3) [5]. For the Fe52Pt30B18 alloy the values of HC = 783 kA/m, Mr = 0.70 T and (BH)max = 88 kJ/m3 were given [7]. The structure of both alloys consisted of the nanosized L10 FePt and
11
Fe2B phases. The large coercivity is attributed to a quite large Pt content in the Fe-Pt-B alloys, which takes part in the formation of the L1 0 FePt phase revealing a high magnetocrystalline anisotropy. As shown by the Mössbauer measurements in the present work (Fig. 6) a significantly larger relative fraction (A) of the spectral component assigned to the hard magnetic L10 FePt phase and a larger coercivity were observed for the Fe52Pt28B20 alloy with larger Pt content (A~57%, HC = 823 kA/m) than for the Fe60Pt25B15 one (A~38%, HC = 437 kA/m). The increase of HC with increasing Pt content was reported in the literature for the L10 FePt/fcc-FePt nanomagnets [19] and for ternary L10 (Fe,Co)Pt thin films [10]. Furthermore, the weakly magnetic FeB or non-magnetic (Fe,Co)B phases may have an influence on increasing the coercivity by acting as pinning centers. It was suggested for L10 (Fe,Co)Pt + fcc-(Fe,Co)Pt thin film structures that a separation of the grains with a very thin non-magnetic material could result in sufficient inter-granular exchange decoupling, sustaining the large coercivity [10]. The larger magnetization observed for the Fe60-yCoyPt25B15 alloys than for the Fe52-xCoxPt28B20 samples is attributed to strong exchange coupling interactions between the hard magnetic L10 FePt and (Fe,Co)Pt and the soft magnetic Fe2B or (Fe,Co)2B nanocrystalline phases. The Fe2B phase is characterized by a significantly larger magnetization than the FeB phase. Cobalt substitution for Fe in the (Fe,Co)2B phase causes a reduction of the magnetization, resulting in decreasing magnetic remanence of the Fe60-yCoyPt25B15 alloys with increasing Co content. Comparing the HC, Mr and (BH)max values of both series of the alloys studied it can be concluded that the magnetization has larger impact on the magnitude of the effective energy product than the coercivity.
4. CONCLUSIONS
Annealing of the rapidly quenched Fe52-xCo xPt28B20 and Fe60-yCoyPt25B15 alloys at temperatures exceeding the main exothermic peak in the DSC curves (840 and 880 K, respectively) resulted in the formation of different nanocomposite structures, depending on the alloy composition. In the Co-free alloys their nanostructures consisted of the magnetically hard ordered L10 tetragonal FePt and magnetically soft FeB (for x = 0) or Fe2B (for y = 0) nanocrystalline phases. A significantly larger relative fraction of the L1 0 FePt component was observed in the Mössbauer spectrum of the Fe52Pt28B20 alloy than that for the Fe60Pt25B15 one. In the case of the (Fe,Co)-Pt-B samples cobalt substituted for iron in the L1 0 FePt phase as well as in both borides. The L10 FePt and L1 0 (Fe,Co)Pt atomic arrangements were distinguished in the Mössbauer spectra. Iron-rich L1 0 FePt environments with a hyperfine field as large as 36 T were
12
observed for the Fe60-yCo yPt25B15 alloys with y = 0 and 10. Cobalt addition was found to decrease the hyperfine field of the (Fe,Co)2B phase as compared with the Fe2B phase. A significant reduction and eventually vanishing of the hyperfine field with increasing Co content was observed for the (Fe,Co)B phase, which became non-magnetic for the Fe52-xCoxPt28B20 alloys with x ≥ 30. The formation of phases with Curie temperatures exceeding 700 K was observed in the temperature dependence of magnetization of the as-quenched samples. Cobalt addition increased the Curie temperature of the L10 (Fe,Co)Pt phase. The room temperature values of the magnetic moment per formula unit calculated for both series of the as-quenched alloys decreased with increasing Co content. The Fe60-yCoyPt25B15 alloys exhibited significantly larger magnetic remanence and maximum energy products but a smaller coercivity (Mr = 0.63-0.87 T, Hc up to 439 kA/m, (BH)max up to 80 kJ/m3) than those observed for the Fe52-xCo xPt28B20 alloys (Mr = 0.53-0.66 T, Hc > 820 kA/m, (BH)max up to 65 kJ/m3). The best hard magnetic properties were obtained for both Co-free alloys. The largest magnetic remanence of 0.87 T and the highest maximum energy product (BH)max = 80 kJ/m3 were recorded for the Co-free Fe52Pt28B20 alloy. Co addition caused a reduction of the magnetization and the energy product in both series of the alloys. Differences in the hard magnetic properties of the nanocomposite alloys are related to different phase compositions in their nanostructures, which may enhance or hinder the interphase exchange coupling interactions. In the Fe52-xCoxPt28B20 alloys the hard magnetic properties origin mainly from the interactions between the ordered L10 FePt and (Fe,Co)Pt nanocrystallites, revealing a large magnetocrystalline anisotropy regardless of Co content. The weakly magnetic FeB or non-magnetic (Fe,Co)B phases may have an influence on increasing coercivity by acting as pinning centers. In the Fe60-yCoyPt25B15 alloys strong exchange coupling interactions between the hard magnetic L1 0 FePt and (Fe,Co)Pt phases and the soft magnetic Fe2B or (Fe,Co)2B phases were observed, which resulted in higher energy products than for the Fe52-xCoxPt28B20 alloys. It is concluded that in the (Fe,Co)-Pt-B alloys the magnetization of the soft phase plays a more important role for the energy product than the coercivity related to the magnetic anisotropy of the hard phase. It was shown that cobalt addition was not effective in the enhancement of the energy product. On the other hand, the relative contents of Pt and B in the alloys had a substantial influence on the effective hard magnetic properties of the nanocomposite (Fe,Co)-PtB alloys. The present results contribute to the knowledge related to tailoring the alloy chemical composition in order to improve hard magnetic properties of the Fe-Pt-B alloys, which are the rare-earth-free alternative for applications as permanent magnets.
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Figure captions
Fig. 1. DSC curves recorded for Fe52-xCo xPt28B20 (x = 0 and 26) and Fe60-yCo yPt25B15 (y = 0 and 40) ribbons at 40 K/min. Fig. 2. XRD patterns of the alloys annealed at 840 K for 30 min.: (a-b) Fe52-xCo xPt28B20 (x = 0 and 26), and (c-d) Fe60-yCo yPt25B15 (y = 0 and 40). The straight vertical lines, marking some positions of the nominal diffraction peaks of the ordered fct-FePt phase, are for eye-guide only. Fig. 3. Mössbauer spectra of Fe52-xCoxPt28B20 (x = 0, 10, 18, 26) samples annealed at 840 K for 30 min. Positions of the subspectral lines related to the L1 0 FePt, L10 (Fe,Co)Pt, FeB and (Fe,Co)B phases are marked by S1, S2, S3a, and D1, respectively. Fig. 4. Mössbauer spectra of Fe60-yCoyPt25B15 (y = 0, 10, 30, 40) samples annealed at 880 K for 30 min. Positions of the subspectral lines related to the L1 0 FePt, L10 (Fe,Co)Pt, Fe2B, fcc-FePt and iron-rich FePt phases are marked by S1, S2, S3b, S4 and S5, respectively. Fig. 5. Hyperfine field values of the L10 FePt (S1), (Fe,Co)Pt (S2), (Fe,Co)B (S3a) and (Fe,Co)2B (S3b) phases as a function of Co content in the alloy. Fig. 6. Relative spectral fractions of the phases present in the annealed (a) Fe52-xCo xPt28B20 and (b) Fe60-yCo yPt25B15 alloys, determined from the Mössbauer spectra in Fig. 5 and Fig. 6, respectively, as a function of Co content in the alloy: S1 - L10 FePt, S2 - L10 (Fe,Co)Pt, S3a (Fe,Co)B, S3b - (Fe,Co)2B, S4 - fcc-FePt, S5 - iron-rich FePt. Fig. 7. Magnetization vs. temperature dependence of the as-quenched (a) Fe52-xCo xPt28B20 and (b) Fe60-yCo yPt25B15 alloys measured at 1.5 T. Fig. 8. Magnetic moment per formula unit, µB, as a function of Fe content in the Fe52-xCoxPt28B20 and Fe60-yCoyPt25B15 alloys. Fig. 9. Room temperature magnetic hysteresis loops recorded for the annealed (a) Fe52-xCoxPt28B20 and (b) Fe60-yCoyPt25B15 alloys.
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Table 1.
Onset (Tx1, Tx2) and peak (Tp1, Tp2) crystallization temperatures of the as-quenched alloys determined from DSC measurements.
Sample
Tx1 (K)
Tp1 (K)
Tx2 (K)
Tp2 (K)
Fe52Pt28B20
648
688
788
801
Fe42Co10Pt28B20
653
700
808
823
Fe34Co18Pt28B20
679
705
811
825
Fe26Co26Pt28B20
686
705
798
810
Fe60Pt25B15
847
865
Fe50Co10Pt25B15
856
877
Fe30Co30Pt25B15
829
873
Fe20Co40Pt25B15
834
869
17
Table 2. Comparison of the average crystallite size of the L10 phase,
Sample Fe52-xCoxPt28B20
Fe60-yCoyPt25B15
< D > (nm) HC (kA/m)
Mr (T)
Mr/Ms* (BH)max (kJ/m3)
x=0
20
823
0.66
0.76
65
x = 10
21
970
0.56
0.76
51
x = 18
24
953
0.54
0.76
52
x = 26
24
829
0.53
0.75
46
y=0
30
437
0.87
0.74
80
y = 10
31
438
0.80
0.71
69
y = 30
30
439
0.68
0.68
51
y = 40
32
360
0.63
0.67
47
*
Ms at 1.6 MA/m
18
19
20
21
22
23
24
25
26
27
•
Nanocomposite alloys were formed by annealing of the rapidly quenched alloys
•
Magnetically hard L1 0 (Fe,Co)Pt and soft (Fe,Co)2B or (Fe,Co)B were formed
•
Mössbauer spectra revealed Co substitution for Fe in L10 FePt, FeB and Fe2B phases
•
Annealed alloys exhibit hard magnetic properties which depend on phase compositions
•
Co addition was found to decrease the magnetization and the energy product
28