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Phase transformations and magnetic properties of melt-spun Pr7Fe88B5 ribbons during annealing
Journal of Alloys and Compounds 299 (2000) 258–263
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Phase transformations and magnetic properties of melt-spun Pr 7 Fe 88 B 5 ribbons during annealing Zuocheng Wang*, Shouzeng Zhou, Yi Qiao, Maocai Zhang, Run Wang State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China Received 29 August 1999; accepted 6 November 1999
Abstract ¨ X-ray diffraction and Mossbauer measurements indicate that melt spinning at different wheel velocities causes as-quenched Pr 7 Fe 88 B 5 ribbons to have distinctive structures. Depending on their as-quenched structures, the phase transformations of the ribbons during annealing may take place in one of the following sequences: (1) amorphous phase (Am)1Pr 2 Fe 14 B1a-Fe→Pr 2 Fe 14 B1a-Fe; (2) Am1a-Fe→Am91a-Fe→a-Fe11:7 phase1Pr 2 Fe 14 B→Pr 2 Fe 14 B1a-Fe; and (3) Am→Am91a-Fe→1:7 phase1a-Fe→ Pr 2 Fe 14 B1aFe. However, for all the ribbon samples, the microstructures after optimal annealing with respect to magnetic properties were found to only consist of two magnetic phases: Pr 2 Fe 14 B and a-Fe. The optimum magnetic properties of Hcj and Jr , and the squareness of the demagnetization curves of the annealed ribbons deteriorate drastically with increasing quenching rate of unannealed precursors. Taking into account TEM results, the values of aex and Neff derived from the temperature dependence of the coercivity, these deteriorated effects can be attributed to the formation of a coarser and more irregular microstructure during annealing in the samples initially melt spun with higher wheel speeds. 2000 Elsevier Science S.A. All rights reserved. Keywords: As-quenched structure; Phase transformation; Metastable phase; Nanocomposite magnets; Magnetic properties
1. Introduction One of the latest exciting developments in hard magnetic materials is nanocomposite permanent magnetic alloys, because they open a way for a new generation of isotropic bonded magnets with unusually high remanence, relatively high coercivity, high energy product, and low cost [1–3]. The excellent magnetic properties of these materials are known to arise from the exchange interaction between magnetically hard and soft phases. In order to achieve the most effective exchange interaction between the two phases, however, the optimal structure of these materials requires a homogeneous and narrow distribution of very fine grains, especially soft grains with diameter comparable to the domain wall width of the hard phase [1,4]. In addition, the predominant microstructure should be substantially free from intergranular phase, which would inhibit the exchange coupling between adjacent grains [5]. It is anticipated, therefore, that the magnetic
*Corresponding author. E-mail address: [email protected] (Z. Wang)
behavior of nanocomposite magnetic materials should be dependent on the intrinsic magnetic properties of hard and soft phases as well as their microstructures. So far, two types of Nd–Fe–B nanocomposite magnets, namely Fe 3 B / Nd 2 Fe 14 B and a-Fe / Nd 2 Fe 14 B [6–9], have been developed. It was found that, in order to achieve high performance, thermal processing of precursor materials, such as crystallizing overquenched ribbons, is usually needed because of the narrow range of quenching parameters over which optimum direct quench magnetic properties can be obtained. The crystallization process of the ribbons upon thermal treatment has been analyzed by different techniques [7,9]. However, in our literature, there are no studies on the influence of the initial structural properties of the overquenched precursor phase on the formation, structure and magnetic properties of the resulting nanocrystalline materials. It is established that the degree of as-cast amorphous state in the ribbons can be varied by changing the quenching of the melt during the flow-cast procedure. In this paper, we present a systematic study on microstructure evolution and magnetic properties of as-quenched amorphous and partially crystallized Pr 7 Fe 88 B 5 ribbons during annealing, and clarify how the
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as-quenched structures influence the crystallization process, the resulting nanocrystalline two-phase structures and their magnetic properties.
2. Experimental procedures Five ribbon samples of composition Pr 7 Fe 88 B 5 were obtained from the same master alloy by melt spinning on a copper wheel using different peripheral velocities (T 1 5 18 m / s, T 2 5 22 m / s, T 3 5 26 m / s, T 4 5 30 m / s, and T 5 5 34 m / s). The annealing of as-quenched ribbons was carried out in an evacuated quartz tube by soaking the ribbons at different temperatures for 10 min and then cooling them with water. The ribbon samples before and after annealing were characterized by X-ray diffraction (XRD) with Cu Ka radiation ( l 5 1.54 nm), electron microscopy (TEM), ¨ and 57 Fe Mossbauer spectroscopy. The structural evolution of the as-quenched ribbons was also investigated using a Perkin-Elmer 7 differential thermal analysis (DTA) system, with a heating rate of 10 K / min. Heating was carried out in flowing purified argon gas during the DTA analysis. The magnetic properties of the ribbon samples were measured with a vibrating sample magnetometer (VSM) and SQUID magnetometer (SQUID) using a maximum applied field of 2400 and 4400 kA / m, respectively.
3. Results and discussion Fig. 1 shows XRD patterns of as-quenched Pr 7 Fe 88 B 5 ribbon samples prepared by melt spinning with different copper wheel velocities (T 1 , T 2 , T 3 , T 4 , and T 5 from low to high). The sample T 1 consists of hard magnetic phase Pr 2 Fe 14 B and soft magnetic phase a-Fe, which indicates a fully crystallized ribbon was obtained at the lowest wheel speed (18 m / s). The XRD pattern of ribbon T 5 melt spun at the highest speed (34 m / s) shows only the amorphous phase. At intermediate wheel velocities, samples T 2 and T 3 consist of amorphous phase and crystalline phases of both Pr 2 Fe 14 B and a-Fe, while sample T 4 comprises mainly amorphous phase and a small amount of a-Fe. A quantitative analysis of the phases present in the samples based on ¨ Mossbauer spectra is shown in Table1. The data in Table 1 indicate that the volume fraction of amorphous phase increases with increasing quenching rate, while the volume fraction of crystalline phases decreases with increasing quenching rate. In order to study the structure evolution of overquenched ribbons (T 2 , T 3 , T 4 , and T 5 ), DTA scans were performed. Fig. 2 presents DTA results for all overquenched ribbon samples. For T 2 and T 3 , there is only one peak starting from 793 K to about 863 K for each sample. For samples T 4 and T 5 , there are two exothermic peaks: the first is a sharp one starting from 823 K to about 893 K. The second peak starts from 913 K to about 973 K.
Fig. 1. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 18 m / s (T 1 ), vs 5 22 m / s (T 2 ), vs 5 26 m / s (T 3 ), vs 5 30 m / s (T 4 ), and vs 5 34 m / s (T 5 ).
Table 1 Principal phases present in as-quenched Pr 7 Fe 88 B 5 ribbons vs (m / s)
Pr 2 Fe 14 B (vol.%)
a-Fe (vol.%)
Pr 1.1 Fe 4 B 4 (vol.%)
18 22 26 30 34
61.7 29.6 20.2
35.9 19.5 11.7 4.2
2.4
Amorphous (vol.%) 50.9 68.1 95.8 100.0
Fig. 2. DTA curves of Pr 7 Fe 88 B 5 samples melt spun with vs 5 22 m / s (T 2 ), vs 5 26 m / s (T 3 ), vs 5 30 m / s (T 4 ), and vs 5 34 m / s (T 5 ).
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Fig. 3. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 22 m / s (a) and then annealed at 873 K for 10 min (b), and 973 K for 10 min (c).
Furthermore, it can be seen that the intensity of the second peak for sample T 4 or T 5 is weaker than that of the corresponding first peak. In order to further study the phase transformation process corresponding to the DTA thermal features, isothermal annealing experiments were carried out, on all over-quenched ribbons, at temperature below or above the exothermic peaks on the DTA curves. Figs. 3–5 shows the XRD results for as-quenched samples after annealing at different temperatures for 10 min. The XRD results for
Fig. 5. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 34 m / s (a) and then annealed at 823 K (b), 893 K (c), and 973 K (d) for 10 min.
sample T 2 indicate that only Pr 2 Fe 14 B and a-Fe phases can be detected, and the amorphous phase is absent after annealing at 873 K, as shown in Fig. 3. For sample T 3 , the X-ray diffraction measurements also show that it consists of only two phases: i.e. Pr 2 Fe 14 B and a-Fe, after annealing at 873 K for 10 min. However, in the case of samples T 4 and T 5 , the a-Fe phase is found after annealing at 823 K, which is the onset temperature of the first DTA peak. There is mainly the metastable phase of TbCu 7 type and a-Fe and a small amount of Pr 2 Fe 14 B in sample T 4 , and only the 1:7 phase and a-Fe phases in sample T 5 after annealing at 893 K. The crystalline phases, after annealing at temperatures .973 K, can be identified as only Pr 2 Fe 14 B and a-Fe phases for all four samples. Based on the results in the previous paragraphs, we deduce the structure evolution of the over-quenched ribbon samples during annealing treatment as the following sequences: For T 2 and T 3 :
873 K
Am 1 Pr 2 Fe 14 B 1 a-Fe → Pr 2 Fe 14 B 1 a-
Fe For T 4 :
823 K
893 K
Am 1 a-Fe → Am9 1 a-Fe → a-Fe 1 973 K
1:7 phase 1 Pr 2 Fe 14 B → Pr 2 Fe 14 B 1 a-Fe For T 5 :
823 K
893 K
973 K
Am → Am9 1 a-Fe → 1:7 phase 1 a-Fe → P-
r 2 Fe 14 B 1 a-Fe
Fig. 4. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 30 m / s (a) and then annealed at 823 K (b), 893 K (c), and 973 K (d) for 10 min.
In the Pr–Fe–B phase diagram [10], the composition of Pr 7 Fe 88 B 5 is very near the Fe–Pr 2 Fe 14 B tie line. At equilibrium, this should lead essentially to two-phase structures of a-Fe and Pr 2 Fe 14 B phases with a small
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amount of Pr 1.1 Fe 4 B 4 or Fe 3 B. Upon annealing, the amorphous hypoeutectic Fe-rich Fe–Pr–B alloy along the Fe–Pr 2 Fe 14 B tie line should transform initially into a-Fe and then directly to Pr 2 Fe 14 B. However, what is very interesting to us is that our results clearly indicate that the phase transformation of Pr 7 Fe 88 B 5 alloy ribbons during annealing is strongly affected by the structures of the as-quenched ribbons. For the samples melt spun at wheel velocities of v 5 22 and 26 m / s, the amorphous phase transforms directly into Pr 2 Fe 14 B and a-Fe phases. But the case is totally different for the samples melt spun at higher wheel speeds. After initial crystallization of a-Fe but prior to the formation of the final mixture of Pr 2 Fe 14 B and a-Fe phases, both the metastable 1:7 phase and the Pr 2 Fe 14 B phase form for the sample melt spun at v 5 30 m / s, and only the metastable 1:7 phase occurs for the sample melt spun at v 5 34 m / s. The specific cause of these changes in phase transformation is of significant interest and needs to be studied further. The dependence of the room-temperature coercivity Hcj and remanence Jr of the samples on annealing temperature T a is shown in Fig. 6. It can be seen that for each sample both Hcj and Jr have a peak value, i.e. Hcj and Jr first increase with increasing T a , reach maximum values simultaneously and then decrease with further increasing T a . However, the temperatures at which both Hcj and Jr reached the maximum were different for different samples (923 K for samples T 2 and T 3 , 973 K for samples T 4 and T 5 ). Furthermore, it is of interest to note that all the samples after optimal annealing exhibit a two-phase structure consisting of magnetically hard Pr 2 Fe 14 B and soft
magnetic a-Fe. These results indicate that the annealing temperature should be sufficiently high to fully crystallize the amorphous phase for samples T 2 and T 3 or to decompose completely the metastable 1:7 phase for samples T 4 and T 5 , but low enough to avoid excessive grain growth. Similarly, Gabay et al. [11] have reported that the presence of the metastable 1:7 phase in Nd 2 Fe 14 B / a-Fe type nanocomposite magnets with composition of Nd 9 (Fe 12x Co x ) 85 B 6 (x 5 0–0.4) was also found to deteriorate the magnetic properties. Fig. 7 presents the demagnetization curves of overquenched samples after optimal annealing treatment. Both Hcj and Jr decrease drastically as the initial quenching rate of the samples increases from 22 to 34 m / s. Although there are two different magnetic phases in sample T 2 , its demagnetization curve exhibits single-phase magnetic behavior, i.e. no two-phase step or kink could be observed. This suggests that, in the sample, hard magnetic grains were strongly exchange-coupled with the neighboring soft grains and the magnetization vectors of hard and soft phases rotate in cluster features during magnetization reversal. However, for samples T 3 , T 4 and T 5 , an apparent kink can be seen along the demagnetization curves near zero field. This may be a consequence of some of the larger soft grains being partly or even completely decoupled from the neighboring grains and reversing independently. Moreover, the kink becomes more pronounced when the wheel speed increases from 26 (T 3 ) to 34 m / s (T 5 ). The microstructure of the annealed ribbons was investigated using TEM. Fig. 8 shows bright-field TEM micrographs of optimally annealed ribbon samples. It can be seen that the grain size of both Pr 2 Fe 14 B and a-Fe phases in the samples increases significantly with increase of the initial quenching rate. For sample T 2 (lowest quenching
Fig. 6. Variation of Hcj and Jr as a function of annealing temperature for Pr 7 Fe 88 B 5 samples melt spun with 22 m / s (a), 26 m / s (b), 30 m / s (c), and 34 m / s (d).
Fig. 7. Demagnetization curves for Pr 7 Fe 88 B 5 samples prepared at vs 5 22 m / s (a), vs 5 26 m / s (b), vs 5 30 m / s (c), and vs 5 34 m / s (d) after optimal annealing treatment.
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Fig. 8. TEM micrographs of ribbons T 2 (a), T 3 (b), T 4 (c), and T 5 (d) after optimal annealing treatment.
rate), the grains are uniformly distributed with an average grain size of around 20 nm. For sample T 5 (highest quenching rate), however, the average grain size increases to about 70 nm and the grain size distribution becomes non-uniform. Large grains up to 160 nm were present in the microstructure of sample T 5 . The microstructure changes in the annealed samples with increasing initial quenching rate (as indicated by TEM) may be determined by following two factors. One is that, in as-quenched ribbons, the volume fraction of crystallite phases, which may act as nucleation centers during crystallization of the amorphous phase, decreases with increasing initial quenching rate. Lack of nucleation sites during the crystallization of highly over-quenched ribbons could lead to a coarsegrained, inhomogeneous microstructure. Another factor is that the occurrence of the metastable phase (1:7 phase) during annealing in the ribbons melt spun at higher speeds enhances the formation temperature of the final mixture of Pr 2 Fe 14 B and a-Fe and consequently leads to grain coarsening. Previous studies [12] have shown that the magnetic properties of nanocomposite permanent magnetic materials are strongly dependent upon the grain size of hard and soft phases. Both Hcj and Jr decrease monotonously with increasing grain size, due to the weakening of the exchange coupling effect between hard and soft grains. In addition, a non-uniform microstructure is predicted to deteriorate the magnetic properties [13]. Therefore, the grain coarsening and more irregular microstructure, as
detected by TEM in Fig. 8, may be responsible for the degradation of the squareness of the demagnetization curves and the decrease in Hcj and Jr of the samples initially melt spun at high velocity. The model proposed and developed by Kronmuller et al. was used successfully to analyze the relationship between the microstructure and coercivity in nanocrystalline twophase magnets [14,15]. According to Kronmuller, the coercivity is expressed as K1 1 K2 Hc 5 aK aex ]]] 2 Neff Ms m0 Ms
(1)
where aK takes into account the reduction in anisotropy in the region near the internal surfaces as grain boundaries and interfaces, aex describes the effect of exchange coupling between neighboring grains on the coercivity force of the magnet, and Neff is related to the local stray field induced by the magnetostatic volume charges and surface charges, and describes the condition of the grain shape and surface. In melt-spun Pr–Fe–B ribbons, aK values were found to vary between 0.7 and 0.9, depending weakly on microstructure [16], so the value of aK is assumed to be 0.80 [15]. In addition, the data for the anisotropy constant K1 1 K2 and the spontaneous magnetization Ms for pure Pr 2 Fe 14 B phase can be derived from Ref. [16]. Fig. 9 presents the relationships Hcj (T ) /Ms (T ) and [K1 (T ) 1 K2 (T )] /m0 M 2s (T ) for optimally annealed T 2 , T 3 ,
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and the Pr 2 Fe 14 B phase form for the case of v 5 30 m / s, and only the metastable 1:7 phase occurs in the ribbon melt spun at v 5 34 m / s. For all ribbons, the microstructure after optimal annealing consists of only magnetically hard Pr 2 Fe 14 B and soft magnetic a-Fe phases. The optimum magnetic properties of Hcj and Jr , and the squareness of demagnetization curves of the annealed ribbons, however, deteriorate drastically with increasing quenching rate of unannealed precursor. These deterioration effects can be attributed to the formation of a coarser and more irregular microstructure during annealing in the samples initially melt spun with higher wheel speeds, which was confirmed by our TEM observations and the analysis of the temperature dependence of coercivity.
Acknowledgements Fig. 9. Hcj /Ms versus (K1 1 K2 ) /m0 M s2 plots to determine the microstructural parameters aex and Neff for optimally annealed Pr 7 Fe 88 B 5 samples.
T 4 , and T 5 ribbon samples. A linear relation is found, implying that Hcj of these samples is controlled by the nucleation of reversed domains. It can be seen that, with increasing initial quenching rate, aex decreases monotonically while Neff increases monotonically. With respect to the microstructure, the decrease in aex can be attributed to the weakening of exchange coupling between hard and soft magnetic phases due to grain coarsening. Furthermore, the increase of Neff indicates stronger demagnetization effects acting on grains. These increased effects may be attributed to the deterioration of grain shape due to the formation of sharp edges and corners. Taking into account the decrease in aex and the increase in Neff , one may conclude that the deterioration in Hcj of the samples with increasing initial quenching rate arises from both the coarser grain size and more irregular microstructures.
4. Conclusions Structure evolution during annealing treatment of Pr 7 Fe 88 B 5 ribbons melt spun at lower wheel speeds (22 or 26 m / s) takes place in one step, i.e. the amorphous phase in the as-quenched ribbons transforms directly into Pr 2 Fe 14 B and a-Fe phases. However, for the ribbons melt spun at higher wheel speeds (30 or 34 m / s), the structures evolve in totally different ways: after initial crystallization of a-Fe but prior to the formation of the final mixture of Pr 2 Fe 14 B and a-Fe phases, both the metastable 1:7 phase
Financial support from the National Natural Science Foundation of China (grant No. 59831010) is gratefully acknowledged.
References [1] E.F. Kneller, R. Hawig, IEEE Trans. Magn. 27 (1991) 3588. [2] R. Skomski, J.M.D. Coey, Phys. Rev. B 48 (1993) 15812. [3] T. Schrefl, H. Kronmuller, J. Fidler, J. Magn. Magn. Mater. 27 (1993) L273. [4] J.M.D. Coey, J. Magn. Magn. Mater. 140–144 (1995) 1040. [5] G.B. Clement, J.E. Keem, J.P. Bradley, J. Appl. Phys. 64 (1988) 5299. [6] R. Coehoorn, D.B. Mooij, J.P.W.B. Duchateau, K.H.J. Buschow, J. Phys. C 49 (1988) 669. [7] G.C. Hadjipanayis, L. Withanawasam, J. Appl. Phys. 76 (1994) 7065. [8] Z. Wang, M. Zhang, F. Li, S. Zhou, R. Wang, W. Gong, J. Appl. Phys. 81 (1997) 5097. [9] T. Hinomura, S. Nasu, H. Kanekiyo, Mater. Trans. JIM 38 (1997) 1106. [10] J.H. Tian, Y.Y. Huang, J.K. Liang, Chin. Sci. A 2 (1987) 159. [11] A.M. Gabay, A.G. Popov, V.S. Gaviko, Y.V. Belozerov, A.S. Yermolenko, N.N. Shchegoleva, J. Alloys Comp. 237 (1996) 101. [12] R. Fischer, T. Schrefl, H. Kronmuller, J. Fidler, J. Magn. Magn. Mater. 153 (1996) 35. [13] R. Fischer, T. Schrefl, H. Kronmuller, J. Fidler, J. Magn. Magn. Mater. 150 (1995) 329. [14] J. Bauer, M. Seeger, A. Zern, H. Kronmuller, J. Appl. Phys. 80 (1996) 1667. [15] D. Goll, M. Seeger, H. Kronmuller, J. Magn. Magn. Mater. 185 (1998) 49. [16] M. Seeger, H. Kronmuller, J. Magn. Magn. Mater. 137 (1994) 18.
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Phase transformations and magnetic properties of melt-spun Pr 7 Fe 88 B 5 ribbons during annealing Zuocheng Wang*, Shouzeng Zhou, Yi Qiao, Maocai Zhang, Run Wang State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China Received 29 August 1999; accepted 6 November 1999
Abstract ¨ X-ray diffraction and Mossbauer measurements indicate that melt spinning at different wheel velocities causes as-quenched Pr 7 Fe 88 B 5 ribbons to have distinctive structures. Depending on their as-quenched structures, the phase transformations of the ribbons during annealing may take place in one of the following sequences: (1) amorphous phase (Am)1Pr 2 Fe 14 B1a-Fe→Pr 2 Fe 14 B1a-Fe; (2) Am1a-Fe→Am91a-Fe→a-Fe11:7 phase1Pr 2 Fe 14 B→Pr 2 Fe 14 B1a-Fe; and (3) Am→Am91a-Fe→1:7 phase1a-Fe→ Pr 2 Fe 14 B1aFe. However, for all the ribbon samples, the microstructures after optimal annealing with respect to magnetic properties were found to only consist of two magnetic phases: Pr 2 Fe 14 B and a-Fe. The optimum magnetic properties of Hcj and Jr , and the squareness of the demagnetization curves of the annealed ribbons deteriorate drastically with increasing quenching rate of unannealed precursors. Taking into account TEM results, the values of aex and Neff derived from the temperature dependence of the coercivity, these deteriorated effects can be attributed to the formation of a coarser and more irregular microstructure during annealing in the samples initially melt spun with higher wheel speeds. 2000 Elsevier Science S.A. All rights reserved. Keywords: As-quenched structure; Phase transformation; Metastable phase; Nanocomposite magnets; Magnetic properties
1. Introduction One of the latest exciting developments in hard magnetic materials is nanocomposite permanent magnetic alloys, because they open a way for a new generation of isotropic bonded magnets with unusually high remanence, relatively high coercivity, high energy product, and low cost [1–3]. The excellent magnetic properties of these materials are known to arise from the exchange interaction between magnetically hard and soft phases. In order to achieve the most effective exchange interaction between the two phases, however, the optimal structure of these materials requires a homogeneous and narrow distribution of very fine grains, especially soft grains with diameter comparable to the domain wall width of the hard phase [1,4]. In addition, the predominant microstructure should be substantially free from intergranular phase, which would inhibit the exchange coupling between adjacent grains [5]. It is anticipated, therefore, that the magnetic
*Corresponding author. E-mail address: [email protected] (Z. Wang)
behavior of nanocomposite magnetic materials should be dependent on the intrinsic magnetic properties of hard and soft phases as well as their microstructures. So far, two types of Nd–Fe–B nanocomposite magnets, namely Fe 3 B / Nd 2 Fe 14 B and a-Fe / Nd 2 Fe 14 B [6–9], have been developed. It was found that, in order to achieve high performance, thermal processing of precursor materials, such as crystallizing overquenched ribbons, is usually needed because of the narrow range of quenching parameters over which optimum direct quench magnetic properties can be obtained. The crystallization process of the ribbons upon thermal treatment has been analyzed by different techniques [7,9]. However, in our literature, there are no studies on the influence of the initial structural properties of the overquenched precursor phase on the formation, structure and magnetic properties of the resulting nanocrystalline materials. It is established that the degree of as-cast amorphous state in the ribbons can be varied by changing the quenching of the melt during the flow-cast procedure. In this paper, we present a systematic study on microstructure evolution and magnetic properties of as-quenched amorphous and partially crystallized Pr 7 Fe 88 B 5 ribbons during annealing, and clarify how the
0925-8388 / 00 / $ – see front matter 2000 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 99 )00741-0
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259
as-quenched structures influence the crystallization process, the resulting nanocrystalline two-phase structures and their magnetic properties.
2. Experimental procedures Five ribbon samples of composition Pr 7 Fe 88 B 5 were obtained from the same master alloy by melt spinning on a copper wheel using different peripheral velocities (T 1 5 18 m / s, T 2 5 22 m / s, T 3 5 26 m / s, T 4 5 30 m / s, and T 5 5 34 m / s). The annealing of as-quenched ribbons was carried out in an evacuated quartz tube by soaking the ribbons at different temperatures for 10 min and then cooling them with water. The ribbon samples before and after annealing were characterized by X-ray diffraction (XRD) with Cu Ka radiation ( l 5 1.54 nm), electron microscopy (TEM), ¨ and 57 Fe Mossbauer spectroscopy. The structural evolution of the as-quenched ribbons was also investigated using a Perkin-Elmer 7 differential thermal analysis (DTA) system, with a heating rate of 10 K / min. Heating was carried out in flowing purified argon gas during the DTA analysis. The magnetic properties of the ribbon samples were measured with a vibrating sample magnetometer (VSM) and SQUID magnetometer (SQUID) using a maximum applied field of 2400 and 4400 kA / m, respectively.
3. Results and discussion Fig. 1 shows XRD patterns of as-quenched Pr 7 Fe 88 B 5 ribbon samples prepared by melt spinning with different copper wheel velocities (T 1 , T 2 , T 3 , T 4 , and T 5 from low to high). The sample T 1 consists of hard magnetic phase Pr 2 Fe 14 B and soft magnetic phase a-Fe, which indicates a fully crystallized ribbon was obtained at the lowest wheel speed (18 m / s). The XRD pattern of ribbon T 5 melt spun at the highest speed (34 m / s) shows only the amorphous phase. At intermediate wheel velocities, samples T 2 and T 3 consist of amorphous phase and crystalline phases of both Pr 2 Fe 14 B and a-Fe, while sample T 4 comprises mainly amorphous phase and a small amount of a-Fe. A quantitative analysis of the phases present in the samples based on ¨ Mossbauer spectra is shown in Table1. The data in Table 1 indicate that the volume fraction of amorphous phase increases with increasing quenching rate, while the volume fraction of crystalline phases decreases with increasing quenching rate. In order to study the structure evolution of overquenched ribbons (T 2 , T 3 , T 4 , and T 5 ), DTA scans were performed. Fig. 2 presents DTA results for all overquenched ribbon samples. For T 2 and T 3 , there is only one peak starting from 793 K to about 863 K for each sample. For samples T 4 and T 5 , there are two exothermic peaks: the first is a sharp one starting from 823 K to about 893 K. The second peak starts from 913 K to about 973 K.
Fig. 1. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 18 m / s (T 1 ), vs 5 22 m / s (T 2 ), vs 5 26 m / s (T 3 ), vs 5 30 m / s (T 4 ), and vs 5 34 m / s (T 5 ).
Table 1 Principal phases present in as-quenched Pr 7 Fe 88 B 5 ribbons vs (m / s)
Pr 2 Fe 14 B (vol.%)
a-Fe (vol.%)
Pr 1.1 Fe 4 B 4 (vol.%)
18 22 26 30 34
61.7 29.6 20.2
35.9 19.5 11.7 4.2
2.4
Amorphous (vol.%) 50.9 68.1 95.8 100.0
Fig. 2. DTA curves of Pr 7 Fe 88 B 5 samples melt spun with vs 5 22 m / s (T 2 ), vs 5 26 m / s (T 3 ), vs 5 30 m / s (T 4 ), and vs 5 34 m / s (T 5 ).
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Fig. 3. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 22 m / s (a) and then annealed at 873 K for 10 min (b), and 973 K for 10 min (c).
Furthermore, it can be seen that the intensity of the second peak for sample T 4 or T 5 is weaker than that of the corresponding first peak. In order to further study the phase transformation process corresponding to the DTA thermal features, isothermal annealing experiments were carried out, on all over-quenched ribbons, at temperature below or above the exothermic peaks on the DTA curves. Figs. 3–5 shows the XRD results for as-quenched samples after annealing at different temperatures for 10 min. The XRD results for
Fig. 5. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 34 m / s (a) and then annealed at 823 K (b), 893 K (c), and 973 K (d) for 10 min.
sample T 2 indicate that only Pr 2 Fe 14 B and a-Fe phases can be detected, and the amorphous phase is absent after annealing at 873 K, as shown in Fig. 3. For sample T 3 , the X-ray diffraction measurements also show that it consists of only two phases: i.e. Pr 2 Fe 14 B and a-Fe, after annealing at 873 K for 10 min. However, in the case of samples T 4 and T 5 , the a-Fe phase is found after annealing at 823 K, which is the onset temperature of the first DTA peak. There is mainly the metastable phase of TbCu 7 type and a-Fe and a small amount of Pr 2 Fe 14 B in sample T 4 , and only the 1:7 phase and a-Fe phases in sample T 5 after annealing at 893 K. The crystalline phases, after annealing at temperatures .973 K, can be identified as only Pr 2 Fe 14 B and a-Fe phases for all four samples. Based on the results in the previous paragraphs, we deduce the structure evolution of the over-quenched ribbon samples during annealing treatment as the following sequences: For T 2 and T 3 :
873 K
Am 1 Pr 2 Fe 14 B 1 a-Fe → Pr 2 Fe 14 B 1 a-
Fe For T 4 :
823 K
893 K
Am 1 a-Fe → Am9 1 a-Fe → a-Fe 1 973 K
1:7 phase 1 Pr 2 Fe 14 B → Pr 2 Fe 14 B 1 a-Fe For T 5 :
823 K
893 K
973 K
Am → Am9 1 a-Fe → 1:7 phase 1 a-Fe → P-
r 2 Fe 14 B 1 a-Fe
Fig. 4. X-ray diffraction patterns of Pr 7 Fe 88 B 5 samples melt spun with vs 5 30 m / s (a) and then annealed at 823 K (b), 893 K (c), and 973 K (d) for 10 min.
In the Pr–Fe–B phase diagram [10], the composition of Pr 7 Fe 88 B 5 is very near the Fe–Pr 2 Fe 14 B tie line. At equilibrium, this should lead essentially to two-phase structures of a-Fe and Pr 2 Fe 14 B phases with a small
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amount of Pr 1.1 Fe 4 B 4 or Fe 3 B. Upon annealing, the amorphous hypoeutectic Fe-rich Fe–Pr–B alloy along the Fe–Pr 2 Fe 14 B tie line should transform initially into a-Fe and then directly to Pr 2 Fe 14 B. However, what is very interesting to us is that our results clearly indicate that the phase transformation of Pr 7 Fe 88 B 5 alloy ribbons during annealing is strongly affected by the structures of the as-quenched ribbons. For the samples melt spun at wheel velocities of v 5 22 and 26 m / s, the amorphous phase transforms directly into Pr 2 Fe 14 B and a-Fe phases. But the case is totally different for the samples melt spun at higher wheel speeds. After initial crystallization of a-Fe but prior to the formation of the final mixture of Pr 2 Fe 14 B and a-Fe phases, both the metastable 1:7 phase and the Pr 2 Fe 14 B phase form for the sample melt spun at v 5 30 m / s, and only the metastable 1:7 phase occurs for the sample melt spun at v 5 34 m / s. The specific cause of these changes in phase transformation is of significant interest and needs to be studied further. The dependence of the room-temperature coercivity Hcj and remanence Jr of the samples on annealing temperature T a is shown in Fig. 6. It can be seen that for each sample both Hcj and Jr have a peak value, i.e. Hcj and Jr first increase with increasing T a , reach maximum values simultaneously and then decrease with further increasing T a . However, the temperatures at which both Hcj and Jr reached the maximum were different for different samples (923 K for samples T 2 and T 3 , 973 K for samples T 4 and T 5 ). Furthermore, it is of interest to note that all the samples after optimal annealing exhibit a two-phase structure consisting of magnetically hard Pr 2 Fe 14 B and soft
magnetic a-Fe. These results indicate that the annealing temperature should be sufficiently high to fully crystallize the amorphous phase for samples T 2 and T 3 or to decompose completely the metastable 1:7 phase for samples T 4 and T 5 , but low enough to avoid excessive grain growth. Similarly, Gabay et al. [11] have reported that the presence of the metastable 1:7 phase in Nd 2 Fe 14 B / a-Fe type nanocomposite magnets with composition of Nd 9 (Fe 12x Co x ) 85 B 6 (x 5 0–0.4) was also found to deteriorate the magnetic properties. Fig. 7 presents the demagnetization curves of overquenched samples after optimal annealing treatment. Both Hcj and Jr decrease drastically as the initial quenching rate of the samples increases from 22 to 34 m / s. Although there are two different magnetic phases in sample T 2 , its demagnetization curve exhibits single-phase magnetic behavior, i.e. no two-phase step or kink could be observed. This suggests that, in the sample, hard magnetic grains were strongly exchange-coupled with the neighboring soft grains and the magnetization vectors of hard and soft phases rotate in cluster features during magnetization reversal. However, for samples T 3 , T 4 and T 5 , an apparent kink can be seen along the demagnetization curves near zero field. This may be a consequence of some of the larger soft grains being partly or even completely decoupled from the neighboring grains and reversing independently. Moreover, the kink becomes more pronounced when the wheel speed increases from 26 (T 3 ) to 34 m / s (T 5 ). The microstructure of the annealed ribbons was investigated using TEM. Fig. 8 shows bright-field TEM micrographs of optimally annealed ribbon samples. It can be seen that the grain size of both Pr 2 Fe 14 B and a-Fe phases in the samples increases significantly with increase of the initial quenching rate. For sample T 2 (lowest quenching
Fig. 6. Variation of Hcj and Jr as a function of annealing temperature for Pr 7 Fe 88 B 5 samples melt spun with 22 m / s (a), 26 m / s (b), 30 m / s (c), and 34 m / s (d).
Fig. 7. Demagnetization curves for Pr 7 Fe 88 B 5 samples prepared at vs 5 22 m / s (a), vs 5 26 m / s (b), vs 5 30 m / s (c), and vs 5 34 m / s (d) after optimal annealing treatment.
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Fig. 8. TEM micrographs of ribbons T 2 (a), T 3 (b), T 4 (c), and T 5 (d) after optimal annealing treatment.
rate), the grains are uniformly distributed with an average grain size of around 20 nm. For sample T 5 (highest quenching rate), however, the average grain size increases to about 70 nm and the grain size distribution becomes non-uniform. Large grains up to 160 nm were present in the microstructure of sample T 5 . The microstructure changes in the annealed samples with increasing initial quenching rate (as indicated by TEM) may be determined by following two factors. One is that, in as-quenched ribbons, the volume fraction of crystallite phases, which may act as nucleation centers during crystallization of the amorphous phase, decreases with increasing initial quenching rate. Lack of nucleation sites during the crystallization of highly over-quenched ribbons could lead to a coarsegrained, inhomogeneous microstructure. Another factor is that the occurrence of the metastable phase (1:7 phase) during annealing in the ribbons melt spun at higher speeds enhances the formation temperature of the final mixture of Pr 2 Fe 14 B and a-Fe and consequently leads to grain coarsening. Previous studies [12] have shown that the magnetic properties of nanocomposite permanent magnetic materials are strongly dependent upon the grain size of hard and soft phases. Both Hcj and Jr decrease monotonously with increasing grain size, due to the weakening of the exchange coupling effect between hard and soft grains. In addition, a non-uniform microstructure is predicted to deteriorate the magnetic properties [13]. Therefore, the grain coarsening and more irregular microstructure, as
detected by TEM in Fig. 8, may be responsible for the degradation of the squareness of the demagnetization curves and the decrease in Hcj and Jr of the samples initially melt spun at high velocity. The model proposed and developed by Kronmuller et al. was used successfully to analyze the relationship between the microstructure and coercivity in nanocrystalline twophase magnets [14,15]. According to Kronmuller, the coercivity is expressed as K1 1 K2 Hc 5 aK aex ]]] 2 Neff Ms m0 Ms
(1)
where aK takes into account the reduction in anisotropy in the region near the internal surfaces as grain boundaries and interfaces, aex describes the effect of exchange coupling between neighboring grains on the coercivity force of the magnet, and Neff is related to the local stray field induced by the magnetostatic volume charges and surface charges, and describes the condition of the grain shape and surface. In melt-spun Pr–Fe–B ribbons, aK values were found to vary between 0.7 and 0.9, depending weakly on microstructure [16], so the value of aK is assumed to be 0.80 [15]. In addition, the data for the anisotropy constant K1 1 K2 and the spontaneous magnetization Ms for pure Pr 2 Fe 14 B phase can be derived from Ref. [16]. Fig. 9 presents the relationships Hcj (T ) /Ms (T ) and [K1 (T ) 1 K2 (T )] /m0 M 2s (T ) for optimally annealed T 2 , T 3 ,
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and the Pr 2 Fe 14 B phase form for the case of v 5 30 m / s, and only the metastable 1:7 phase occurs in the ribbon melt spun at v 5 34 m / s. For all ribbons, the microstructure after optimal annealing consists of only magnetically hard Pr 2 Fe 14 B and soft magnetic a-Fe phases. The optimum magnetic properties of Hcj and Jr , and the squareness of demagnetization curves of the annealed ribbons, however, deteriorate drastically with increasing quenching rate of unannealed precursor. These deterioration effects can be attributed to the formation of a coarser and more irregular microstructure during annealing in the samples initially melt spun with higher wheel speeds, which was confirmed by our TEM observations and the analysis of the temperature dependence of coercivity.
Acknowledgements Fig. 9. Hcj /Ms versus (K1 1 K2 ) /m0 M s2 plots to determine the microstructural parameters aex and Neff for optimally annealed Pr 7 Fe 88 B 5 samples.
T 4 , and T 5 ribbon samples. A linear relation is found, implying that Hcj of these samples is controlled by the nucleation of reversed domains. It can be seen that, with increasing initial quenching rate, aex decreases monotonically while Neff increases monotonically. With respect to the microstructure, the decrease in aex can be attributed to the weakening of exchange coupling between hard and soft magnetic phases due to grain coarsening. Furthermore, the increase of Neff indicates stronger demagnetization effects acting on grains. These increased effects may be attributed to the deterioration of grain shape due to the formation of sharp edges and corners. Taking into account the decrease in aex and the increase in Neff , one may conclude that the deterioration in Hcj of the samples with increasing initial quenching rate arises from both the coarser grain size and more irregular microstructures.
4. Conclusions Structure evolution during annealing treatment of Pr 7 Fe 88 B 5 ribbons melt spun at lower wheel speeds (22 or 26 m / s) takes place in one step, i.e. the amorphous phase in the as-quenched ribbons transforms directly into Pr 2 Fe 14 B and a-Fe phases. However, for the ribbons melt spun at higher wheel speeds (30 or 34 m / s), the structures evolve in totally different ways: after initial crystallization of a-Fe but prior to the formation of the final mixture of Pr 2 Fe 14 B and a-Fe phases, both the metastable 1:7 phase
Financial support from the National Natural Science Foundation of China (grant No. 59831010) is gratefully acknowledged.
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