Nanoscale precipitates and phase transformations in a rapidly-solidified Fe–Pt–B amorphous alloy

Journal of Alloys and Compounds 402 (2005) 78–83

Nanoscale precipitates and phase transformations in a rapidly-solidified Fe–Pt–B amorphous alloy Dmitri V. Louzguine-Luzgin ∗ , Wei Zhang, Akihisa Inoue Laboratory for Non-equilibrium Materials, Institute for Materials Research, Tohoku University, Katahira 2-1-1, Aoba-Ku, Sendai 980-8577, Japan Received 8 March 2005; accepted 31 March 2005 Available online 15 June 2005

Abstract The phase transformations in a rapidly solidified (Fe0.75 Pt0.25 )75 B25 alloy on heating were studied by X-ray diffractometry, differential scanning and isothermal calorimetries. An as-solidified structure was examined by transmission electron microscopy (TEM). The results of TEM, including nano-beam diffraction and high-resolution imaging, clearly indicate the existence of nano-scale cubic cF4 Fe(Pt) solid solution particles within the as-solidified amorphous matrix. The particles of the cF4 Fe(Pt) phase start growing at elevated temperature and undergo ordering forming a tP4 FePt compound followed by the precipitation of a tI12 Fe2 B phase from the residual amorphous matrix. The ordering of the cubic phase takes place after an incubation period which allows to anticipate a nucleation and growth mechanism. This process is followed by further growth of the particles. The transformation behavior on heating is studied in detail. © 2005 Elsevier B.V. All rights reserved. Keywords: Amorphous materials; Nano-structures; Liquid quenching; Calorimetry; Transmission electron microscopy

1. Introduction Hard ferromagnetic alloys have been obtained in different systems. In particular, Fe–Pt alloys exhibit good permanent magnetic properties [1–3]. In particular, Fe61.5 Pt38.5 alloy being solution treated and quenched from the disordered fcc cF4 (Pearson symbol) ␥ solid solution phase region and subsequently aged at 773 K exhibits good hard magnetic properties of Br = 1.08 T, Hc = 340 kA/m, and (BH)max = 159 kJ m−3 [3], where Br is remanence, Hc the coercivity and (BH)max is maximum energy product. Its microstructure consists of an ordered ␥1 tP4 phase [3,4] in accordance with the phase diagram [5]. Good hard magnetic properties are obtained due to a transformation from the ␥ cF4 Fe solid solution phase to the ␥1 tP4 FePt phase having a high crystal magnetic anisotropy [4,6].

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The Fe–Pt films with smaller grain size and somewhat better magnetic properties were produced by sputtering [7–9]. It was also reported that by forming nano-scale multilayer and rapid thermal processing, a nano-composite structure consisting of the hard tP4 FePt and the soft cF4 Fe-rich phase has been formed. The intergrain exchange coupling mechanism has been used for developing high energy products [9]. It has been also found that the (Fe0.65 Pt0.35 )83 B17 alloy possesses higher coercivity in the annealed state compared to binary Fe–Pt alloys [10]. B is also known to be an appropriate alloying element for obtaining amorphous and nanocrystalline Fe-based alloys [11]. Br , Mr /Ms , Hic , and (BH)max of the rapidly solidified Fe80−x Ptx B20 (x = 20, 22, 24) ribbons in the annealed state are in the range of 0.93–1.05 T, 0.79–0.82, 375–487 kA/m, and 118–127 kJ/m3 , respectively (Hic is intrinsic coercivity, M is magnetization). The good hard magnetic properties are interpreted to result from the exchange magnetic coupling between the nano-scale hard ␥1 tP4 FePt and soft magnetic ␥ cF4 Fe(Pt) solid solution as well as Fe2 B phases [12]. Fe–Pt–P rapidly solidified alloys were also found to possess good magnetic properties [13].

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The melt-spun (Fe0.75 Pt0.25 )80–75 B25–30 ribbons obtained recently were found to be X-ray amorphous, that is, no sharp diffraction peaks except for the broad peaks were observed [14]. These alloys were also found to possess good hard magnetic properties including high intrinsic coercivity values up to 400 kA/m in the nano-crystallized state [14]. These alloys are promising to be applied as nano-composite permanent magnets. Thus, the detailed study of the phase transformations on heating related to the devitrification process is required. In the present work we study the structure and devitrification behavior of a rapidly solidified (Fe0.75 Pt0.25 )75 B25 alloy.

2. Experimental procedure An ingot of (Fe0.75 Pt0.25 )75 B25 alloy (alloy composition is given in nominal at.%) was prepared by arc-melting mixtures of Fe 99.9 mass% purity, Pt 99.9 mass% purity, and B 99.9 mass% purity in an argon atmosphere. From this ingot, ribbon samples of about 20 ␮m in thickness and 0.8 mm in width were prepared by rapid solidification of the melt on a single copper roller at a roller tangential velocity of 42 m/s. The structure of the samples was examined by X-ray diffractometry (XRD) with monochromatic Cu K␣ radiation. The phase transformations were studied by differential scanning calorimetry (DSC) at a heating rate of 0.67 K/s and differential isothermal calorimetry (DIC). During isothermal calorimetry the samples were heated up to the testing temperature at the highest possible heating rate of 1.67 K/s. The temperature and enthalpy for the DSC were calibrated within ±0.2 K and ±0.1 J/g using the melting of In as a standard. Transmission electron microscopy (TEM) investigation was carried out using a JEM 2010 (JEOL) microscope operating at 200 kV having a nano-beam diffraction (NBD) facility. The sample for TEM was prepared by the ion polishing technique.

3. Results Only broad diffraction peaks are observed in the XRD pattern of the rapidly solidified (Fe0.75 Pt0.25 )75 B25 alloy (all data in this work are related to this alloy) as shown in Fig. 1(a). As found by the Lorentzian fitting the first and the second broad maximums are located at 42.3 and 76.3◦ of 2θ corresponding to wave vectors of 29.4 and 50.4 nm−1 , respectively. The nano-particles of 2–6 nm in size are seen in the bright and dark-field TEM images Fig. 2(a) and (b) of the as-solidified sample. Smaller particles can be hardly identified. These particles do not produce the detectable diffraction peaks in Fig. 1(a) due to their small size and limited volume fraction. The selected-area electron diffraction (SAED) pattern (Fig. 2(a) insert) shows an amorphous halo and welldefined diffraction rings, yet significantly broadened, which can be well indexed with the fcc ␥ cF4 Fe(Pt) lattice with a lattice parameter of 0.38 nm. ␥ cF4 Fe(Pt) is a continuous solid

Fig. 1. X-ray diffraction patterns of the studied alloy: (a) in as-solidified state, (b) after DIC at 647 K for 4 ks, after DIC at 795 K (c) for 330 s and (d) for 760 s.

solution phase between Fe and Pt. The nano-beam electron diffraction (NBD) pattern (Fig. 2(b) insert) taken at a probe size of 5 nm also can be indexed according to a cubic cF4 lattice. The high-resolution TEM (HRTM) image (Fig. 3) also confirms co-existence of the amorphous and nano-crystalline cF4 Fe(Pt) phases. The DSC trace (Fig. 4) on heating exhibits a broad exothermic peak with an onset temperature of 665 K and a sharp exothermic peak consisting of two shoulders starting at 820 K. No exothermic peak but just a variation of the baseline is observed during isothermal calorimetry trace taken at 647 K (Fig. 5). At the same time, broad but clear diffraction peaks of the fcc cF4 Fe(Pt) phase are seen in Fig. 1(b). One should note a weak peak at about 54◦ of 2θ corresponding to the prohibited reflection (2 1 0) which may indicate partial ordering of the fcc cF4 Fe(Pt) solid solution phase. Using a multi-peak fitting by the Lorentzian function it was possible to separate and index the diffraction peaks in Fig. 1(b). The lattice parameter of the cF4 Fe(Pt) solid solution is found to be 0.379 nm. The isothermal calorimetry trace obtained at 795 K is shown in Fig. 6(a). One can see that the shoulders of the sharp exothermic peak observed using DSC at 820 K (Fig. 4) become quite well separated. The transformation at 795 K starts after an incubation period of about 60 s which indicates nucleation and growth process. Two XRD patterns taken from the samples subjected to the DIC measurements at 795 K up to the completion of the first (Fig. 1(c)) and the second (Fig. 1(d)) exothermic peak indicate the formation of the tP4 FePt and tI12 Fe2 B compounds, respectively. Although the

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Fig. 2. (a) Bright-field TEM image (the insert: SAED pattern) and (b) dark-field TEM image (the insert: NBD pattern; camera length is different from that in SAED) taken in as-solidified state.

Fig. 5. DIC trace obtained at 647 K.

Fig. 3. High-resolution TEM image in as-solidified state. The insert: NBD pattern from the amorphous phase.

Fig. 4. DSC trace of the studied alloy scanned at 0.67 K/s.

Fig. 6. (a) DIC trace obtained at 795 K, (b) Avrami plots corresponding to two peaks in (a).

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samples after DSC and DIC were clean and did not exhibit a visible oxide film on the surface, a weak peak at 28◦ of 2θ (Fig. 1(d)) is considered to be caused by partial oxidation as no such peak is observed for the samples annealed in vacuum. However, it is very difficult to keep exact temperature and the heating rate in the furnace as in the DIC device while a slight change in temperature influences significantly the transformation kinetics. Thus, DIC is used for annealing of the samples followed by XRD investigations. The lattice parameters of the tP4 FePt phase in Fig. 1(d) generated by using the least-squares fitting of the four highangle peaks (3 1 1, 1 1 3, 2 2 2, and 3 1 2) are a = 0.3846 and c = 0.3706 nm. These values are very close to that of the equiatomic FePt phase [15]. It was impossible to confirm precisely the lattice parameters of the tI12 Fe2 B compound as only the strongest low-angle peaks are observed (Fig. 1(c) and (d)). However, it is known that this phase is a line compound. Fig. 7 also shows a number of the XRD patterns in the annealed state. For example, Fig. 7(b) illustrates that the tI12 Fe2 B phase is absent in the first half of the peak 1 (Fig. 6(a)) that argues for an independent nucleation of the tP4 FePt and tI12 Fe2 B phases. The size of the tP4 FePt grains can be estimated using broadening of the XRD peaks. The procedure is discussed in details in Refs. [16,17], for example. It is widely used and the results match well with those of TEM observation [18]. The resulted average particle sizes calculated using (1 1 1) and (2 2 2) diffraction peaks in Fig. 1(c) are 14 and 15 nm, respectively. The β1 1 1 /β2 2 2 ratio of 0.733 where β is an

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intrinsic broadening of the corresponding diffraction peaks is close to cos θ 2 2 2 /cos θ 1 1 1 ratio of 0.760 which indicates that the broadening is mostly caused by a small grain size. The instrumental broadening of the peaks was obtained from the XRD pattern (Fig. 7(d)) of the reference sample annealed for 36 ks at 973 K and cooled down with furnace. The grain size of the tP4 FePt phase increases with annealing time and temperature. For example, it reaches 26 nm after annealing for 760 s at 795 K. The tP4 FePt and tI12 Fe2 B phases are the only observed phases in the sample annealed for 760 s at 795 K as well as for 36 ks at 973 K (Figs. 1 and 7(d)). Smaller tP4 FePt grains are obtained after annealing at 795 K for 180 and 220 s, i.e. in the early stages of cF4 Fe(Pt) → tP4 FePt transformation (Fig. 7(b) and (c)) related to peak 1 in Fig. 6(a). The average particle (grain) sizes calculated using (1 1 1) and (2 2 2) diffraction peaks in Fig. 7(c) are 8.5 and 9 nm, respectively. Seven nano-meters grain size was obtained from Fig. 7(b) (only (1 1 1) diffraction was used). Thus, during the first phase transformation the tP4 FePt phase nucleates and grows at the expense of the amorphous matrix. The overlapped DIC peaks at 795 K (Fig. 6(a)) were fitted with two Gaussian functions having the maxima at 183 and 425 s. According to the following kinetic law [19] for the volume fraction (x) transformed as a function of time (t): x(t) = 1 − exp[−Kt n ]

(1)

the Avrami plots of ln[−ln (1 − x)] versus ln (t) are obtained (Fig. 6(b)). The Avrami exponent n values [20,21] of 2.9 and 3.1 were obtained from the least-squares fittings.

4. Discussion

Fig. 7. X-ray diffraction patterns of the studied alloy: (a) after DIC at 647 K for 4 ks. After DIC at 795 K (b) for 180 s and (c) for 220 s. (d) The sample annealed in furnace for 36 ks at 973 K.

The results of TEM, including NBD and HRTEM imaging clearly indicate the formation of a limited volume fraction of the nano-scale cubic cF4 Fe(Pt) solid solution phase within the as-solidified amorphous matrix on rapid solidification. This phase exhibits small growth rate and the largest particles attain a size of about 6 nm. The cF4 Fe(Pt) solid solution phase starts growing intensively at 665 K and transforms to tP4 FePt ordered compound of nearly equiatomic composition at higher temperature. At the same time, it is shown that the residual amorphous phase is enriched in B, because the tI12 Fe2 B compound precipitates from the residual amorphous phase right after the cF4 Fe(Pt) → tP4 FePt transformation. Although the Avrami plot (Fig. 6(b)) for the cF4 Fe(Pt) → tP4 FePt transformation monitored by DIC as peak 1 in Fig. 6(a) is linear this reaction actually consists of two parts: cF4 Fe(Pt) → tP4 FePt transformation which is finished in the beginning of the peak 1 (see Fig. 7(b)) and further intensive growth of the tP4 FePt particles at the expense of the amorphous matrix. Both these processes are related to peak 1 of DIC in Fig. 6(a). Peak 2 in Fig. 6(b) is related to the residual amorphous → tI12 Fe2 B transformation. Co-existence of two

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The low growth rates and formation of a nano-structure were also observed in the Fe–Nb–B alloy on primary devitrification during annealing for 1.2 ks at 900 K [23]. Earlier fine structure was also obtained in the Fe–B–Si alloy [24]. The tP4 FePt phase has a wide area of homogeneity, while the tI12 Fe2 B phase is a line compound [25]. Thus, according to the alloy composition one can assume that FePt phase is somewhat enriched in Pt.

5. Conclusions Fig. 8. Average size of the tP4 FePt particles at 795 K as a function of time.

intermetallic compounds: tP4 FePt and tI12 Fe2 B in Fig. 1(c) indicates an overlapping of two exothermic reactions (peaks 1 and 2). The precipitation of the tI12 Fe2 B phase starts before the completion of peak 1 (which is cF4 Fe(Pt) → tP4 FePt transformation and further growth of tP4 FePt phase) in Fig. 6(a) as the residual amorphous phase becomes gradually enriched in B. The kinetics of the second phase transformation related to peak 2 (residual amorphous → tI12 Fe2 B) allows its precipitation conjointly with growth of tP4 FePt particles. The incubation period of 60 s obtained for the cF4 Fe(Pt) → tP4 FePt transformation at 795 K indicates nucleation and the first-order type transformation which is also supported by the existence of the two-phase areas around the tP4 FePt phase in the Fe–Pt phase diagram [5] while the obtained Avrami exponent values may indicate diffusioncontrolled growth of nuclei at an increasing nucleation rate. An interface-controlled 3-dimensional growth at the obtained Avrami exponent values close to 3 would lead to a zero nucleation rate. The average size of the tP4 FePt particles after annealing for 760 s at 795 K is about 26 nm indicating their small growth rate. Usually small particle size is obtained when particles exhibit diffusion-controlled growth which is consistent with the above data of the kinetics analysis. Fig. 8 shows a plot of the tP4 FePt particles size as a function of time indicating non-linear growth. Pt may be the element limiting growth rate of the cF4 Fe(Pt) and tP4 FePt phases. At the same time, the diffusion coefficient of Pt in fcc Fe, for the lowest temperature of 1233 K at which it has been measured, is only three times less than the self-diffusion coefficient of fcc Fe [22]. However, one should take into account quite low temperatures (<850 K) at which growth of the cF4 Fe(Pt) and tP4 FePt phases takes place. By extrapolation of the high-temperature values to lower temperatures [22] the diffusion coefficient of Pt in cF4 Fe is predicted to be in the√order of 10−23 m2 /s at 850 K. Thus, the diffusion length 2Dτ in 60 s would be only 10−10 m or 0.1 nm. Of course, this is valid only for the bulk crystalline materials while in amorphous alloys diffusion kinetics is different. However, it makes a qualitative explanation of the observed low growth rates.

The structure of the rapidly solidified (Fe0.75 Pt0.25 )75 B25 alloy contains a limited volume fraction of the nano-scale cubic cF4 Fe(Pt) solid solution particles of about 4 nm in size embedded in the amorphous matrix. The nano-particles of the cF4 Fe(Pt) phase start growing at the elevated temperature and then undergo ordering forming the tP4 FePt compound of about 15 nm in size which is followed by the formation of the tI12 Fe2 B phase from the residual amorphous matrix. The ordering of the cF4 phase takes place after an incubation period, the existence of which allows to anticipate a nucleation and growth mechanism. This process takes place conjointly with further growth of the tP4 FePt particles at the expense of the residual amorphous matrix. The formation of the tI12 Fe2 B phase from the residual amorphous matrix begins when the residual amorphous matrix is enriched in B up to a certain value. The observed low growth rate of the tP4 FePt phase leading to the formation of a nano-structure is obtained due to the diffusion-controlled growth mechanism and very low calculated diffusion coefficients of Fe and Pt in cF4 Fe at the temperature of transformation (795–860 K). The obtained results can be used for designing new amorphous, nano-structured and -composite magnetic alloys in the Fe–Pt–B system.

Acknowledgement This work was supported by the Grant-in-Aid of Ministry of Education, Sports, Culture, Science and Technology, Japan, Priority Area on “Materials Science on Bulk Metallic Glasses” and the Grant-in-Aid (Wakate B) of Ministry of Education, Sports, Culture, Science and Technology, Japan N: 16760559.

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