DEVELOPMENT OF NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS

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1 Trans. Indian Inst. Met. Vol.58, No. 6, December 2005, pp TP 2016 DEVELOPMENT OF NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS Amitava Mitra and Ashis Kr. Panda Magnetism Group, National Metallurgical Laboratory, Jamshedpur (Received 29 October 2004 ; in revised form 20 October 2005) ABSTRACT The structural and soft magnetic properties of Fe and FeCo-based nanocrystalline alloys have been investigated. These alloys were initially prepared in the form of amorphous ribbons with consistent properties by optimising different melt-spinning parameters. The structural behavior and soft magnetic properties depended on the alloy chemistry. The effect of metalloids in the Fe 73.5 Nb 3 Si 22.5-X B X (X = 5, 9, 10, and 19 at %) was studied. X-ray diffractograms showed that formation of a-fe(si) and /orfe 3 Si nanoparticles were responsible for the superior soft magnetic properties of the alloy with 9 at% Boron. All other alloys (X=5,10,11.25 and 19at%) exhibited early appearance of highly magnetocrystalline anisotropic boride phases leading to deterioration in soft magnetic properties. The role of extra alloying elements Al and Mn in the FeNbCuSiB system was also investigated. The alloy exhibited superior soft magnetic properties with a coercivity value of 0.32A/m (~ 4mOe) when heat-treated at 790K for 15min. Transmission electron microscopy study showed that this was due to the formation of ~ 6.0nm sized a-fe(si,al) and /or Fe 3 (Si,Al) nanoparticles. However, this alloy has a limitation on its use for high temperature soft magnetic application due to its low Curie temperature in nanocrystalline state. Hence, a new Fe 40 alloy system was developed in which the Curie temperature of 736K in amorphous state increased above 1000K on annealing to nanocrystalline state. The saturation magnetization of the annealed alloy was also found to increase due to the formation of nanocrystalline -(Fe,Co)(Si,Al) phase with higher magnetization suggesting the suitability of the alloy for high temperature soft magnetic applications. 1. INTRODUCTION Over the past several decades, magnetic amorphous alloys have been investigated for applications in the areas of sensors, transformer cores, inductive devices, etc. Most recently, research interest on nanocrystalline alloys has dramatically increased due to the enhanced magnetic properties compared to their amorphous state. The benefits found in the nanocrystaline alloys stem from their chemical and structural variations on a nanoscale which are important for developing optimal magnetic properties. The schematic representation of magnetic properties of various soft magnetic alloys is shown in Fig. 1. The nanocrystalline Fe-based soft magnetic materials were first observed in late 80 s in melt-spun Fe-Si-B alloys containing small amounts of Nb and which were named as FINEMET. The Fig. 1 : Relationship between Saturation Induction and permeability of various soft magnetic materials. Fe-Si-B-Nb-Cu amorphous phase transforms to a bodycentered cubic (bcc) Fe-Si solid solution with grain sizes of about 10 nm during annealing around the

2 TRANS. INDIAN INST. MET., VOL. 58, NO. 6, DECEMBER 2005 crystallization temperature. The presence of small amounts of Cu helps to increase the nucleation rate of the bcc phase while Nb retards the grain growth. These Finemet alloys provide low core losses (even lower than amorphous soft magnetic alloys such as Co-Fe-Si-B), exhibit saturation induction of about 1.2 T, and also exhibit very good properties at high frequencies, comparable to the best Co-based amorphous alloys. While many of the soft magnetic properties of Finemet-type nanocrystalline alloys are superior, yet they exhibit lower saturation inductions than Fe-metalloid amorphous alloys, mainly because of the lower Fe content to attain amorphization. In order to remedy this problem, another class of Febased nanocrystalline alloys was developed by Inoue and co-workers at Tohoku University 2. These Nanoperm alloys based on the Fe-Zr-B system contain larger concentrations of Fe (83-89 at.%) compared to the Finemet alloys (~ 74 at.% Fe) and have higher values of saturation induction (~ T). 2. EXPERIMENTAL About 150g of cylindrical ingot of FeNbCuSiB based alloy was taken in a quartz crucible for melt-spinning. The melt was ejected using Ar-gas pressure of 0.5bar. The nozzle diameter and the separation between the nozzle and the spinning wheel were optimized as 1.5mm and 0.5mm respectively for obtaining continuous ribbon 3. Quench rate was optimized by evaluating the crystallization behavior and magnetic properties of the ribbons prepared at different quenching wheel velocity. Fig. 2 shows the variation of crystallization temperature and the coercivity with quenching wheel velocity. Thickness (t) of the ribbon also depends on the wheel velocity (V s ) and follows the relationship 4 t ~ Vs m (1) where 0.5 m 1. The ideal cooling condition is at m = 0.5 whereas when m=1, Newtonian cooling occurs. The variation of ribbon thickness with wheel velocity is shown in Fig. 3. The experimental results suggest that the ribbon prepared with the wheel velocities ranging between 27m/sec and 42m/sec. exhibits consistent properties 5. However, in most cases the wheel velocity was kept at 35m/sec. Magnetic properties of the materials were determined in an open flux configuration using a system developed at National Metallurgical Laboratory 6. Different magnetic parameters are schematically shown in Fig. 4 where the magnetic induction (B) has been plotted against the magnetising field (H). Excellent soft magnetic materials should have high initial permeability ( i =B/H) H 0 =1+c i, c i is the susceptibility) and low coercivity, H C. Curie temperature was measured from thermal variation of ac susceptibility plots. The crystallisation behavior was studied using Differential Scanning Calorimeter, DSC (Perkin Elmer DSC-7) and Thermal Electrical Resistivity (Sinku-Riko TER-2000). Microstructure and phases were identified using Transmission Electron Microscope (Philips-CM-200) and X-ray diffractometer (Siefert PTS 3003). Fig. 2 : Variation of Crystallisation onset temperature and Coercivity with wheel velocity Fig. 3 : Variation of ribbon thickness with wheel velocity 982

3 MITRA AND PANDA : NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS Fig. 4 : Magnetic hysteresis loop showing different parameters 3. EFFECT OF METALLOIDS IN Fe-Nb- Cu-Si-B The magnetic softness after nano-crystallization of FeNbCuSiB based alloys is caused by the random orientation of -Fe(Si) or ordered Fe 3 Si nanoparticles dispersed in an amorphous matrix which averages out the magnetocrystalline anisotropy energy 7. The magnetic softness is further enhanced by the vanishingly small magnetostriction constant (l) of the system originating from the balance between the magnetostriction constant of the crystallites (lcr < 0) and the amorphous matrix (lam > 0) 8. The values of the lcr, lam and the specific magnetization of -Fe(Si) depend on the silicon content within the material 9. Hence, Si plays a crucial role on the mechanism of nanocrystallization 10. The ordered Fe 3 Si phase with DO 3 superstructure and also the superior soft magnetic properties are found when the Si content of the alloy is above a critical value 11. The influence of metalloid on the transport and the magnetic properties has been investigated for the Fe 73.5 Nb 3 Si 22.5-X B X amorphous alloys with X = 5, 9, 10, and 19 at % Magnetic properties Figure 5 shows the temperature variation of susceptibility of the as-cast ribbons after normalizing by their room temperature value. The sharp drop in the susceptibility indicated the transformation of the Fig. 5 : Thermal variation of normalized ac. Susceptibility plots of the alloys with different metalloid content material from ferromagnetic to paramagnetic state. The Curie temperature was estimated from the derivative of the normalized susceptibility plot and indicated by an arrow in Fig. 5. Except in Alloy#1 where T C = 631K, the Curie temperature was found to increase slightly from 625K at X = 9 to 630K at X = 19. Such low variation indicated that Curie temperatures of the measured alloys were weakly dependent on the Si and B-content. Magnetic coercivity and susceptibility of the sample annealed at different temperatures are shown in Fig. 6a and b respectively. At the initial stage of annealing, a decrease in coercivity and an increase of susceptibility were observed in all the measured alloys which were due to the relaxation of internal stress generated within the sample during preparation of the alloy by rapid solidification technique 12,13. After softening of magnetic properties at the initial stage of annealing, there was a tendency of magnetic hardening for all the measured alloys except for X = 19 alloy. In X=19 alloy continuous change in improvement in 983

4 TRANS. INDIAN INST. MET., VOL. 58, NO. 6, DECEMBER 2005 Fig. 6a: Annealing behavior of coercivity for the alloys with different metalloid content Fig. 6b: Annealing behavior of susceptibility for the alloys with different metalloid content properties observed up to the annealing temperature of 725K above which the magnetic properties deteriorated. In other measured alloys, further improvements in magnetic properties were observed after intermediate hardening. The annealing temperature at which superior magnetic properties were observed was dependent on the composition of the alloy. Annealing above stress relaxation temperature, the magnetic hardening (Fig. 6 a and b) in alloys X= 5,9,10 and was due to superparamagnetic behavior of small volume fraction of very fine nanoparticles which were weakly coupled and dispersed in large volume fraction of amorphous matrix 14. After magnetic hardening, an improvement in soft magnetic properties was observed which was primarily due to the formation of Fe 3 Si and/or a-fe(si) phase 7. The extent of magnetic softening was maximum in X = 9 alloy with Hc = 1.56A/m and c i ~ 1.35x 10 5 when the alloy was annealed around 800K. Above this temperature, rapid deterioration in soft magnetic properties were observed. Critical size and the volume fraction of Fe 3 Si and/or -Fe(Si) phase were important for the superior soft-magnetic properties 15. Fig. 7 : Temperature variation of electrical resistivity of the alloys with different metalloid content. Scan rate 10K min

5 MITRA AND PANDA : NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS 3.2. Transport properties Temperature variation of electrical resistivity of the measured alloys is shown in Fig. 7. The results are plotted after normalizing with the room temperature resistivity values of the as-cast samples. The Curie temperature (T C ) of as-cast alloy (Fig. 5) is also indicated in the resistivity plot. Above T C, the first change in slope of resistivity observed at a temperature designated by T xr as shown in Fig. 7, is due to the formation of crystalline phase. Resistivity was found to increase above crystallization temperature for 5 X 10 alloys whereas a decrease in resistivity was found in X = and 19 alloys. The resistivity of the amorphous alloy can be divided into three temperature regions (i) below Debye ( D ) temperature (ii) between Debye temperature ( D ) and Curie temperature (T C ) (iii) between Curie temperature (T C ) and crystallization temperature (T xr ). Applying diffraction 16 and spin order model 17, the temperature dependence of total resistivity of amorphous alloys may be described as; (T) = a 01 + a 21 T 2 T< D a 02 + a 12 T + a 22 T 2 D <T< T C (2) a 03 + a 13 T T C <T< T xr Debye temperature D was determined by best fitting of the experimental data of resistivity. The mathematical fitting of the equations was done for different temperature ranges with the correlation factor Fig. 8 : XRD plots of the measured alloys annealed at 825K more than 0.9. The temperature coefficient of resistivity at different regions for the measured alloys is shown in Table-I. Debye temperature ( D ) was found to increase with the decrease in Si content of Table 1 TEMPERATURE COEFFICIENTS OF RESISTIVITY DERIVED FROM THERMAL ELECTRICAL RESISTIVITY MEASUREMENT X D (K) T C (K) T xr (K) Coefficients of Resistivity a 01 a 02 a 03 a 12 a 13 a 21 a 22 ( 10-4 ) ( 10-4 ) ( 10-4 ) ( 10-4 )

6 TRANS. INDIAN INST. MET., VOL. 58, NO. 6, DECEMBER 2005 Fig.10 : Variation of particle size with the annealing temperature the alloy up to X = at%. Very low value of D = 418K was found for X = 19 alloy. 3.3 Evaluation of different phases To find the formation of different phases, X-ray diffractograms (XRD) were taken after annealing the samples at 825K and is shown in Fig. 8. It appeared from X-ray diffractograms that except for X = 9 alloy where nanocrystalline a FeSi and/ or Fe 3 Si phase formed, there existed two or more crystalline phases when the materials were annealed at 825K. The phases formed together with the nanocrystalline Fe 3 Si and/ or a FeSi after annealing at 825K were Fe 23 B 6 in X=5 alloy, Fe 3 B in X=10 alloy, Fe 2 B and Fe 3 B in X=11.25 alloy. In case of X=19 alloy Fe nanocrystalline phase was formed along with Fe 3 B and Fe 2 B. 4. INFLUENCE OF Al AND Mn IN FeNbCuSiB ALLOY 4.1 Structural evaluation Fig. 9 : TEM micrograph of the alloy annealed at 775K (a), 825K (b) and 875K (c) Microstructure of the materials was studied by Transmission Electron Microscopy for the as-cast and different heat treated samples. Fig. 9a shows microstructure of the sample annealed at 775K indicating the initiation of nanocrystallization. The SAD pattern in the inset of Fig. 9a shows the lattice 986

7 MITRA AND PANDA : NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS images of -Fe(Si,Al) and/or Fe 3 (Si,Al) planes from which a lattice parameter of nm was calculated. This value of lattice parameter is higher than the reported value in -Fe 3 Si (0.2835nm nm) 18. Thus presence of Al increases the lattice parameter of nanoparticles. The average grain size was measured on an assembly of 100 grains and was found to be 5.1± 0.4 nm for 775K /15min annealed sample. With the increase of the annealing temperature, the grain size increased as shown in Fig. 9 (b) and (c). Fig. 10 shows the variation of grain size with the annealing temperature. The maximum observed grain size was 11.0 ± 0.5 nm at 925 K/15 min annealed sample, which was much lower, compared to the reported value for FeNbCuSiB sample 19. At this stage the cause of such slow grain growth process is not known. 4.2 Magnetic behavior Fig.11 : Thermal variation of normalised a.c susceptibility plots of alloy in the as-cast state and at different annealing temperatures, T A Thermal variation of a.c. susceptibility was measured for the as-cast and heat treated samples and is shown in Fig. 11. The two peaks observed at 825K, 850K and 875K annealed sample were due to the presence of two phased structures i.e., amorphous and nanocrystalline in annealed samples. The variation of Curie temperature with annealing is replotted in Fig. 12. The Curie temperature was found to increase with annealing at the initial stage. Change of Curie temperature was also observed due to the stress relaxation in amorphous alloy. However, in the present case the observed variation was about 2% for 790K Fig.12 : Variation of Curie temperature of the alloy with annealing temperature Fig.13 : Variation of coercivity (a) and susceptibility (b) of the alloy with annealing temperature 987

8 TRANS. INDIAN INST. MET., VOL. 58, NO. 6, DECEMBER 2005 annealed sample, which was more than the expected change due to stress relaxation. The change of chemical composition due to formation of small percentage of very fine nanoparticles was the cause of such large change. TEM micrograph for 775K annealed sample (Fig. 9b) showed the formation of such fine nanoparticles although the Curie temperature measurement did not reveal such two-phased structure even after annealing at 790K. Thus the observed increase in Curie temperature was due to the stress relaxation as well as due to change of chemical composition of the system. As the annealing temperature was increased, the size and volume fraction of the nanoparticles increased resulting in two distinct ferromagnetic regions of different Curie temperature. The Curie temperature decreased in the amorphous phase due to depletion of Fe in it where as it increased in the nanocrystalline phase. Figure 13 (a) and (b) represent the annealing behavior of coercivity and susceptibility of the measured alloy. After annealing at 675K the coercivity decreased to 1.49± 0.08 A/m (~19±1 moe) from its as-cast value of 1.75 ± 0.08 A/m (~22 ± 1 moe). Similarly the susceptibility value also increased. The magnetic softness after initial stage of annealing was due to the relaxation of quenched-in internal stress introduced during preparation. However, on further annealing at 700K, magnetic hardness i.e., increase in coercivity and decrease in susceptibility, was observed. The cause of such magnetic hardness is not clear at this stage although some authors correlated it with the formation superparamagnetic 19 particles. Magnetic hardness may also be caused by precipitation of Cu, which may form clustering at a temperature much lower than the crystallization temperature. Such precrystallization clustering of Cu atom in FeNbCuSiB system took place after annealing at 675K/60min and was observed by Hono et al. 20 using 3-dimensional atom probe (3DAP) microscopy. The Cu-clustering induced internal stresses and also changed the exchange correlation length between nearest Fe-atom in the amorphous state, which made the material magnetically harder. After further annealing, coercivity decreased sharply up to an annealing temperature of 750K. The susceptibility was also found to increase from 3.0 x10 4 to 6.0 x10 4. TEM study of 725K annealed sample still exhibited amorphous nature of the alloy. However at this stage the -Fe started nucleating heterogeneously at the Cu clustering regions and thereby reducing the effective nonmagnetic inclusions resulting in decrease in coercivity. When the annealing temperature was further increased Si and Al would dissolve in -Fe forming -Fe(Si,Al) nanoparticles as evident from the TEM micrograph at 775K. The size of these nanoparticles was controlled by high atomic radius of Nb, which accumulated at the nanoparticle-amorphous interface and restricted the growth of the nano-particles by limiting diffusions 21. The coercivity further decreased due to the formation of nanoparticles and showed a minimum value of 0.32A/m (4mOe) at 790K. Large increase in susceptibility in this region was also observed with the peak value of 2.0 x10 5. TEM micrograph showed that the volume fraction of the nanoparticles was less than 50% in the present case as compared to 70% 22 in Fe 73.5 Nb 3 Si 13.5 B 9 alloy at an optimum heat-treatment condition, which resulted in minimum in coercivity. Some Al atom that partitioned in the Cu-rich sites as observed by Warren et al. 23 by 3-dimensional atom probe mapping might have interacted with Mn producing AlMn phase and hence reducing the number of nucleation centers for -Fe(Si,Al) phase. This might explain the low volume fraction of the nanoparticles in the optimum heat-treated sample for superior soft magnetic properties for the present alloy. However, the existence of such phases were not observed in the X-ray diffractograms due to the overlapping of reflections with the existing phases like Fe 23 B 6, -Fe(Si,Al). The minimum value of coercivity for the present alloy was 0.32A/m (~4.0mOe) at 790 K 24 which was also much lower than the reported value of 1A/m for Fe 73.5 Nb 3 Si 13.5 B 9 system. The particle size at this annealing temperature was 6.0 ± 0.5nm which was also lower compared to the Fe-Nb- Cu-Si-B alloy system (particle size ~ nm) which does not contain Al and Mn 7. It has been reported 21 that particle size was independent of Nb when it was more than 3at%. As the present alloy is containing Nb = 3.7at%, the lower particle size (~ 6nm) was due to the addition of Al. Such reduction of particle size in Al-containing alloy was also found by others 25. It is to be mentioned that any effect on the particle size on annealing time greater than 15min was not observed for the present alloy system which was supported by work of Zhou et al. 26. Thus, lower particle size in the present alloy 988

9 MITRA AND PANDA : NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS was due to partial dissolution of Al into Mn, which also prevented the growth of the nanophase. The low coercivity at the optimum heat-treatment condition could be explained by random anisotropy model DEVELOPMENT OF NANO- CRYSTALLINE MATERIALS FOR HIGH TEMPERATURE APPLICA- TIONS Fig.14 : DSC curve of Fe 40 alloy showing different stages of crystallization. The nanoparticles in the amorphous matrix are coupled through the amorphous phase. With the increase of the volume fraction of the nano-particles in FeNbCuSiB type alloys, the Curie temperature of the amorphous phase decreased (Fig. 12). Hence the materials are not suitable for high temperature applications. To overcome this, a new type of material was developed where Fe is partially replaced by Co in nanocrystalline soft magnetic alloys which enhances the Curie temperature of the alloy. The present work involves the characterisation of material with the nominal composition Fe 40 prepared by melt-spinning technique. 5.1 Crystallization and structure Fig.15 : XRD of as-spun and annealed Fe 40 alloy showing the formation of -(Fe,Co)(Si,Al) phase after annealing at 873K/15min. Figure 14 shows the DSC curve of melt-spun and annealed ribbons measured at a scanning rate of 40K/ min. The on-set of primary crystallization occurs at 820K. The XRD pattern (Fig. 15) shows that the asspun ribbon is fully amorphous and a-(fe,co)(sial) phase is formed after annealing at 873K for 15 mins. Fig. 16 shows the TEM micrograph of the annealed specimen with the selected area electron diffraction (SEAD) pattern at the inset. TEM study revealed that nanocrystalline particles were formed after annealing and are dispersed in the amorphous matrix. The nanoparticle size with annealing is shown in Fig Magnetic properties Fig.16 : TEM micrograph of Fe 40 alloy annealed for 15min. at 873K showing the formation of nanoparticles. Inset shows the selected-area electron diffraction (SAED) pattern The Curie temperature (T c ) of the materials was measured from the temperature variation of magnetization at a magnetic field of 2.8kA/m and is shown in Fig. 18. Magnetization in the as-spun state decreased rapidly with temperature and the Curie temperature of the amorphous phase was 736K. The material was annealed at 873K for 15 min and the Curie temperature was found above 1000K. The 989

10 TRANS. INDIAN INST. MET., VOL. 58, NO. 6, DECEMBER 2005 Fig.17 : Variation of particle size with annealing temperature. Fig.18 : Temperature variation of magnetization for estimation of Curie temperature measurement of as-spun and annealed Fe 40 alloy measured at a field of 2.8kA/m. Table 2 COMPARISON OF ROOM TEMPERATURE AC MAGNETIC PROPERTIES Property HIPERCO-50 Fe 44 Co 44 Zr 7 B 4 Fe 40 (HITPERM) annealed annealed at at 873K K Coercivity 320A/m a) 160A/m at 4kHz 95±8 A/m up to 20kHz (H ac c ) at 10kHz b) 175A/m at 10kHz F( // peak) 1kHz measured a) 2kHz measured at 80A/m 20 khz measured at 5A/m at 80A/m b) 20kHz measured at 200A/m Permeability 1800 measured at a) 700 measured at 200A/m and 200Hz 5A/m and 1kHz b) 6400 measured at 110A/m and 1kHz saturation magnetization of the annealed alloy was also found to increase 27. The formation of nanocrystalline -(Fe,Co)(Si,Al) phase with higher magnetization value was the cause of such increase in saturation magnetization. The other soft magnetic properties of the alloy were measured and a comparison of those properties with the crystalline (HIPERCO-50) and other nanocrystalline material (HITPERM) are shown in the Table II. 6. CONCLUSIONS The nanocrystalline soft magnetic materials were initially prepared in the form of amorphous ribbons by using a melt spinner. The melt spinning process 990 parameters like Crucible design, crucible-wheel separation and a critical range of quenching wheel velocity were optimized to get amorphous ribbons with consistent properties. Besides these processing conditions, the soft magnetic and electrical properties of nanocrystalline Fe 73.5 Nb 3 Si 22.5-X B X (X = 5, 9, 10, and 19) alloys also depended on the metalloid content. Temperature coefficient of resistivity after nanocrystallisation was found to depend on the Si content within a-fe(si) nanoparticles. The formation of a-fe(si) and / or Fe 3 Si nanoparticles were responsible for superior soft magnetic properties of X = 9 alloy annealed around 800K. All the alloys except X = 9 showed deterioration in magnetic properties at lower annealing temperature due to the

11 MITRA AND PANDA : NANOCRYSTALLINE MATERIALS FOR SOFT-MAGNETIC APPLICATIONS formation of strongly anisotropic boride phases like Fe 2 B, Fe 3 B and Fe 23 B 6. The soft magnetic properties of FeNbCuSiB alloy can be further improved by addition of extra alloying elements Al and Mn in the system. With optimum additions of these elements, the alloy exhibited ultrasoft magnetic properties with minimum coercivity of 0.32A/m (4mOe) and susceptibility of 2.0 x10 5. This improvement was due to the lowering of the magnetocrystalline anisotropy by a-fe(si,al) nanophase. Inspite of these excellent soft magnetic properties, the FeNbCuAlMnSiB alloy was unsuitable for high temperature soft magnetic applications due to reduction of Curie temperature of the amorphous phase upon nanocrystallisation. This limitation was overcome by partially replacing Fe by Co in a nanocrystalline soft magnetic alloys which enhances the Curie temperature of the material. The study indicated that the Fe 40 alloy subjected to annealing at 873K for 15min. developed a nanocrystalline structure with a-(fe,co)(si,al) nanoparticles dispersed in the amorphous matrix. The alloy exhibited high Curie temperature, high saturation magnetization and good ac soft magnetic properties. These properties made the alloy a potential candidate for high temperature soft magnetic applications. ACKNOWLEDGEMENT The authors wish to express their gratitude to Dr. I Chattoraj and Dr.S.R.Singh of National Metallurgical Laboratory, Jamshedpur for their interest in this work and valuable suggestions in explaining the experimental results. Thanks are also due to Director, NML in giving permission to publish the work. REFERENCES 1. Yoshizawa Y, Oguma S and Yamauchi K, J.Appl.Phys. 64 (1988) Makino A, Hatnai T, Inoue A and Masumoto T, Mat.Sci.Eng.A 226 (1997) Mitra A, Panda A K, Rao V, Singh S R and Ramachandrarao P, J. App. Surf. Sci. 182 (2001) Shingu P H and Ozaki R, Met. Trans. A 6 (1995) Panda A K, Roy S, Singh S R, Rao V, Pramanik S, Chattoraj I, Mitra A and Ramachandrarao P, Mater.Sci.Engg.A (2001) Panda A K, Basu S and Mitra A, J.Magn. Magn. Mat. 261 (2003) Herzer G, IEEE Trans.Magn. 25 (1989) Herzer G, Mater.Sci.Engg.A 133 (1990) 1 9. Zhu J, Pradell T, Clavaguera N and Claveguera-Mora M T, Mat. Res. Soc. Symp. Pro 285 (1997) Conde C F, Conde A, Nanostruct. Mat. 6 (1995) Muller M, Mattern N and Z Illgen, Metallkde 82 (1991) Herzer G, Phys.Scripta 49 (1993) Panda A K, Chattoraj I and Mitra A, J.Magn. Magn. Mat. 222 (2000) Skorvanek I, Handley R C O, J.Magn. Magn. Mat (1995) Panda A K, Ravikumar B, Basu S and Mitra A, J.Magn.Magn.Mater. 260 (2003) Cote P J, Miesel LV, Phys.Rev.Lett. 40 (1978) Bergmann G, Marquardt. P, Phys.Rev.B 17 (1978) Kulik T and Hernando A, Mat. Sci. Forum (1995) Kulik T, and Hernando A, J.Mag.Mag.Mater138 (1994) Hono K, Ping D H, Ohnuma M and Onodera H, Acta.Mater. 47 (1999) Herzer G Nanocrystalline soft magnetic alloys, in hand book of magnetic materials, vol.10, ed. by Bushow (Elsevier Sc.B.V, 1997) 22. Marin P, Vazquez M and Hernando A, J.Magn.Magn.Mater (1999) Warren P J, Todd I, Davies H A Cerezo A, Gibbs M R J, Kendall D and Major R V, Scripta.Mat. 41 (1999) Mitra A, Panda A K, Singh S R and Pramachandrarao, Phil.Mag. 83 (2003) Lim S H, Pi W K, Noh T H, Kim H J and Kang I K, J.Appl.Phys. 73 (1993) Zhou F, He K Y and Lu K, Nanosturct.Mater. 9 (1997) Mitra A, Kim H Y, Louzguine D V, Nishiyama N, Shen B and Inoue A, J.Magn.Magn.Mater. 278 (2004)