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aPhys. Status Solidi C 7, No. 5, 1331–1335 (2010) / DOI 10.1002/pssc.200983375

FeCoBSiNb bulk metallic glasses with Cu additions Mihai Stoica*,1, Stefan Roth2, Jürgen Eckert1,**, Trisha Karan1,2, Shanker Ram2, Gavin Vaughan3, and Alain Reza Yavari4 1 IFW Dresden, Helmholtzstr. 20, 01069 Dresden, Germany 2 IIT Kharagpur, West Bengal 721302, India 3 European Synchrotron Radiation Facilities ESRF, 38042 Grenoble, France 4 SIMAP, INP Grenoble, BP 75, Saint Martin d’Hères Campus 38402, France

Received 22 October 2009, revised 2 December 2009, accepted 3 December 2009 Published online 19 April 2010

Keywords Fe-alloys, glasses, structure, annealing, phase transition, magnetic properties * Corresponding author: e-mail m.stoica@ifw-dresden.de Phone: +49 351 4659 644, Fax: +49 351 4659 452 **also at University of Technology Dresden, Germany

The effect of Cu additions (1, 2 and 3 at.%) on the glass forming ability (GFA) and magnetic properties of [(Fe0.5Co0.5)0.75Si0.05B0.20]96Nb4 is studied. As the Cu con-tent increases, the GFA deteriorates. The structure of the samples, as well as the crystallization behavior was in-situ studied by means of synchrotron beam. The as-cast 1 mm diameter rods are characterized by a mixture of (nano)-crystals and amorphous matrix. Upon heating, the pre-existing nuclei grow, without forming new crystalline phases. The magnetic properties are characteristic for soft

magnetic materials. Despite that several crystalline mag-netic phases can be identified in the as-cast samples, only one single magnetic transition prior to complete crystalli-zation could be identified upon heating. The values for saturation magnetization and for Curie temperature are relatively close for all samples. The initial permeability of the Cu containing samples seems to be higher than that for the fully amorphous [(Fe0.5Co0.5)0.75Si0.05B0.20]96Nb4 alloy.

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1 Introduction Since the first ferromagnetic metallic glass Fe-C-P was found in 1967 [1], Fe- and Co-based amorphous alloys and the resulting crystalline alloys, pro-duced through crystallization of the corresponding glassy precursors, are regarded as attractive industrial alloys be-cause of their excellent soft magnetic properties [2]. Based on their magnetic advantages, many kinds of products con-sisting of ferromagnetic metallic glasses have been widely used, such as anti-theft labels in supermarkets and libraries and high efficient magnetic transformers [3, 4]. Multi-component glassy alloys decrease the critical cooling rate of glass formation, and promote successfully the formation of bulk metallic glasses (BMGs), which give ferromagnetic metallic glasses potential application as advanced struc-tural material because of their high strength and good cor-rosion resistance [5-8]. Recently, (Fe-Co)-Si-B-Nb BMGs

are regarded as one of the most excellent candidates combining the advantages of functional and structural ma-terials because of their high glass-forming ability (GFA), good mechanical and magnetic properties [9]. Minor addi-tion of Cu in the corresponding glassy precursors is well known as a good method to increase the soft magnetic properties resulting in ultrafine crystalline alloys, such as in case of FINEMET alloys [10]. In the (Fe-Co)-B-Si-Nb alloy system, some researchers tried to elucidate the influence of minor addition of Cu on GFA, mechani- cal and magnetic properties of the resulting alloys. Jia et al. reported on the GFA of {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 0, 0.5, 0.6, 0.7 and 1.0) alloys [11]. Shen et al. revealed that in situ formation of (Fe,Co) and (Fe,Co)23B6 microcrystalline grains during the solidification

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process can improve the ductility of the {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}99.75Cu0.25 BMG com-posite [12]. Inoue et al. reported on the effect of crystalli-zation of Fe–Co–B–Si–Nb–Cu glassy alloys on their soft magnetic properties by heat treatment [13]. Li et al., in the case of {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}98Cu2 amor-phous ribbons, found out that during the solidification process the primary crystalline phase is the o-(Fe,Co)3B metastable phase, which is replaced by α-(Fe,Co), (Fe,Co)23B6 and (Fe,Co)2B phases under slower cooling conditions [14]. The precipitation of α-(Fe,Co) is benefi-cial for the improvement of soft magnetic properties of as-cast rods.

In this paper, we report on the effect of Cu addition in {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 1, 2 and 3) alloys on GFA and the evolution of crystalline phases in 1 mm diameter rods. The relationship between the phase constituents and the magnetic properties is also studied.

2 Experimental procedure Alloy ingots of

{[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 1, 2 and 3) were produced by induction-melting mixtures of pure Fe (99.9 mass %), Co (99.9 mass %), Si (99.99 mass %), B (99.9mass %), 25Fe75Nb eutectic pre-alloy (99.9 mass %) and Cu (99.9 mass %) in a quartz cup under a high-purity argon atmosphere. Pieces of the alloyed ingots were remelted in quartz tubes and the molten liquid was then in-jected into a water-cooled copper mold in a high-purity ar-gon atmosphere to produce rod-shaped specimens with 1 mm diameter. The melting temperature of the alloys was measured by an infrared pyrometer. The structure of the as-cast rods, as well as the in-situ crystallization behavior was examined by x-ray diffraction in transmission configu-ration, using a high-energy high intensity monochromatic synchrotron radiation (λ = 0.017615 nm) at ID11 of ESRF Grenoble. The heating in the beam was possible with the samples closed under vacuum in a capillary tube and placed in a LINKAM hot stage device. The heating was performed from room temperature up to 1100 K, at a con-stant heating rate of 0.33 K/s. The thermal stability and the melting behavior of the specimens were evaluated with a NETZSCH DSC 404 C differential scanning calorimeter (DSC) at heating and cooling rates of 0.33 K/s. For mag-netic measurements, M-H hysteresis loops were measured with a vibrating sample magnetometer (VSM) at ambient temperature. The Curie temperature (Tc) for the rod sam-ples was determined using an in-house developed Faraday magnetometer. In order to minimize the errors, the meas-ured data were analyzing using the method proposed by Herzer [15]. The coercivity was measured using a Foerster Coercimat under an applied field high enough to saturate the samples. All magnetic properties were measured under DC magnetic field.

3 Results and discussions Cu additions to the

master alloy [(Fe0.5Co0.5)0.75Si0.05B0.20]96Nb4 (for simplicity,

this alloy will be further named M.A.) decrease the GFA. From the previous published papers [11-14] it clearly re-sults that for more than 0.5 at% Cu only fully amorphous ribbons can be produced. In this light, we tried to study (nano)composite samples with a small crystalline fraction, i.e. thicker than the ribbons but still far from the maximum (5 or even 7 mm rod diameter [9, 16]) which can be pro-duced by casting using the M.A. without Cu. As it was ex-pected, the 1 mm diameter rods are not fully amorphous, but they still contain a large quantity of the amorphous ma-trix, as resulted from the x-ray diffraction studies presented in Fig. 1. In fact, such relatively small volume fraction of (nano)crystalline inclusions cannot be surely ruled out us-ing only the regular x-ray diffraction in Bragg-Brentano configuration. As seen in Fig. 1, the crystalline volume fraction increases as the Cu content increases. The phases which could be identified are marked in the plot and mainly consist on bcc FeCo solid solution and the tetrago-nal (Fe,Co)2B. Additionally, some of the peaks may be at-tributed to pure Cu and to other fcc phase, the complex Fe23B6-type.

Figure 1 X-ray diffraction patterns using the synchrotron radia-tion of the as-cast 1 mm diameter rods produced from the {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 1, 2 and 3) al-loys.

The presence of Cu in the diffraction pattern is not un-usual, because the enthalpy of mixing of Cu with the other metals from this composition is positive [17], so Cu will not form intermetallic phases with other elements. In the case of FINEMET-type alloys [10], which are fully amor-phous, the first precipitate phase upon heat treatment of amorphous ribbons is always Cu, with the crystalline grains of about 10 nm diameter. In our case it seems that Cu precipitates directly upon casting. This could be ex-plained by assuming two factors: the positive heat of mix-ing and a slightly higher cooling rate than the critical one.

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The thermal behavior upon heating at a constant rate of 0.33K/s is presented in Fig. 2. The DSC traces of the Cu-containing alloys are presented together with the one cor-responding to the fully amorphous M.A. 1 mm diameter rod. The M.A. undergoes clearly a glass transition event, followed by the supercooled liquid region and crystalliza-tion. Further, two other small exothermic peaks can be ob-served. It was published previously that this type of alloy forms by primary crystallization only one complex fcc phase of the type Fe23B6 [12, 18]. The other peaks are as-sociated to the final transformation of the remaining matrix.

Figure 2 DSC curves of {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96 · Nb0.04}100-xCux (x = 0, 1, 2 and 3) 1 mm diameter rods.

Once the Cu is added to the M.A., the glass transition

event is not anymore clearly seen. Instead, a relatively large exothermic event, usually associated to the nanocrys-tallization of Cu [19] is present. For all three Cu-containing samples the crystallization takes place in two steps and finishes at a lower temperature than in the case of M.A. The last endothermic event, at around 1200K, can be attributed to the α−γ transition of the FeCo solid solution [20] (which is also a magnetic transition). The event is pre-sent also in the cooling curves (not presented here), but it does not appear on the case of M.A. This is due by the fact that upon crystallization, only one crystalline phase forms, the Fe23B6-type phase [12, 18], and its magnetic order-disorder transition takes place without structural modifica-tion. Such transitions do not necessary trigger a change in the DSC signal. The melting behavior is not much changed; judging from DSC melting traces the Cu-containing alloys and the M.A. as well, are close to eutec-tic.

Upon complete crystallization, the phases in the Cu-containing samples do not change. The diffraction patterns of the crystallized rods, as recorded at 1100 K, are pre-

sented in Fig. 3. The peak corresponding to the FeCo solid solution increases in intensity, indicating a larger volume fraction of this phase. An increasing of the volume fraction of the tetragonal (Fe,Co)2B and fcc Fe23B6-type phases could be observed as well. The peaks which were identi-fied at the beginning (prior heating) as fcc Cu are still pre-sent. However, the crystallization behavior is completely different in comparison with the one known for the M.A. [9], [12].

Figure 3 X-ray diffraction patterns using the synchro- tron radiation of the 1 mm diameter {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 1, 2 and 3) rods, recorded with the samples at 1100 K.

Figure 4 Saturation magnetization data as a function of tempera-ture for: (a) {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 1, 2 and 3) 1 mm diameter rods, transformed data, heating curves, (b) as-measured data, heating and cooling, only for x = 1.

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Figure 4 (a) shows the variation of the saturation mag-netization upon heating, for all 3 Cu-containing samples. Here should be pointed that the plot were transformed us-ing the algorithm proposed by Herzer, when the tempera-tures approaches the Curie temperature [15]. In order to minimize the errors, the experimental results were plotted as (MS)1/ß versus T, with the exponent β = 0.36 (the Heisenberg exponent) [15]. The Curie temperature was considered to be the temperature were the (MS)1/ß deviated from linearity. Due to the device limitation, the measure-ments were done up to 900 K, i.e. a temperature slightly above the first exothermic event observed upon DSC measurements (see Fig. 2). Figure 4(b) shows the raw measured data, heating and cooling completely, measured for the alloy with 1 at% Cu. Figure 5 shows the magnetiza-tion curves recorded at room temperature for all Cu-containing samples, as well as for the M.A. The inset shows the all hysteresis behavior. The Curie temperatures for 1 mm diameter rod samples, together with the coerciv-ity and saturation magnetization as measured by VSM, are summarized in Table 1.

Figure 5 Magnetic behaviour at room temperature of {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 0, 1, 2 and 3) 1 mm diameter rods. The inset shows the entire hysteresis loops.

Table 1 Coercivity Hc, Curie temperature Tc and saturation magnetization Ms for {[(Fe0.5Co0.5)0.75Si0.05B0.20]0.96Nb0.04}100-xCux (x = 0, 1, 2 and 3) as-cast 1 mm diameter rods.

Hc [A/m] Tc [K] Ms [Am2/kg] M.A. 4 720 112 (M.A.)99Cu1 5 700 107 (M.A.)98Cu2 65 713 109 (M.A.)97Cu3 140 708 107

As it can be observed from the Figs. 4 and 5, the sam-ples show a different saturation magnetization at room temperature, which decreases as the temperature increases. However, the saturation values are comparable (see Table I) within the measurement errors. This statement is valid also for Curie temperatures, which remain very close (their values differ with maximum 2,5 %). This could be under-stand if one assume that the magnetism of the samples come in fact mainly from the amorphous matrix. Interest-ing is that here, as can be observed from Fig. 5, the slope of the first magnetization curves is higher for the Cu-containing alloys than for the mother composition. The co-ercivities are very small and also comparable. It is known that the coercivity strongly depends on the structure of the sample [21, 22]; so it is possible to say that the M.A. and (M.A.)99Cu1 samples are quite similar from structural point of view. Another interesting feature is the presence of only one single magnetic transition in thermomagnetic curves (Fig. 4(a)), even though all samples show beside the amor-phous matrix several other magnetic phases. By analyzing the as-measured data (Fig. 4(b)) one can see that the satu-ration magnetization become almost zero and starts to in-crease only when the temperature exceeds 850 K. Just upon cooling one can see the two magnetic phase behavior. For simplicity, in Fig. 4 (b) only the data corresponding to 1 at% Cu containing sample are plotted, because the other two samples behave similarly.

The temperature at which the saturation magnetization starts to increase again corresponds to the first exothermic peak seen in DSC (see Fig. 2). So, this event cannot be as-sociated only to the growth of Cu, as was previously sup-posed. This behavior could be explained if one assumes that the crystal nuclei embedded in amorphous matrix are small and the distance between them is smaller than the exchange length. In this way their contribution to the total magnetic moment is negligible and the measured Curie temperatures are in fact the magnetic-paramagnetic transi-tion temperatures of the amorphous matrix. This explana-tion is consistent with the coercivity values, with the dif-fraction data and with the calorimetric measurements. The absence of the domain pinning could explain the small val-ues of the coercivity, as well as the increased slope of the first magnetization curve. However, in order to confirm this asumption, TEM investigations are needed.

4 Conclusions The effect of Cu additions (1, 2 and 3

at.%) on the glass forming ability and magnetic properties of [(Fe0.5Co0.5)0.75Si0.05B0.20]96Nb4 was studied. The GFA decreases as Cu content increases. However, 1 mm diame-ter composites rods with a large volume content of amor-phous matrix were possible to cast from all Cu-containing compositions. Cu additions change not only the crystalliza-tion behavior, but also the magnetic properties. Upon heat-ing, the crystalline nanoscale phases which form directly

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during the casting grow without further transformation. The magnetic behavior up to 850 K is characteristic for a material which shows only one single magnetic phase. The Cu additions seem to increase the initial permeability and, up to a level, to decrease the coercivity, at this moment the mechanism being not fully ruled out.

Acknowledgements The authors thank R. Li, S. Pauly, S. Scudino for stimulating discussions and B. Bartusch, S. Donath, M. Frey, H. Schulze for technical help.. This work was supported by the EU within the framework of the European Network of Ex-cellence on Complex Metallic Alloys (NoE CMA) and of the re-search and training networks on Ductile BMG Composites (MRTN-CT-2003-504692). The support of A. von Humboldt as-sociation and DAAD is acknowledged as well.

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