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Solidification, microstructural refinement and magnetism in Nd2Fe14B
Journal of Magnetism and Magnetic Materials 241 (2002) 144–155
Topical Review
Solidification, microstructural refinement and magnetism in Nd2Fe14B M.J. Kramera,*, L.H. Lewisb, L.M. Fabiettic, Y. Tanga,1, W. Millerb,2, K.W. Dennisa, R.W. McCalluma a
USDOE and Department of Materials Science and Engineering, Ames Laboratory, Iowa State University, 37 Wilhelm Hall, Ames, IA 50011, USA b Materials Science Department, Brookhaven National Laboratory, Upton, NY 11973-5000, USA c Facultad de Matematica Astronomia y Fisica, Ciudad Universitaria, Cordoba 5000, Argentina Received 25 May 2001; received in revised form 24 October 2001
Abstract Nanocrystalline Nd2Fe14B (2-14-1) exhibits very favorable properties conferred by rapid solidification processes, including: simplified processing, good corrosion resistance and high magnetic hardness provided by a resultant grain size that is on the order of the single-domain size. While the ideal microstructural state of melt-spun 2-14-1 is a homogeneous dispersion of nanocrystallites embedded within a vitrified matrix, in practice it is very difficult to control the microstructural development to a uniform nanocrystalline state. This as-yet unmet challenge underlies the desire to understand, and ultimately to control, the solidification path taken by Nd–Fe–B-based melts during rapid solidification processing in order to obtain a homogeneous microstructure. Processing factors that influence the homogeneity and hence magnetic performance of the rapidly quenched microstructure of Nd2Fe14B are evaluated and categorized. These results emphasize both the complexities of rapid solidification processing as well as the potential for control of a highly chaotic processing technique. A qualitative solidification model that incorporates recalesence is presented. This model details the evolution of microstructures related to the thermal gradient as well as the phase selection process undertaken during solidification of rapidly solidified Nd–Fe–B alloys. The model is novel in that it explains the observed microstructures, from nanophased equiaxed grains transitioning to large elongated grains, and provides both diagnostic and predictive information concerning rapid solidification processing methods. r 2002 Elsevier Science B.V. All rights reserved. Keywords: Solidification; Peritectic; Recalesence; Coercivity
1. Introduction *Corresponding author. Tel.: +1-515-294-0276; fax: +1515-294-4291. E-mail address: [email protected] (M.J. Kramer). 1 Now at Argonne National Laboratory, 9700 South Cass Ave. Argonne, IL 60439, USA. 2 Now at University of Pennsylvania School of Engineering & Applied Science, 113 Towne Building 220 So. 33rd Street, Philadelphia, PA 19104-6391, USA.
Rare-earth-based permanent magnets, with coercivities (H ci ) generally exceeding 5000 Oe, are found in a wide variety of applications with demand for rare-earth permanent magnets growing by 10–15% per annum [1]. Today, the permanent magnet of choice for most advanced
0304-8853/02/$ - see front matter r 2002 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 0 1 ) 0 0 9 5 5 - 6
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applications is based on the intermetallic compound Nd2Fe14B (2-14-1). It is anticipated that rapid solidification processing will assume greater importance in the near future as it is the main technique used to produce bulk magnets with a nanoscale microstructure. The extremely small scale of microstructures produced by rapid solidification techniques approaches that of the domain wall widths in these materials [2] and thus it is possible for exchange interactions, in addition to the ever-present magnetostatic interactions, to magnetically couple majority and minority phases. Such effects can create magnetic mesostructures that consist of many exchange-coupled grains, producing ‘‘interaction domains’’ [3–5] that are much larger than the individual crystallographically defined grains. In some instances the intergranular exchange coupling is exploited to create two-phase ‘‘exchange–spring’’ magnets [6] that consist of a magnetically soft phase intimately mixed with a hard phase to form a magnetic composite that possesses the best properties of each constituent phase. Exchange–spring magnets are candidates for the next generation of permanent magnets whose properties can be optimized through better control of processing. The goal is to attain a better fundamental understanding of the factors that influence the microstructure, and hence magnetic properties, of rapidly solidified advanced magnetic materials. Melt–spinning is a widely used production method to produce rapidly solidified materials with a metastable microstructure [7,8]. Unlike traditional casting techniques which result in significant amounts of macrosegregation, the high cooling rates obtained in melt-spinning are typically believed to eliminate macrosegregation during solidification. However, in the Nd–Fe–B system significant microsegregation still occurs in rapidly solidified ribbons of Nd2Fe14B [9–11]. The grain size from the wheel side (chill surface) to the free side can vary by 2–3 orders of magnitude in a region less than 75 mm through the ribbons thickness [10]. The existence of processing-induced microstructural inhomogeneities in Nd2Fe14B, and the associated adverse effect upon the magnetic properties, have been known for many years [12]. However, few attempts have been made to system-
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atically identify, categorize, model and compensate for these defects in this family of materials. It is empirically recognized that microsegregation can be minimized by undercooling the melt below the glass transition temperature Tg where conventional nucleation and growth processes, which lead to segregation, can be eliminated. However, heat flow limitations and stochastic effects frustrate the complete transformation of the melt into a uniform glass in ways that are not completely understood. The magnetic properties of rapidly solidified materials are especially sensitive to microstructure where control to the nanometer length scale is necessary. While the ideal microstructural state of melt-spun Nd–Fe–B is a homogeneous dispersion of nanocrystallites embedded within a vitrified matrix, in practice it is very difficult to control the structural development to the nanocrystalline state. This challenge underlies the desire to understand, and ultimately to control, the solidification path taken by Nd–Fe– B-based melts during rapid solidification processing. Recent results, outlined in this paper, obtained from a wide variety of experimental techniques underscore the dependence of the phase selection and microstructure of the final meltquenched Nd–Fe–B-based product on the condition of the as-quenched state. A solidification model is presented which qualitatively explains the wide range of microstructures observed and the phase selection process undertaken during solidification in rapidly quenched Nd–Fe–B alloys. The model explains the observed microstructures, from nanophased equiaxed grains transitioning to large elongated grains. These results emphasize both the complexities of the solidification in this system as well as the potential for controlling the highly chaotic melt-spinning processing technique.
2. Approach 2.1. Sample preparation Ingots of Nd–Fe–B with initial compositions of stoichiometric Nd2Fe14B were prepared by arcmelting in an Ar atmosphere. Charges of approximately 10 g were loaded in quartz tubes,
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inductively melted to about 13751C and ejected through an orifice of 0.8 mm in diameter with a 150 Torr Ar or He over-pressure onto a single rotating Cu wheel or a double roller system with stainless steel quench wheels. Melt-spinning was performed at wheel speeds from 5–40 m/s, providing resultant microstructures which spanned the range of crystallinity from coarsely crystalline at low speeds (under-quenched) to nanophased at intermediate wheel speeds (optimally quenched) and amorphous at high wheel speeds (overquenched). The chamber pressure and gas were 750 Torr of Ar or He, 250 Torr He and active vacuum [13].
2.2. Characterization 2.2.1. Thermal analyses Thermal analyses were performed on asquenched samples to assess the microstructural stability and to evaluate the effect of quenching atmosphere on the melt-spun microstructure. Differential thermal analysis (DTA) was carried out with a Perkin Elmer System 7 in flowing ultrahigh purity Ar using a flow rate of 50 cc/min and a heating rate of 101C/min. Differential scanning calorimetry (DSC) was done on a Perkin Elmer Pyris DSC at a scanning rate of 401C/min. Samples for thermal analysis were ground from melt-spun ribbon to o50 mm powder in a N2 glove box.
2.2.2. X-ray diffraction (XRD) The evolution of the ribbon microstructure was monitored across the thickness from the bottom of the ribbon that was in contact with the quenching wheel surface (‘‘wheel side’’) to the top of the ribbon in contact with the quenching atmosphere (‘‘free side’’). To determine the degree of variability in microstructure from wheel side to free side, X-ray diffraction was performed on 10–20 ribbon fragments placed on a glass slide using double-sided tape with the ribbons’ ‘‘wheel side surface’’ facing up or down. Standard 2y scans were then performed from 101 to 1001 with step size of 0.051 using Cu Ka radiation.
2.2.3. Microscopies An in-depth study of the quenched ribbon microstructure was carried out with electron and atomic force microscopies. Transmission electron microscopy (TEM) was performed on crosssection specimens of the ribbons. The samples were prepared by inserting several ribbons into 3 mm quartz tubes which were then filled with epoxy, cured then sliced into 3 mm disks. The disks were mechanically polished to about 20 mm thick then Ar ion milled at 5 kV with a LN2 cooled stage to perforation. Planar TEM samples were prepared by preferential polishing and ion milling from one side. TEM evaluations were conducted on Philips CM30 transmission electron microscope operated at 300 keV. Atomic/magnetic force microscopy (AFM/MFM) was used to examine the cross-sectional topographic and magnetic structure of the ribbons. AFM/MFM images were obtained with a Digital Instruments Nanoscope III AFM operated in tapping mode using commercial standard magnetic tips (MESP, Digital Instruments Inc.) which provided the best magnetic contrast for these samples. Each ribbon was examined with AFM/MFM in more than 10 regions to ensure a representative image of the ribbons. Representative examples of the crosssectional microstructure and associated magnetic domain structures are provided in Fig. 1. In Fig. 1(a) magnetic domains of opposite polarization are delineated by dark domain walls with a domain contrast that varies from grain to grain. In Fig. 1(b) the magnetic domains appear in light and dark contrast, and the grain contrast is more difficult to discern. The difference in the domain and grain contrast between Fig. 1(a) and (b) is due to the type of MFM tip utilized during the experiments. 2.2.4. Magnetic characterization Extrinsic technical magnetic properties such as remanence Br and coercivity Hci are highly dependent upon microstructural details. Inhomogeneities in the microstructure of rapidly solidified Nd2Fe14B ribbons give rise to a variety of magnetic phases with diverse technical magnetic properties. This is despite the fact that the overall chemical composition and sometimes crystal
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tion of the number and type of magnetic phases in the sample provides diagnostic information concerning the homogeneity of the sample. To this end, hysteresis loop measurements were performed using a Quantum Design SQUID magnetometer operated at 300 K. The room-temperature demagnetization data were fitted to an empirical model using the general curve fit option in KaleidaGraph software that delineates the ferromagnetic components present in the data. The utilized curve fit is empirical, not derived from first principals, and is not at all intended provide insight into crystallographic alignment, anisotropy or magnetic coupling, which are extremely challenging parameters to model, especially in a multi-phase alloy. MðHÞ ¼
X 2M S
F
n
Fig. 1. MFM images of a polished cross-sections for (a). 2-14-1 melt spun at 10 m/s in 750 Torr He and (b). 15 m/s in 250 Torr He. The lower wheel speed ribbons are thicker and show only a limited region of small magnetic domains. Both images are the same magnification.
structure of these phases may be identical. For instance, depending upon the processing details it is possible to find in a quenched ribbon glassy or amorphous Nd–Fe–B as well as Nd2Fe14B and aFe in their nanocrystalline, single-domain-type and coarse micron-scale forms. Although each of these magnetic phases provides a unique coercivity signature, the saturation magnetization value is not a diagnostic parameter because of the uncertainty in the volume fractions of the different magnetic phases. In general, magnetic phases with low coercivity and low remanence are attributed to amorphous Nd–Fe–B or large grains of Nd2Fe14B. High-coercivity magnetic phases likely consist of single-domain-type nanocrystalline Nd2Fe14B, while a low-coercivity yet high remanence signature may be attributed to micron-scale Nd2Fe14B or to a-Fe. Despite the uncertainty in the actual identity of the magnetic phase, broad quantifica-
p
tan1
ðH7Hc Þ pS tan : Hc 2
ð1Þ
The above equation represents a hysteresis curve of n ferromagnetic magnetic phases of saturation magnetization MFS ; with the offset along the field H axis identified by the coercivity Hc . S is the squareness factor, which is the ratio of the remanent magnetization BR ðH ¼ 0Þ to the saturation magnetization MFS [14]. In this manner, the number of magnetic phases in the samples are assessed and presented in Table 1. Weight fractions of different magnetic phases are delineated by their coercivity value and are listed in Table 1 as p1; p2 and p3 in order of decreasing phase coercivity. The percentage of a given magnetic phase was determined by its moment fraction. This representation of the hysteresis curve involves no assumptions concerning crystallographic texture or interphase exchange coupling present in the samples. Fig. 2 is a representative experimental demagnetization curve (data markers) and the constituent ferromagnetic phases as determined by this curve fitting. It is noted that there are three ferromagnetic phases, of different remanence and coercivity, present in this sample; this determination correlates well with the microstructural data presented in Section 3.2. The results obtained from the room-temperature magnetic characterization are presented in Table 1.
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Table 1 Results of the decomposition of room-temperature hysteresis loops into constituent magnetic phases p1, p2, p3, given in order of magnetic hardness. MS is the saturation magnetization and HC is the coercivity. The percentage of a given phase is determined by its moment fraction Quenching atmosphere
Wheel speed (m/s)
MS (p1) (emu/g)
% p1
1 atm Ar
5 10 15 20 30
57 68 72 38 19
56 67 69 39 15
978 2842 5011 9215 6790
45 31 33 51 53
5 10 15 20 30
47 75 49 57 20
42 75 46 47 16
1755 1328 7036 8453 10050
65 25 35 26 106
5 10 15 20 30
58 23 12 47.5 40
37 23 11 39 34
525 14000 8944 9290 7466
50 50 55 27 15
1 atm He
1/3 atm He
HC (p1) (Oe)
MS (p2) (emu/g)
% p2
44 33 31 52 42
HC (p2) (Oe) 271 505 730 852 80
68 25 30 21.5 84
395 354 1516 1462 202
32 51 51 22 13
215 5068 3095 4332 114
MS (p3) (emu/g)
% p3
HC (p3) (Oe)
19 54
19 43
12 10
23.5 38
22 31
627 20
25.5 41 47 62
26 38 39 53
815 632 10 10
3. Results
Fig. 2. Decomposed hysteresis loop of stoichiometric Nd2Fe14B melt-spun at 20 m/s in 1 atm Ar. The constituent magnetically hard and magnetically soft phases are shown and are identified in the figure legend.
The compilation and analysis of the experimental data obtained on the microstructural, phase constitution and magnetic properties of the rapidly solidified Nd–Fe–B-based ribbons demonstrate the complex interaction of microstructures and properties. They reveal the importance of the quenched-in nanoscale-structure of the melt-spun ribbons in the determination of the degree of homogeneity, and hence the magnetic performance, of the ribbons. Representative microstructures are illustrated in Fig. 3. As expected, exothermic data derived from thermal analyses indicate that a larger glass fraction is present in ribbons quenched at higher wheel speeds (Fig. 4). This conclusion is broadly supported by the magnetic and microstructural investigations. However, significant differences in the quenched microstructures are produced by small changes in the processing variables other than the wheel speed.
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Fig. 3. TEM micrographs taken at various distances from the ribbon wheel side on a sample that was obliquely thinned (i.e., polished surface B201 off the surface normal). (a) Regions near the wheel side showed nearly circular coarse-grained 2-14-1 undergoing what appears to be dynamic recrystallization, with the nanophased matrix increasing in grain size with distance from the approximately cellular 2-14-1 grains. (b) Slightly larger but more uniform 2-14-1 grains and occasional smaller Fe grains. (c) Coarser 2-14-1 with nearly continuous Fe grain-boundary phase. (d) Large-grained matrix of 2-14-1 with aligned Fe cells.
Fig. 4. Results of DSC measurements for stoichiometric 2-14-1 melt spun from 5 to 40 m/s under varying chamber pressures. While the sample quenched in an atmosphere of 250 Torr He possessed the largest glass fraction based on the area under the curve of the DSC exotherm during devitrification, the sample quenched under active vacuum had the lowest glass fraction.
The overall microstructure of the melt-spun ribbons can be grouped into three broad categories. The first category is the under-quenched state, where there exists a substantial fraction of grains with diameters greater than the optimal (B100 nm) grain size necessary for good hard magnetic properties. The second category is the optimally quenched state where the majority of the grains have a near-optimal grain size (B20 nm) to support a high coercivity, and the third category is the over-quenched state where a sizable volume fraction of the material is magnetically soft because the matrix is predominantly amorphous. The appropriate circumferential quenching wheel speeds where these transition points are observed depend upon the alloy composition, the thermal conductivity of the quench wheel, chamber gas
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and pressure and the temperature of the melt prior to rapid solidification (i.e., superheat) of the alloy [15]. 3.1. Under-quenched microstructure The microstructure of an under-quenched 2-141 ribbon is highly variable both through its thickness and along its length. Optical microscopy, SEM and TEM studies on these ribbons confirmed that the microstructure varies from nanophased to coarsely crystalline (grain size dgrain > 10 mm) [10]. Typically, underquenched microstructures are obtained in the Nd–Fe–B system with low wheel speeds (5–10 m/s) and display two distinct magnetic phases with coercivities in the intermediate (Hci o5000 Oe) and low (Hci o700 Oe) ranges. The cross-sectional microstructure has four distinct layers, shown in more detail by TEM (Fig. 3) and partially observable in the MFM images (Fig. 1). The wheel side layer (in Fig. 3(a)) is thin (B0.3 mm) and consists of fine equiaxed 2-14-1 grains and larger cells of 2-14-1, which have apparently undergone recrystallization. A thicker region follows, up to 20 mm thick in some samples, where the equiaxed 2-14-1 grains coarsen systematically from less than 100 nm in size to about 300 nm (Fig. 1(b)). In this region smaller a-Fe particles frequently exist at triple conjunctions of the 2-14-1 grains (in Fig. 3(b)), yet no grain boundary phase has been observed. Further away from the wheel side surface, the grains coarsen and achieve sizes up to 1 mm or more in diameter. As the grains coarsen, occasional a-Fe grains are seen to coalesce and give rise to a thin continuous layer of an a-Fe intergranular phase (Fig. 3(c)). This intermediate layer gives way to a microstructure composed of large equiaxed 2-14-1 grains (dgrain > 10 mm) (Fig. 2) with aligned a-Fe inclusions that exist close to the free side in single roller experiments and in the ribbon center in double roller experiments [16] (Fig. 3(d)). Texturing of the 2-14-1 with the c-axis parallel to the ribbon surface has been observed in both single and twin roller experiments [17–20]. While the texturing has been attributed to columnar growth, this explanation is at odds with preferred basal plane growth observed in single crystals [21].
It should be emphasized that the demarcation between these zones and the coercivity values of these layers are approximate and the thickness of these various layers is dependent upon processing conditions. SEM and MFM images over a large number of cross-sections for samples processed in varying chamber atmospheres and pressures demonstrate that in general the thickness of the coarse grained regions decreased with increasing wheel speed and decreasing the chamber pressure to 250 Torr He. 3.2. Optimally quenched microstructure An increase in the quenching wheel speed produces an increase in the volume fraction of the glassy phase as measured by the increase in the total heat evolved during crystallization in the DSC (Fig. 4) [13]. In addition, the remaining crystalline microstructure becomes finer with increasing wheel speed as indicated by TEM investigation and Bragg peak broadening in the XRD data. For wheel speeds in the range of 15– 20 m/s and a chamber atmosphere of 750 Torr Ar, a near optimally quenched microstructure with grain sizes in the range of 80–100 nm dominates. Typical coercivity values of the hard magnetic phase in optimally processed ribbons are in the range 8–14 kOe, although these materials still exhibit two or more magnetic phases of lower coercivity. Reducing the chamber pressure to 250 Torr (0.333 atm) He allowed the wheel speed necessary to achieve an optimal microstructure to be reduced by B5 m/s, Table 1 (Fig. 4). Assuming the existence of a positive thermal gradient across the ribbon thickness, the nucleation rate (either heterogeneous or homogeneous) is dependent on the degree of undercooling and should vary through the ribbon thickness, resulting in a varying grain size through the ribbon thickness. This microstructure is demonstrated in Fig. 5 where the wheel side of the ribbon is amorphous as determined by electron diffraction in the TEM. Yet the cross-sectioned sample examined by TEM reveals a gradual increase in the size and density of the nanocrystalline particles towards the free surface of the ribbon, along the main thermal gradient. Deconstruction of the
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Fig. 5. TEM cross-section of 22.5 m/s stoichiometric 2-14-1 melt spun under 750 Torr He which shows an increase in proportion of nanophased 2-14-1 and a reduction in the glass fraction from (a) wheel side to (b) free side. The insets are selected area diffraction patterns taken in the region of the bright field images.
quenched ribbons, a-Fe particles are still frequently observed at triple points of 2-14-1 grains. The presence of such a highly dispersed a-Fe phase is inconsistent with homogeneous nucleation of the stoichiometric composition and suggests competing energetics between nucleating Fe and 2-14-1 phases. 3.3. Over quenched microstructure
Fig. 6. TEM micrograph of an over quenched ribbon showing small nuclei are 2-14-1 grains, as confirmed by convergent beam diffraction. (inset upper left).
hysteresis loops measured on the same samples demonstrates that the optimal hard magnetic fraction varies systematically with wheel speed and chamber pressure as well as supports the observation that mixed microstructures occur at all wheel speeds (Table 1). The maximum in the hard magnetic phase fraction agrees well with the microstructural observations of predominately nanophased 2-14-1. As found in the under-
Further increases in the circumferential wheel speed produces microstructures with much finer crystalline 2-14-1 grains (on the order of 20 nm or less) and larger phase fractions of Nd–Fe–B glass of very low coercivity Hci o50 Oe (Table 1). For ribbons melt spun at 30 m/s, about 70 wt% (which is close to 70 vol%) of the microstructure is amorphous. Other studies [22] have demonstrated that amorphization is nearly complete for wheel speeds in excess of 30 m/s. A reduced chamber pressure reduces the quench rate necessary to achieve a substantially glassy matrix [13] as identified by XRD, TEM selected area diffraction patterns (SADP) and thermal analysis techniques. Meanwhile, within nominally glassy regions, extremely small nuclei, with diameters dp5 nm have been observed using conical bright field (BF) and dark field (DF) TEM techniques performed on the same region. Utilization of the convergent beam electron diffraction (CBED) TEM technique revealed the nuclei to consist of the Nd2Fe14B phase (Fig. 6). It should be noted that for our
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melt-spinner at high wheel speeds, the contact between the melt pool and the wheel is destabilized. This was most likely due to wheel vibrations, resulting in regions of entrained gas where the quenching rate is lower than expected. While these regions are occasionally observed in the TEM images, the DTA and XRD results show that the bulk of the alloy melt-spun at wheel speeds above 20 m/s in 250 Torr He is amorphous. The amorphous nature of the alloy is also confirmed by the high volume fraction of the low coercivity component for samples melt-spun above 20 m/s (Table 1).
4. Discussion There is a considerable body of literature on rapidly solidified R2Fe14B (R=Pr, Nd, Tb, Dy) based alloys. Most of the research has concentrated on altering the alloy composition to enhance certain magnetic properties, i.e., the Curie temperature Tc ; ðBHÞmax ; Hci ; etc. [12]. Only a small portion of the research has concentrated on understanding the effects of solidification on phase selection [10,23]. Yet, the magnetic properties are intimately connected to the as-quenched microstructures since local phase assemblages, as well as the overall chemistry, influence the crystallization pathway. The magnetic properties, in turn, are linked to the size and distribution of the primary 214-1 matrix and the secondary phase distribution. For instance, Table 1 shows that coercivity peaks in the region of moderate quenching wheel speeds (10–15 m/s) and low chamber pressures of helium. In order to optimize the technical magnetic properties, better understanding and control of the solidification process must be obtained. The effects of major processing parameters on the quenched microstructure are discussed below. 4.1. Effect of atmosphere The species and pressure of the quenching gas also have a significant effect on the homogeneity, microstructure and phase constitution of the quenched ribbon. In general it is found that a low pressure of quenching gas possessing a
relatively high specific heat delivers the best and most uniform quench [13]. This conclusion is supported by detailed analysis of the longitudinal cross-sections of the ribbons processed in 750 Torr chamber pressure which reveals rather startling microstructural inhomogenities for all wheel speeds. Profilometer measurements of the wheel side topography indicate the increased presence of elongated regions of deep surface depressions up to 18 mm in depth throughout the wheel surface [13,24]. These regions were elongated in a direction parallel to the direction of the tangential motion of the quench wheel. SEM observations revealed that the regions on the ribbon free side directly above the wheel-side depressions show evidence of coarsely crystalline material for nearly all wheel speeds. Since it is suspected that these poorly quenched regions arose from poor contact between the melt pool and the wheel surface due to trapped gas pockets [13], DSC (Fig. 4), XRD and profilometer measurements were repeated on ribbons melt-spun from 5 to 30 m/s in the atmospheres of 250 Torr of He and an active vacuum. The wheel surface appeared both visually and by profilometer measurements to be much smoother, with depressions of less than 5 mm in depth. XRD demonstrates that the formation of nanograined 214-1 was enhanced while DSC shows that the relative proportion of glass was increased by lowering the chamber pressure to 250 Torr He, but not to an active vacuum. In fact, vacuumprocessed ribbons had the largest fraction of crystalline material as compared to Ar- or Heprocessed material at any wheel speed regardless of the chamber pressure. The explanation for poor quenching in vacuum is quite simple: as the ribbons leave the wheel glowing (in excess of 7001C), the final cooling of the ribbon occurs by conduction/convection into the surrounding environment during free flight, a mechanism which is minimally operative in vacuum but is still quite effective in 250 Torr He. 4.2. Heat flow considerations Achievement of through-thickness homogeneity in the ribbon microstructure appears to be limited by unidirectional heat flow through the solidified
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portion of the ribbon to the Cu quench wheel. Local perturbations to the overall heat flow during solidification have a dramatic effect on the phase content of the quenched product, as discussed above for the under-quenched alloys. Reduction in the chamber pressure decreases the occurrence of the trapped gas between the liquid pool and quench wheel, a phenomenon which locally impedes heat flow to the wheel; the use of He gas as a quenching medium improves the quench by virtue of its higher heat capacity. However, even with the resultant improved uniformity of the quench along the ribbon surface obtained by meltspinning in 250 Torr He, the optimally and overquenched alloys still show microstructural variations through the thickness of the ribbon. At higher wheel speeds, the heat extraction is sufficient to cool the liquid below the effective glass transition temperature Tg in only the first B20 mm of the ribbon. Since the intermetallic compound Nd2Fe14B has poor thermal conductivity [25], the middle-to-upper portions of the ribbon towards the free side fail to cool below the effective Tg resulting in successively increasing growth of pre-existing nuclei through the remaining portion of the ribbon (Fig. 5). Further increases in the quenching wheel speed do not appreciably increase the low coercivity glassy fraction since increasing the wheel speed above 30 m/s did not decrease the ribbon thickness [15] and thus do not change the total heat flow. 4.3. General peritectic rapid solidification model In those rapid solidification processes that do not achieve sufficient undercooling to form a glass, the phase selection process is complex. The experimental data presented in the previous section supports a general solidification model pertinent to processing of peritectic systems, such as the Nd2Fe14B alloy system. This model, although largely qualitative, is unique in that it provides a consistent explanation of the wide variety of observed microstructures and allows predictive and corrective control of processing conditions. This model’s most significant feature is that it explains the puzzling observation of c-axis growth of Nd2Fe14B in melt-quenched ribbons, in
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apparent contradiction to the single crystal studies that show the preferred growth is in the basal (a–baxis) plane. While it is possible that under nonequilibrium conditions the growth direction would change, the explanation presented in the peritectic solidification model described below is much simpler. The starting composition and the heat flow (which controls the solidification front velocity) are both important parameters in the determination of the final microstructure of the melt-spun ribbons. Under the conditions discussed in this paper, the primary solidifying phases are Nd2Fe14B [26] and Fe (Fig. 3). In the optimally and over-quenched ribbons the instantaneous quench at the wheel surface appears to be sufficient to suppress heterogeneous nucleation and growth. However, in the under-quenched ribbons the large variability in microstructure supports the hypothesis of nucleation-controlled grain growth at the liquid–solid (gas layer) interface. Conventional solidification theory for these high growth rate conditions would suggest the promotion of a cellular 2-14-1 morphology and possibly the existence of a small amount of fine cellular Fe which grows into the melt in the direction of the thermal gradient [27]. However, this result is not observed. As the growth front advances, the heat of formation of the 2-14-1 phase causes the temperature of the crystallized material to rise. The heat transport out of the ribbon lags behind the heat produced by recalesence and the temperature of the crystallized material rises because the thermal coupling between the quenching wheel and the molten alloy is imperfect. If the heat produced by recalesence causes the local temperature to rise above the metastable extension of peritectic temperature (Tp ) [23], the 2-14-1 solid phase must decompose to Fe+liquid [28]. If the temperature rise in the solidified alloy behind the growth front reaches the metastable peritectic temperature before the growth front stops at the free side of the ribbon, two conditions occur (Fig. 7). First, the primary solidifying phase is no longer 2-14-1 but rather is g-Fe. Second, the primary solidification structure, typically cellular 2-14-1, is no longer stable and will undergo peritectic decomposition. Since this
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Fig. 7. Graphical representation of the proposed solidification sequence in the under-quenched ribbons. The horizontal axis represents the time evolution of the under-quenched microstructure starting with columnar growth of 2-14-1 and Fe evolving to smaller but equiaxed microstructure due to dendritic breakup near the wheel surface. As the solidification front moves upward, the growth velocity slows while the temperature rises behind the front. The primary solidifying phase is Fe but switches quickly back to 2-14-1 due to changes in liquid composition and cooling.
decomposition is an endothermic event, it is most likely a transient phenomenon and is kinetically limited. This conclusion is evidenced by the very fine grains surrounding the decomposing cells shown in Fig. 3(a) which coarsen rapidly from the wheel side into the interior of the ribbon where the resident time above Tp is longer. This partial melting of the primary solidification microstructure results in ‘‘dendritic breakup’’ when the solidified cells are no longer stable in the liquid due to surface tension effects and are free to rotate, eliminating the prior microstructure [29]. The dendritic breakup model has been proposed to account for grain refinement in other alloy systems [30]. In the phase region where T > Tp; Fe is the primary solidifying phase. For very low quench rate conditions, well-developed dendritic growth of Fe has been observed, but only in the upper parts of the ribbons, near the free side. TEM observations have confirmed an interphase epitaxy between the 2-14-1 phase and the Fe phase where the orientational geometry is [1 1 0]ð1% 1 0ÞaFe8
or ½1% 1 0(1 1 0)aFe8½1 1% 0 [0 0 1]ð3% 3 0Þ2-14-1 (0 0 6)2-14-1. This epitaxial relationship produces a tendency for 2-14-1 texturing in a direction orthogonal to its fast growth direction, which supports the hypotheses of recalesence above the metastable peritectic temperature under low quench rate conditions [10]. Under optimal quenching conditions, cells of g-Fe grow in the /1 1 1S direction parallel to the thermal gradient, which is normal to the ribbon surface. A second growth stage for the 2-14-1 phase occurs when the semi-molten alloy cools relatively slowly below the peritectic by radiant cooling of the free surface. Since nucleation and growth occurring on an epitaxial surface is energetically favored over homogeneous nucleation in the peritectic liquid, epitaxial growth of the 2-14-1 phase on the existing Fe nuclei dominates the microstructural development. The preferred growth direction of the 2-14-1 phase is parallel to the basal planes in the plane of the ribbon surface, not via columnar growth from the wheel surface as previously suggested [17–20] (Fig. 7).
5. Conclusions The solidification of melt-spun 2-14-1 is a complex, stochastic process dependent on initial composition, quench wheel speed and chamber pressure. The phase selection process is complex but can be understood in general terms within the framework of solidification theory. Underquenched alloys have a heterogeneous microstructure by virtue of the diverse time-temperature profiles throughout the cross-section of the ribbon. The coarse microstructure found in the low quench rate alloys is attributed to the occurrence of recalesence above the peritectic temperature combined with rapid epitaxial growth of the 2-14-1 phase on Fe. The very fine scale microstructure of the optimally and over-quenched alloys is attributed to an instantaneous quench that occurs below the effective glass transition temperature, suppressing heterogeneous nucleation and grain growth at the wheel surface. However, retarded heat flow through the already-solidified portion of the ribbon results in lower cooling rates for the upper
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portions of the ribbon, near the free surface, and allows critical 2-14-1 nuclei to grow larger with increasing distance from the wheel surface. This model explains well the magnetic structures and microstructures of the as-quenched ribbons subjected to a variety of processing conditions. To better quantify the solidification models proposed, research into understanding the heat flow characteristics is underway.
Acknowledgements We would like to acknowledge Les Reed for his assistance in sample preparation. Research was performed at BNL under the auspices of the US Department of Energy, Division of Materials Sciences, Office of Basic Energy Sciences under contract No. DE-AC02-98CH10886. Ames Laboratory is operated by Iowa State University for the Department of Energy under Contract No. W7405-ENG-82. We gratefully acknowledge support from the DOE National Energy Research Undergraduate Laboratory Fellowship Program (ERULF). Work on double roller was partially supported by the Argentina Cooperative Research Program funded by CONICET, and by NSF under grant INT-9872792.
References [1] Roskill Information Services, The Economics of Rare Earths and Yttrium, 11th Edition, Clapham Road, London SW9 0JA, UK. [2] E.F. Kneller, R. Hawig, IEEE Trans. Magn. 27 (1991) 3588. [3] R.W. McCallum, A.M. Kadin, G.B. Clemente, J.E. Keem, J. Appl. Phys. 61 (1987) 3577; A.M. Kadin, R.W. McCallum, G.B. Clemente, J.E. Keem, in: M. Tenhover (Ed.), Science and Technology of Rapidly Quenched Alloys, MRS Symposium Proceedings, Vol. 80, 1987, pp. 385. [4] F. Pinkerton, J. Appl. Phys. 63 (11) (1988) 5427. [5] L.H. Lewis, J. Gao, D.C. Jiles, D.O. Welch, J. Appl. Phys. 79 (8) (1996) 6470.
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[6] F.K. Eckart, R. Hawig, IEEE Trans. Magn. 27 (4) (1991) 3588. [7] R. Pond, R. Maddin, Trans. Met. Soc. AIEM 245 (1969) 2475. [8] M.C. Narasimham, US Patent No. 4212343, 1980. [9] M.A. Otooni, in: Elements of Rapid Solidification, Springer Series in Materials Science, Vol. 29, Springer, Berlin, 1998. [10] M.J. Kramer, C.P. Li, K.W. Dennis, R.W. McCallum, in: S.P. Marsh, J.A. Dantzig, W. Hofmeister, R. Trivedi, M.G. Chu, E.J. Lavernia, J.H. Chun (Eds.), Rapid Solidification, Models and Experiments, Vol. 213, 1998. [11] D. Cochet-Muchy, S. Paidassi, Int. J. Rapid Solidif. 5 (1) (1989) 11. [12] E. Burzo, Rep. Prog. Phys. 61 (1998) 1099. [13] M.J. Kramer, Y. Tang, K.W. Dennis, R.W. McCallum, Material Research Society Symposium Proceedings, Advanced Hard and Soft Magnetic Materials, Vol. 577, 1999, pp. 57. [14] M.B. Stearns, Y. Cheng, J. Appl. Phys. 75 (10) (1994) 6894. [15] Y.L. Tang, M.J. Kramer, K.W. Dennis, R.W. McCallum, L.H. Lewis, in preparation. [16] M.J. Kramer, K.W. Dennis, R.W. McCallum, L. Fabietti, M. Villafuerte, in preparation. [17] D. Dadon, Y. Gefen, M.P. Dariel, IEEE Trans. 23 (1987) 3605. [18] R. Coehoorn, J. Duchateau, Mat. Sci. Eng. 99 (1988) 131. [19] C.R. Paik, M. Okada, M. Homma, IEEE Trans. 26 (1990) 1730. [20] T. Harada, T. Ando, R.C. O’Handley, N.J. Grant, J. Appl. Phys. 68 (1990) 4728. [21] P.C. Canfield, personal communication. [22] R.K. Mishra, J. Magn. Magn. Mater. 54–57 (1986) 450. [23] R. Hermann, W. Loser, J. Appl. Phys. 83 (1998) 6399. [24] C.L. Harland, H.A. Davies, B.E. Ward, F. Leccabue, in: L. Schultz, K.-H. Muller (Eds.), Proceedings of the 15th International Workshop on Rare Earth Magnets and Their Applications, Dresden, Germany, 1998, pp. 263–269. [25] R. Hanitsch, in: J.M.D. Coey (Ed.), Rare-Earth Iron Permanent Magnets, Oxford Scientific Publications, Oxford, 1996, p. 452. [26] T. Umeda, T. Okane, W. Kurtz, Acta Mater. 4 (1996) 4209. [27] W. Kurz, B. Giovanola, R. Trivedi, Acta Mater. 34 (1986) 823. [28] D.S. Tsai, T.S. Chin, S.E. Hsu, IEEE Trans. Magn. 23 (1987) 3609; G. Schneider, et al., Z. Metallkunde 77 (1986) 755. [29] M. Schwarz, A. Karma, K. Eckler, D.M. Herlach, Phys. Rev. Lett. 73 (1994) 1380. [30] A. Munitz, G.J. Abbaschian, in: E.W. Collings, C.C. Kocj (Eds.), Undercooled Alloy Phases AIME, 1986, pp. 23.
Topical Review
Solidification, microstructural refinement and magnetism in Nd2Fe14B M.J. Kramera,*, L.H. Lewisb, L.M. Fabiettic, Y. Tanga,1, W. Millerb,2, K.W. Dennisa, R.W. McCalluma a
USDOE and Department of Materials Science and Engineering, Ames Laboratory, Iowa State University, 37 Wilhelm Hall, Ames, IA 50011, USA b Materials Science Department, Brookhaven National Laboratory, Upton, NY 11973-5000, USA c Facultad de Matematica Astronomia y Fisica, Ciudad Universitaria, Cordoba 5000, Argentina Received 25 May 2001; received in revised form 24 October 2001
Abstract Nanocrystalline Nd2Fe14B (2-14-1) exhibits very favorable properties conferred by rapid solidification processes, including: simplified processing, good corrosion resistance and high magnetic hardness provided by a resultant grain size that is on the order of the single-domain size. While the ideal microstructural state of melt-spun 2-14-1 is a homogeneous dispersion of nanocrystallites embedded within a vitrified matrix, in practice it is very difficult to control the microstructural development to a uniform nanocrystalline state. This as-yet unmet challenge underlies the desire to understand, and ultimately to control, the solidification path taken by Nd–Fe–B-based melts during rapid solidification processing in order to obtain a homogeneous microstructure. Processing factors that influence the homogeneity and hence magnetic performance of the rapidly quenched microstructure of Nd2Fe14B are evaluated and categorized. These results emphasize both the complexities of rapid solidification processing as well as the potential for control of a highly chaotic processing technique. A qualitative solidification model that incorporates recalesence is presented. This model details the evolution of microstructures related to the thermal gradient as well as the phase selection process undertaken during solidification of rapidly solidified Nd–Fe–B alloys. The model is novel in that it explains the observed microstructures, from nanophased equiaxed grains transitioning to large elongated grains, and provides both diagnostic and predictive information concerning rapid solidification processing methods. r 2002 Elsevier Science B.V. All rights reserved. Keywords: Solidification; Peritectic; Recalesence; Coercivity
1. Introduction *Corresponding author. Tel.: +1-515-294-0276; fax: +1515-294-4291. E-mail address: [email protected] (M.J. Kramer). 1 Now at Argonne National Laboratory, 9700 South Cass Ave. Argonne, IL 60439, USA. 2 Now at University of Pennsylvania School of Engineering & Applied Science, 113 Towne Building 220 So. 33rd Street, Philadelphia, PA 19104-6391, USA.
Rare-earth-based permanent magnets, with coercivities (H ci ) generally exceeding 5000 Oe, are found in a wide variety of applications with demand for rare-earth permanent magnets growing by 10–15% per annum [1]. Today, the permanent magnet of choice for most advanced
0304-8853/02/$ - see front matter r 2002 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 0 1 ) 0 0 9 5 5 - 6
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applications is based on the intermetallic compound Nd2Fe14B (2-14-1). It is anticipated that rapid solidification processing will assume greater importance in the near future as it is the main technique used to produce bulk magnets with a nanoscale microstructure. The extremely small scale of microstructures produced by rapid solidification techniques approaches that of the domain wall widths in these materials [2] and thus it is possible for exchange interactions, in addition to the ever-present magnetostatic interactions, to magnetically couple majority and minority phases. Such effects can create magnetic mesostructures that consist of many exchange-coupled grains, producing ‘‘interaction domains’’ [3–5] that are much larger than the individual crystallographically defined grains. In some instances the intergranular exchange coupling is exploited to create two-phase ‘‘exchange–spring’’ magnets [6] that consist of a magnetically soft phase intimately mixed with a hard phase to form a magnetic composite that possesses the best properties of each constituent phase. Exchange–spring magnets are candidates for the next generation of permanent magnets whose properties can be optimized through better control of processing. The goal is to attain a better fundamental understanding of the factors that influence the microstructure, and hence magnetic properties, of rapidly solidified advanced magnetic materials. Melt–spinning is a widely used production method to produce rapidly solidified materials with a metastable microstructure [7,8]. Unlike traditional casting techniques which result in significant amounts of macrosegregation, the high cooling rates obtained in melt-spinning are typically believed to eliminate macrosegregation during solidification. However, in the Nd–Fe–B system significant microsegregation still occurs in rapidly solidified ribbons of Nd2Fe14B [9–11]. The grain size from the wheel side (chill surface) to the free side can vary by 2–3 orders of magnitude in a region less than 75 mm through the ribbons thickness [10]. The existence of processing-induced microstructural inhomogeneities in Nd2Fe14B, and the associated adverse effect upon the magnetic properties, have been known for many years [12]. However, few attempts have been made to system-
145
atically identify, categorize, model and compensate for these defects in this family of materials. It is empirically recognized that microsegregation can be minimized by undercooling the melt below the glass transition temperature Tg where conventional nucleation and growth processes, which lead to segregation, can be eliminated. However, heat flow limitations and stochastic effects frustrate the complete transformation of the melt into a uniform glass in ways that are not completely understood. The magnetic properties of rapidly solidified materials are especially sensitive to microstructure where control to the nanometer length scale is necessary. While the ideal microstructural state of melt-spun Nd–Fe–B is a homogeneous dispersion of nanocrystallites embedded within a vitrified matrix, in practice it is very difficult to control the structural development to the nanocrystalline state. This challenge underlies the desire to understand, and ultimately to control, the solidification path taken by Nd–Fe– B-based melts during rapid solidification processing. Recent results, outlined in this paper, obtained from a wide variety of experimental techniques underscore the dependence of the phase selection and microstructure of the final meltquenched Nd–Fe–B-based product on the condition of the as-quenched state. A solidification model is presented which qualitatively explains the wide range of microstructures observed and the phase selection process undertaken during solidification in rapidly quenched Nd–Fe–B alloys. The model explains the observed microstructures, from nanophased equiaxed grains transitioning to large elongated grains. These results emphasize both the complexities of the solidification in this system as well as the potential for controlling the highly chaotic melt-spinning processing technique.
2. Approach 2.1. Sample preparation Ingots of Nd–Fe–B with initial compositions of stoichiometric Nd2Fe14B were prepared by arcmelting in an Ar atmosphere. Charges of approximately 10 g were loaded in quartz tubes,
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inductively melted to about 13751C and ejected through an orifice of 0.8 mm in diameter with a 150 Torr Ar or He over-pressure onto a single rotating Cu wheel or a double roller system with stainless steel quench wheels. Melt-spinning was performed at wheel speeds from 5–40 m/s, providing resultant microstructures which spanned the range of crystallinity from coarsely crystalline at low speeds (under-quenched) to nanophased at intermediate wheel speeds (optimally quenched) and amorphous at high wheel speeds (overquenched). The chamber pressure and gas were 750 Torr of Ar or He, 250 Torr He and active vacuum [13].
2.2. Characterization 2.2.1. Thermal analyses Thermal analyses were performed on asquenched samples to assess the microstructural stability and to evaluate the effect of quenching atmosphere on the melt-spun microstructure. Differential thermal analysis (DTA) was carried out with a Perkin Elmer System 7 in flowing ultrahigh purity Ar using a flow rate of 50 cc/min and a heating rate of 101C/min. Differential scanning calorimetry (DSC) was done on a Perkin Elmer Pyris DSC at a scanning rate of 401C/min. Samples for thermal analysis were ground from melt-spun ribbon to o50 mm powder in a N2 glove box.
2.2.2. X-ray diffraction (XRD) The evolution of the ribbon microstructure was monitored across the thickness from the bottom of the ribbon that was in contact with the quenching wheel surface (‘‘wheel side’’) to the top of the ribbon in contact with the quenching atmosphere (‘‘free side’’). To determine the degree of variability in microstructure from wheel side to free side, X-ray diffraction was performed on 10–20 ribbon fragments placed on a glass slide using double-sided tape with the ribbons’ ‘‘wheel side surface’’ facing up or down. Standard 2y scans were then performed from 101 to 1001 with step size of 0.051 using Cu Ka radiation.
2.2.3. Microscopies An in-depth study of the quenched ribbon microstructure was carried out with electron and atomic force microscopies. Transmission electron microscopy (TEM) was performed on crosssection specimens of the ribbons. The samples were prepared by inserting several ribbons into 3 mm quartz tubes which were then filled with epoxy, cured then sliced into 3 mm disks. The disks were mechanically polished to about 20 mm thick then Ar ion milled at 5 kV with a LN2 cooled stage to perforation. Planar TEM samples were prepared by preferential polishing and ion milling from one side. TEM evaluations were conducted on Philips CM30 transmission electron microscope operated at 300 keV. Atomic/magnetic force microscopy (AFM/MFM) was used to examine the cross-sectional topographic and magnetic structure of the ribbons. AFM/MFM images were obtained with a Digital Instruments Nanoscope III AFM operated in tapping mode using commercial standard magnetic tips (MESP, Digital Instruments Inc.) which provided the best magnetic contrast for these samples. Each ribbon was examined with AFM/MFM in more than 10 regions to ensure a representative image of the ribbons. Representative examples of the crosssectional microstructure and associated magnetic domain structures are provided in Fig. 1. In Fig. 1(a) magnetic domains of opposite polarization are delineated by dark domain walls with a domain contrast that varies from grain to grain. In Fig. 1(b) the magnetic domains appear in light and dark contrast, and the grain contrast is more difficult to discern. The difference in the domain and grain contrast between Fig. 1(a) and (b) is due to the type of MFM tip utilized during the experiments. 2.2.4. Magnetic characterization Extrinsic technical magnetic properties such as remanence Br and coercivity Hci are highly dependent upon microstructural details. Inhomogeneities in the microstructure of rapidly solidified Nd2Fe14B ribbons give rise to a variety of magnetic phases with diverse technical magnetic properties. This is despite the fact that the overall chemical composition and sometimes crystal
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147
tion of the number and type of magnetic phases in the sample provides diagnostic information concerning the homogeneity of the sample. To this end, hysteresis loop measurements were performed using a Quantum Design SQUID magnetometer operated at 300 K. The room-temperature demagnetization data were fitted to an empirical model using the general curve fit option in KaleidaGraph software that delineates the ferromagnetic components present in the data. The utilized curve fit is empirical, not derived from first principals, and is not at all intended provide insight into crystallographic alignment, anisotropy or magnetic coupling, which are extremely challenging parameters to model, especially in a multi-phase alloy. MðHÞ ¼
X 2M S
F
n
Fig. 1. MFM images of a polished cross-sections for (a). 2-14-1 melt spun at 10 m/s in 750 Torr He and (b). 15 m/s in 250 Torr He. The lower wheel speed ribbons are thicker and show only a limited region of small magnetic domains. Both images are the same magnification.
structure of these phases may be identical. For instance, depending upon the processing details it is possible to find in a quenched ribbon glassy or amorphous Nd–Fe–B as well as Nd2Fe14B and aFe in their nanocrystalline, single-domain-type and coarse micron-scale forms. Although each of these magnetic phases provides a unique coercivity signature, the saturation magnetization value is not a diagnostic parameter because of the uncertainty in the volume fractions of the different magnetic phases. In general, magnetic phases with low coercivity and low remanence are attributed to amorphous Nd–Fe–B or large grains of Nd2Fe14B. High-coercivity magnetic phases likely consist of single-domain-type nanocrystalline Nd2Fe14B, while a low-coercivity yet high remanence signature may be attributed to micron-scale Nd2Fe14B or to a-Fe. Despite the uncertainty in the actual identity of the magnetic phase, broad quantifica-
p
tan1
ðH7Hc Þ pS tan : Hc 2
ð1Þ
The above equation represents a hysteresis curve of n ferromagnetic magnetic phases of saturation magnetization MFS ; with the offset along the field H axis identified by the coercivity Hc . S is the squareness factor, which is the ratio of the remanent magnetization BR ðH ¼ 0Þ to the saturation magnetization MFS [14]. In this manner, the number of magnetic phases in the samples are assessed and presented in Table 1. Weight fractions of different magnetic phases are delineated by their coercivity value and are listed in Table 1 as p1; p2 and p3 in order of decreasing phase coercivity. The percentage of a given magnetic phase was determined by its moment fraction. This representation of the hysteresis curve involves no assumptions concerning crystallographic texture or interphase exchange coupling present in the samples. Fig. 2 is a representative experimental demagnetization curve (data markers) and the constituent ferromagnetic phases as determined by this curve fitting. It is noted that there are three ferromagnetic phases, of different remanence and coercivity, present in this sample; this determination correlates well with the microstructural data presented in Section 3.2. The results obtained from the room-temperature magnetic characterization are presented in Table 1.
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Table 1 Results of the decomposition of room-temperature hysteresis loops into constituent magnetic phases p1, p2, p3, given in order of magnetic hardness. MS is the saturation magnetization and HC is the coercivity. The percentage of a given phase is determined by its moment fraction Quenching atmosphere
Wheel speed (m/s)
MS (p1) (emu/g)
% p1
1 atm Ar
5 10 15 20 30
57 68 72 38 19
56 67 69 39 15
978 2842 5011 9215 6790
45 31 33 51 53
5 10 15 20 30
47 75 49 57 20
42 75 46 47 16
1755 1328 7036 8453 10050
65 25 35 26 106
5 10 15 20 30
58 23 12 47.5 40
37 23 11 39 34
525 14000 8944 9290 7466
50 50 55 27 15
1 atm He
1/3 atm He
HC (p1) (Oe)
MS (p2) (emu/g)
% p2
44 33 31 52 42
HC (p2) (Oe) 271 505 730 852 80
68 25 30 21.5 84
395 354 1516 1462 202
32 51 51 22 13
215 5068 3095 4332 114
MS (p3) (emu/g)
% p3
HC (p3) (Oe)
19 54
19 43
12 10
23.5 38
22 31
627 20
25.5 41 47 62
26 38 39 53
815 632 10 10
3. Results
Fig. 2. Decomposed hysteresis loop of stoichiometric Nd2Fe14B melt-spun at 20 m/s in 1 atm Ar. The constituent magnetically hard and magnetically soft phases are shown and are identified in the figure legend.
The compilation and analysis of the experimental data obtained on the microstructural, phase constitution and magnetic properties of the rapidly solidified Nd–Fe–B-based ribbons demonstrate the complex interaction of microstructures and properties. They reveal the importance of the quenched-in nanoscale-structure of the melt-spun ribbons in the determination of the degree of homogeneity, and hence the magnetic performance, of the ribbons. Representative microstructures are illustrated in Fig. 3. As expected, exothermic data derived from thermal analyses indicate that a larger glass fraction is present in ribbons quenched at higher wheel speeds (Fig. 4). This conclusion is broadly supported by the magnetic and microstructural investigations. However, significant differences in the quenched microstructures are produced by small changes in the processing variables other than the wheel speed.
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Fig. 3. TEM micrographs taken at various distances from the ribbon wheel side on a sample that was obliquely thinned (i.e., polished surface B201 off the surface normal). (a) Regions near the wheel side showed nearly circular coarse-grained 2-14-1 undergoing what appears to be dynamic recrystallization, with the nanophased matrix increasing in grain size with distance from the approximately cellular 2-14-1 grains. (b) Slightly larger but more uniform 2-14-1 grains and occasional smaller Fe grains. (c) Coarser 2-14-1 with nearly continuous Fe grain-boundary phase. (d) Large-grained matrix of 2-14-1 with aligned Fe cells.
Fig. 4. Results of DSC measurements for stoichiometric 2-14-1 melt spun from 5 to 40 m/s under varying chamber pressures. While the sample quenched in an atmosphere of 250 Torr He possessed the largest glass fraction based on the area under the curve of the DSC exotherm during devitrification, the sample quenched under active vacuum had the lowest glass fraction.
The overall microstructure of the melt-spun ribbons can be grouped into three broad categories. The first category is the under-quenched state, where there exists a substantial fraction of grains with diameters greater than the optimal (B100 nm) grain size necessary for good hard magnetic properties. The second category is the optimally quenched state where the majority of the grains have a near-optimal grain size (B20 nm) to support a high coercivity, and the third category is the over-quenched state where a sizable volume fraction of the material is magnetically soft because the matrix is predominantly amorphous. The appropriate circumferential quenching wheel speeds where these transition points are observed depend upon the alloy composition, the thermal conductivity of the quench wheel, chamber gas
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and pressure and the temperature of the melt prior to rapid solidification (i.e., superheat) of the alloy [15]. 3.1. Under-quenched microstructure The microstructure of an under-quenched 2-141 ribbon is highly variable both through its thickness and along its length. Optical microscopy, SEM and TEM studies on these ribbons confirmed that the microstructure varies from nanophased to coarsely crystalline (grain size dgrain > 10 mm) [10]. Typically, underquenched microstructures are obtained in the Nd–Fe–B system with low wheel speeds (5–10 m/s) and display two distinct magnetic phases with coercivities in the intermediate (Hci o5000 Oe) and low (Hci o700 Oe) ranges. The cross-sectional microstructure has four distinct layers, shown in more detail by TEM (Fig. 3) and partially observable in the MFM images (Fig. 1). The wheel side layer (in Fig. 3(a)) is thin (B0.3 mm) and consists of fine equiaxed 2-14-1 grains and larger cells of 2-14-1, which have apparently undergone recrystallization. A thicker region follows, up to 20 mm thick in some samples, where the equiaxed 2-14-1 grains coarsen systematically from less than 100 nm in size to about 300 nm (Fig. 1(b)). In this region smaller a-Fe particles frequently exist at triple conjunctions of the 2-14-1 grains (in Fig. 3(b)), yet no grain boundary phase has been observed. Further away from the wheel side surface, the grains coarsen and achieve sizes up to 1 mm or more in diameter. As the grains coarsen, occasional a-Fe grains are seen to coalesce and give rise to a thin continuous layer of an a-Fe intergranular phase (Fig. 3(c)). This intermediate layer gives way to a microstructure composed of large equiaxed 2-14-1 grains (dgrain > 10 mm) (Fig. 2) with aligned a-Fe inclusions that exist close to the free side in single roller experiments and in the ribbon center in double roller experiments [16] (Fig. 3(d)). Texturing of the 2-14-1 with the c-axis parallel to the ribbon surface has been observed in both single and twin roller experiments [17–20]. While the texturing has been attributed to columnar growth, this explanation is at odds with preferred basal plane growth observed in single crystals [21].
It should be emphasized that the demarcation between these zones and the coercivity values of these layers are approximate and the thickness of these various layers is dependent upon processing conditions. SEM and MFM images over a large number of cross-sections for samples processed in varying chamber atmospheres and pressures demonstrate that in general the thickness of the coarse grained regions decreased with increasing wheel speed and decreasing the chamber pressure to 250 Torr He. 3.2. Optimally quenched microstructure An increase in the quenching wheel speed produces an increase in the volume fraction of the glassy phase as measured by the increase in the total heat evolved during crystallization in the DSC (Fig. 4) [13]. In addition, the remaining crystalline microstructure becomes finer with increasing wheel speed as indicated by TEM investigation and Bragg peak broadening in the XRD data. For wheel speeds in the range of 15– 20 m/s and a chamber atmosphere of 750 Torr Ar, a near optimally quenched microstructure with grain sizes in the range of 80–100 nm dominates. Typical coercivity values of the hard magnetic phase in optimally processed ribbons are in the range 8–14 kOe, although these materials still exhibit two or more magnetic phases of lower coercivity. Reducing the chamber pressure to 250 Torr (0.333 atm) He allowed the wheel speed necessary to achieve an optimal microstructure to be reduced by B5 m/s, Table 1 (Fig. 4). Assuming the existence of a positive thermal gradient across the ribbon thickness, the nucleation rate (either heterogeneous or homogeneous) is dependent on the degree of undercooling and should vary through the ribbon thickness, resulting in a varying grain size through the ribbon thickness. This microstructure is demonstrated in Fig. 5 where the wheel side of the ribbon is amorphous as determined by electron diffraction in the TEM. Yet the cross-sectioned sample examined by TEM reveals a gradual increase in the size and density of the nanocrystalline particles towards the free surface of the ribbon, along the main thermal gradient. Deconstruction of the
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Fig. 5. TEM cross-section of 22.5 m/s stoichiometric 2-14-1 melt spun under 750 Torr He which shows an increase in proportion of nanophased 2-14-1 and a reduction in the glass fraction from (a) wheel side to (b) free side. The insets are selected area diffraction patterns taken in the region of the bright field images.
quenched ribbons, a-Fe particles are still frequently observed at triple points of 2-14-1 grains. The presence of such a highly dispersed a-Fe phase is inconsistent with homogeneous nucleation of the stoichiometric composition and suggests competing energetics between nucleating Fe and 2-14-1 phases. 3.3. Over quenched microstructure
Fig. 6. TEM micrograph of an over quenched ribbon showing small nuclei are 2-14-1 grains, as confirmed by convergent beam diffraction. (inset upper left).
hysteresis loops measured on the same samples demonstrates that the optimal hard magnetic fraction varies systematically with wheel speed and chamber pressure as well as supports the observation that mixed microstructures occur at all wheel speeds (Table 1). The maximum in the hard magnetic phase fraction agrees well with the microstructural observations of predominately nanophased 2-14-1. As found in the under-
Further increases in the circumferential wheel speed produces microstructures with much finer crystalline 2-14-1 grains (on the order of 20 nm or less) and larger phase fractions of Nd–Fe–B glass of very low coercivity Hci o50 Oe (Table 1). For ribbons melt spun at 30 m/s, about 70 wt% (which is close to 70 vol%) of the microstructure is amorphous. Other studies [22] have demonstrated that amorphization is nearly complete for wheel speeds in excess of 30 m/s. A reduced chamber pressure reduces the quench rate necessary to achieve a substantially glassy matrix [13] as identified by XRD, TEM selected area diffraction patterns (SADP) and thermal analysis techniques. Meanwhile, within nominally glassy regions, extremely small nuclei, with diameters dp5 nm have been observed using conical bright field (BF) and dark field (DF) TEM techniques performed on the same region. Utilization of the convergent beam electron diffraction (CBED) TEM technique revealed the nuclei to consist of the Nd2Fe14B phase (Fig. 6). It should be noted that for our
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melt-spinner at high wheel speeds, the contact between the melt pool and the wheel is destabilized. This was most likely due to wheel vibrations, resulting in regions of entrained gas where the quenching rate is lower than expected. While these regions are occasionally observed in the TEM images, the DTA and XRD results show that the bulk of the alloy melt-spun at wheel speeds above 20 m/s in 250 Torr He is amorphous. The amorphous nature of the alloy is also confirmed by the high volume fraction of the low coercivity component for samples melt-spun above 20 m/s (Table 1).
4. Discussion There is a considerable body of literature on rapidly solidified R2Fe14B (R=Pr, Nd, Tb, Dy) based alloys. Most of the research has concentrated on altering the alloy composition to enhance certain magnetic properties, i.e., the Curie temperature Tc ; ðBHÞmax ; Hci ; etc. [12]. Only a small portion of the research has concentrated on understanding the effects of solidification on phase selection [10,23]. Yet, the magnetic properties are intimately connected to the as-quenched microstructures since local phase assemblages, as well as the overall chemistry, influence the crystallization pathway. The magnetic properties, in turn, are linked to the size and distribution of the primary 214-1 matrix and the secondary phase distribution. For instance, Table 1 shows that coercivity peaks in the region of moderate quenching wheel speeds (10–15 m/s) and low chamber pressures of helium. In order to optimize the technical magnetic properties, better understanding and control of the solidification process must be obtained. The effects of major processing parameters on the quenched microstructure are discussed below. 4.1. Effect of atmosphere The species and pressure of the quenching gas also have a significant effect on the homogeneity, microstructure and phase constitution of the quenched ribbon. In general it is found that a low pressure of quenching gas possessing a
relatively high specific heat delivers the best and most uniform quench [13]. This conclusion is supported by detailed analysis of the longitudinal cross-sections of the ribbons processed in 750 Torr chamber pressure which reveals rather startling microstructural inhomogenities for all wheel speeds. Profilometer measurements of the wheel side topography indicate the increased presence of elongated regions of deep surface depressions up to 18 mm in depth throughout the wheel surface [13,24]. These regions were elongated in a direction parallel to the direction of the tangential motion of the quench wheel. SEM observations revealed that the regions on the ribbon free side directly above the wheel-side depressions show evidence of coarsely crystalline material for nearly all wheel speeds. Since it is suspected that these poorly quenched regions arose from poor contact between the melt pool and the wheel surface due to trapped gas pockets [13], DSC (Fig. 4), XRD and profilometer measurements were repeated on ribbons melt-spun from 5 to 30 m/s in the atmospheres of 250 Torr of He and an active vacuum. The wheel surface appeared both visually and by profilometer measurements to be much smoother, with depressions of less than 5 mm in depth. XRD demonstrates that the formation of nanograined 214-1 was enhanced while DSC shows that the relative proportion of glass was increased by lowering the chamber pressure to 250 Torr He, but not to an active vacuum. In fact, vacuumprocessed ribbons had the largest fraction of crystalline material as compared to Ar- or Heprocessed material at any wheel speed regardless of the chamber pressure. The explanation for poor quenching in vacuum is quite simple: as the ribbons leave the wheel glowing (in excess of 7001C), the final cooling of the ribbon occurs by conduction/convection into the surrounding environment during free flight, a mechanism which is minimally operative in vacuum but is still quite effective in 250 Torr He. 4.2. Heat flow considerations Achievement of through-thickness homogeneity in the ribbon microstructure appears to be limited by unidirectional heat flow through the solidified
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portion of the ribbon to the Cu quench wheel. Local perturbations to the overall heat flow during solidification have a dramatic effect on the phase content of the quenched product, as discussed above for the under-quenched alloys. Reduction in the chamber pressure decreases the occurrence of the trapped gas between the liquid pool and quench wheel, a phenomenon which locally impedes heat flow to the wheel; the use of He gas as a quenching medium improves the quench by virtue of its higher heat capacity. However, even with the resultant improved uniformity of the quench along the ribbon surface obtained by meltspinning in 250 Torr He, the optimally and overquenched alloys still show microstructural variations through the thickness of the ribbon. At higher wheel speeds, the heat extraction is sufficient to cool the liquid below the effective glass transition temperature Tg in only the first B20 mm of the ribbon. Since the intermetallic compound Nd2Fe14B has poor thermal conductivity [25], the middle-to-upper portions of the ribbon towards the free side fail to cool below the effective Tg resulting in successively increasing growth of pre-existing nuclei through the remaining portion of the ribbon (Fig. 5). Further increases in the quenching wheel speed do not appreciably increase the low coercivity glassy fraction since increasing the wheel speed above 30 m/s did not decrease the ribbon thickness [15] and thus do not change the total heat flow. 4.3. General peritectic rapid solidification model In those rapid solidification processes that do not achieve sufficient undercooling to form a glass, the phase selection process is complex. The experimental data presented in the previous section supports a general solidification model pertinent to processing of peritectic systems, such as the Nd2Fe14B alloy system. This model, although largely qualitative, is unique in that it provides a consistent explanation of the wide variety of observed microstructures and allows predictive and corrective control of processing conditions. This model’s most significant feature is that it explains the puzzling observation of c-axis growth of Nd2Fe14B in melt-quenched ribbons, in
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apparent contradiction to the single crystal studies that show the preferred growth is in the basal (a–baxis) plane. While it is possible that under nonequilibrium conditions the growth direction would change, the explanation presented in the peritectic solidification model described below is much simpler. The starting composition and the heat flow (which controls the solidification front velocity) are both important parameters in the determination of the final microstructure of the melt-spun ribbons. Under the conditions discussed in this paper, the primary solidifying phases are Nd2Fe14B [26] and Fe (Fig. 3). In the optimally and over-quenched ribbons the instantaneous quench at the wheel surface appears to be sufficient to suppress heterogeneous nucleation and growth. However, in the under-quenched ribbons the large variability in microstructure supports the hypothesis of nucleation-controlled grain growth at the liquid–solid (gas layer) interface. Conventional solidification theory for these high growth rate conditions would suggest the promotion of a cellular 2-14-1 morphology and possibly the existence of a small amount of fine cellular Fe which grows into the melt in the direction of the thermal gradient [27]. However, this result is not observed. As the growth front advances, the heat of formation of the 2-14-1 phase causes the temperature of the crystallized material to rise. The heat transport out of the ribbon lags behind the heat produced by recalesence and the temperature of the crystallized material rises because the thermal coupling between the quenching wheel and the molten alloy is imperfect. If the heat produced by recalesence causes the local temperature to rise above the metastable extension of peritectic temperature (Tp ) [23], the 2-14-1 solid phase must decompose to Fe+liquid [28]. If the temperature rise in the solidified alloy behind the growth front reaches the metastable peritectic temperature before the growth front stops at the free side of the ribbon, two conditions occur (Fig. 7). First, the primary solidifying phase is no longer 2-14-1 but rather is g-Fe. Second, the primary solidification structure, typically cellular 2-14-1, is no longer stable and will undergo peritectic decomposition. Since this
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Fig. 7. Graphical representation of the proposed solidification sequence in the under-quenched ribbons. The horizontal axis represents the time evolution of the under-quenched microstructure starting with columnar growth of 2-14-1 and Fe evolving to smaller but equiaxed microstructure due to dendritic breakup near the wheel surface. As the solidification front moves upward, the growth velocity slows while the temperature rises behind the front. The primary solidifying phase is Fe but switches quickly back to 2-14-1 due to changes in liquid composition and cooling.
decomposition is an endothermic event, it is most likely a transient phenomenon and is kinetically limited. This conclusion is evidenced by the very fine grains surrounding the decomposing cells shown in Fig. 3(a) which coarsen rapidly from the wheel side into the interior of the ribbon where the resident time above Tp is longer. This partial melting of the primary solidification microstructure results in ‘‘dendritic breakup’’ when the solidified cells are no longer stable in the liquid due to surface tension effects and are free to rotate, eliminating the prior microstructure [29]. The dendritic breakup model has been proposed to account for grain refinement in other alloy systems [30]. In the phase region where T > Tp; Fe is the primary solidifying phase. For very low quench rate conditions, well-developed dendritic growth of Fe has been observed, but only in the upper parts of the ribbons, near the free side. TEM observations have confirmed an interphase epitaxy between the 2-14-1 phase and the Fe phase where the orientational geometry is [1 1 0]ð1% 1 0ÞaFe8
or ½1% 1 0(1 1 0)aFe8½1 1% 0 [0 0 1]ð3% 3 0Þ2-14-1 (0 0 6)2-14-1. This epitaxial relationship produces a tendency for 2-14-1 texturing in a direction orthogonal to its fast growth direction, which supports the hypotheses of recalesence above the metastable peritectic temperature under low quench rate conditions [10]. Under optimal quenching conditions, cells of g-Fe grow in the /1 1 1S direction parallel to the thermal gradient, which is normal to the ribbon surface. A second growth stage for the 2-14-1 phase occurs when the semi-molten alloy cools relatively slowly below the peritectic by radiant cooling of the free surface. Since nucleation and growth occurring on an epitaxial surface is energetically favored over homogeneous nucleation in the peritectic liquid, epitaxial growth of the 2-14-1 phase on the existing Fe nuclei dominates the microstructural development. The preferred growth direction of the 2-14-1 phase is parallel to the basal planes in the plane of the ribbon surface, not via columnar growth from the wheel surface as previously suggested [17–20] (Fig. 7).
5. Conclusions The solidification of melt-spun 2-14-1 is a complex, stochastic process dependent on initial composition, quench wheel speed and chamber pressure. The phase selection process is complex but can be understood in general terms within the framework of solidification theory. Underquenched alloys have a heterogeneous microstructure by virtue of the diverse time-temperature profiles throughout the cross-section of the ribbon. The coarse microstructure found in the low quench rate alloys is attributed to the occurrence of recalesence above the peritectic temperature combined with rapid epitaxial growth of the 2-14-1 phase on Fe. The very fine scale microstructure of the optimally and over-quenched alloys is attributed to an instantaneous quench that occurs below the effective glass transition temperature, suppressing heterogeneous nucleation and grain growth at the wheel surface. However, retarded heat flow through the already-solidified portion of the ribbon results in lower cooling rates for the upper
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portions of the ribbon, near the free surface, and allows critical 2-14-1 nuclei to grow larger with increasing distance from the wheel surface. This model explains well the magnetic structures and microstructures of the as-quenched ribbons subjected to a variety of processing conditions. To better quantify the solidification models proposed, research into understanding the heat flow characteristics is underway.
Acknowledgements We would like to acknowledge Les Reed for his assistance in sample preparation. Research was performed at BNL under the auspices of the US Department of Energy, Division of Materials Sciences, Office of Basic Energy Sciences under contract No. DE-AC02-98CH10886. Ames Laboratory is operated by Iowa State University for the Department of Energy under Contract No. W7405-ENG-82. We gratefully acknowledge support from the DOE National Energy Research Undergraduate Laboratory Fellowship Program (ERULF). Work on double roller was partially supported by the Argentina Cooperative Research Program funded by CONICET, and by NSF under grant INT-9872792.
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