Evolution of magnetic properties and microstructure of Hf 2 Co 11 B alloys

Transcription

1 Evolution of magnetic properties and microstructure of Hf 2 Co 11 B alloys Michael A. McGuire and Orlando Rios Oak Ridge National Laboratory, Oak Ridge, Tennessee USA Amorphous Hf 2 Co 11 B alloys produced by melt-spinning have been crystallized by annealing at 5 8 C, and the products have been investigated using magnetization measurements, x-ray diffraction, and scanning electron microscopy. The results reveal the evolution of the phase fractions, microstructure, and magnetic properties with both annealing temperature and time. Crystallization of the phase denoted HfCo 7, which is associated with the development of coercivity, occurs slowly at 5 C. Annealing at intermediate temperatures produces mixed phase samples containing some of the HfCo 7 phase with the highest values of remanent magnetization and coercivity. The equilibrium structure at 8 C contains HfCo 3 B 2, Hf 6 Co 23 and Co, and displays soft ferromagnetism. Maximum values for the remanent magnetization, intrinsic coercivity, and magnetic energy product among the samples are approximately 5.2 kg, 2. koe, and 3.1 MGOe, respectively, which indicates that the significantly higher values observed in crystalline, melt-spun Hf 2 Co 11 B ribbons are a consequence of the non-equilibrium solidification during the melt-spinning process. Application of high magnetic fields during annealing is observed to strongly affect the microstructural evolution, which may provide access to higher performance materials in Zr/Hf-Co hard ferromagnets. The crystal structure of HfCo 7 and the related Zr analogues is unknown, and without knowledge of atomic positions powder diffraction cannot distinguish among proposed unit cells and symmetries found in the literature. I. INTRODUCTION For a ferromagnet to be useful as a permanent magnet material, it must not only have a large magnetic moment, but also a Curie temperature well above the expected operating temperature, and high coercivity to resist demagnetizing fields. The strength of a permanent magnet is often quantified by the magnetic energy product denoted BH max, which is the maximum absolute value of the product of the magnetic induction B and the applied magnetic field H realized in the second quadrant of the magnetic hysteresis loop [1]. Current state of the art permanent magnets are based on Nd 2 Fe 14 B with BH max up to about 55 MGOe [2, 3], and SmCo 5, with BH max up to about 3 MGOe [4]. The next-best-performing commercial magnet is AlNiCo [5], with BH max up to about 1 MGOe [1]. Thus, there is interest not only in new permanent magnet materials which can compete with the highest performance rare-earth magnets, but also for materials which can outperform AlNiCo, and therefore displace rare-earth magnets in applications which require intermediate energy products (1 < BH < max 25), as well as those which can operate at higher temperatures than NdFeB-based magnets. In efforts to explore alternative materials for permanent magnet applications which do not contain rare earth elements [6], cobalt rich alloys with Zr and Hf have attracted some attention. These alloys have compositions ranging from MCo 5 to MCo 7 (M = Zr, Hf), Curie temperatures near 5 C [7 11] and are often made by meltspinning [9, 1, 12 18] and more recently by cluster beam deposition [11, 19, 2]. Energy products of 16 2 MGOe McGuireMA@ornl.gov have been reported for samples made by cluster deposition [2], while the highest reported values for melt-spun materials range from 5 8 MGOe [9, 12, 13, 21]. The Hf/Zr-Co magnet alloys typically contain multiple chemical phases, sometimes including multiple ferromagnetic phases. There is some evidence that the combination of hard and soft ferromagnetic materials plays an important role in producing the best permanent magnet performance [9, 2, 22]. In our previous study of melt-spun Hf 2 Co 11 B ribbons [9], we found that by altering the quenching rate during the melt-spinning process we could produce either amorphous, magnetically soft material or crystalline, magnetically hard material. The quenching rates were controlled by varying the speed of the spinning copper wheel onto which the molten alloy is ejected, and amounts essentially to an in-situ thermal treatment. Here we report the results of ex-situ thermal treatment of the amorphous precursor material, produced by melt spinning with a surface velocity of 24 m/s [9]. The purpose of the present study is threefold: (1) determine whether thermal processing of amorphous precursor material can produce hard magnetic properties similar to the melt-spun-crystalline material in a more controllable way, (2) examine the phase and microstructural evolution of the amorphous material associated with the evolution of the magnetic behavior, and (3) evaluate the ability of powder diffraction to distinguish between the proposed crystal structures for the hard magnetic phase in (Zr/Hf)Co 5 7 alloys. In keeping with the published phase diagram [23] and the majority of the literature cited above, we will refer to the hard ferromagnetic phase in the Hf-Co system as HfCo 7. We report results of magnetization measurements, powder x-ray diffraction, and microstructural analysis using scanning electron microscopy for materials pro-

2 2 duced by annealing amorphous ribbons of composition Hf 2 Co 11 B at temperatures ranging from 5 to 8 C and times ranging from.125 to 26 hours. Effects of high magnetic field applied during the annealing process are also demonstrated and discussed. The temperature range investigated spans the two exothermic phase transitions at 565 and 63 C in the amorphous material, which were identified in thermal analysis and magnetization measurements [9]. Our findings reveal how the magnetic anisotropy, chemical phase fractions, and microstructural morphology evolve. None of the thermal treatments produced magnetic coercivity as high as that observed in the optimally melt-spun materials, indicating the non-equilibrium conditions produce a unique microstructure which is most favorable for hard ferromagnetism. Before presenting and discussing the experimental results, we give here a brief review of what is known regarding the Hf-Co equilibrium phase diagram and the magnetic properties of the chemical phases encountered during the course of this study. The Hf Co phase diagram is found in Ref. 23. The intermediate phases on this equilibrium phase diagram that are found in the samples studied in this work are Hf 6 Co 23, and HfCo 7. The phase labeled HfCo 7 is the magnetically hard phase of interest to permanent magnet research. It is reported to be thermodynamically stable only between 15 and 1245 C. The crystal structure and precise stoichiometry of this phase is unknown. The Zr-Co phase diagram shows a related phase which is known as Zr 2 Co 11. Powder diffraction patterns from hard ferromagnetic samples of HfCo 7 and Zr 2 Co 11 are quite similar [8 11, 15 18, 21], suggesting these compounds are structurally related and perhaps isostructural with the same or similar stoichiometries. As noted above, we will use here HfCo 7 to denote the hard ferromagnetic phase in the Hf-Co system. This compound has a Curie temperature near 77 K and the saturation magnetization (J S = 4πM S ) of samples containing this as the main phase range from about 8 11 kg, corresponding to a theoretical maximum energy product (JS 2 /4) of MGOe [9, 1, 19]. Of the remaining compounds in the Hf-Co system, elemental cobalt and Hf 6 Co 23 were found to be present in some of the samples studied in the present work. Cobalt is ferromagnet with a Curie temperature of 1388 K, J S = 18 kg, and an anisotropy field near 1 koe [24]. Hf 6 Co 23 is also a ferromagnet, with a reported Curie temperature of 75 K, and is expected to have little magnetic anisotropy due to its cubic crystal structure [25]. No further investigation of the magnetic properties of this material were found in the literature. In this study, boron was found in the form of HfCo 3 B 2 and Co 23 B 6. No results of magnetization measurements were located in the literature for either of these materials. In the rare earth (R) containing RCo 3 B 2 phases, cobalt has been found to be paramagnetic at room temperature [26], and the same may be expected for HfCo 3 B 2 though this is not certain. Theoretical calculations predict ferromagnetic order in Co 23 B 6 [27]. II. EXPERIMENTAL DETAILS Amorphous ribbons of composition Hf 2 Co 11 B were prepared by melt-spinning as described in Ref. 9. Sections of this material were sealed in evacuated silica ampoules for the annealing studies. The sealed tubes were placed inside electrical box furnaces preheated to the desired annealing temperatures, left for a specified amount of time, and then quenched in water. Annealing times ranged from 1 to 26 hours (about 11 days) at 5 C, from.5 to ours at 6 and 7 C, and from.125 to ours at 8 C. In this work the samples are identified by the furnace temperature and the time spent in the furnace. It is understood that some time is required for the sample to come into thermal equilibrium with the furnace, and thus the durations at the specified temperatures are shorter than the specified times. Thermomagnetic processing was performed in a 9 Tesla, 5 inch diameter warm bore superconducting magnet with a 12 inch uniform field zone fitted with an RF induction heating insert. The materials were placed in a sealed ampoule within a 1 inch diameter 34 stainless steel susceptor located within the RF coil. Electromagnetic energy is supplied by a 9 kw power supply coupled with an applicator that is 6 inches in length. The thickness of the susceptor was selected to be greater that 1 times the skin depth, therefore the sample is shielded from the intense RF energy and heated by only by radiant energy supplied uniformly by the tubular heating element. Thermomagnetic processing was performed using actively controlled thermal profiles that approximately reproduce the transient conditions experienced by the samples annealed by insertion into a preheated tube furnace. The encapsulated material was loaded into the electromagnetic processing insert within the superconducting magnet at 9 Tesla while in persistent mode. The hot zone within the insert is heated from room temperature to 6 C in 3 minutes and held for ours. The sample was extracted from the insert at 6 C and immediately water quenched. The phase compositions of the annealed materials were characterized using x-ray diffraction (Cu K α radiation) from the free surfaces of the ribbons, the surface not in contact with the melt-spinner wheel. A PANalytical X Pert Pro MPD diffractometer was used for the diffraction measurements. The microstructure was examined using a JEOL 65 FEG-type scanning electron microscope (SEM) operating at 1 kv accelerating voltage in back scattered electron mode using a 14 mm working distance in order to observe slight compositional variations between phases in within the microstructure. The melt spun ribbons were mounted in cross-section (the wheel face normal to the polish plane) in room temperature curing epoxy resin. As discussed in Ref. 9 the microstructure of the melt spun ribbon can vary across

3 3 TABLE I. Magnetic properties determined from the magnetization curves shown in Fig. 1 for samples produced by annealing amorphous Hf 2Co 11B ribbons. T anneal time J 6T B r i H C B H C BH max ( C) h (kg) (kg) (koe) (koe) (MGOe) the sample with a distinct yet well-defined region adjacent to the wheel/ribbon interface. In the melt spun amorphous materials this interfacial region was on average 3 µm and is characterized by a less compositional variation in comparison to the bulk. Since we expect the magnetic properties of the bulk regions dominate the macroscopic physical properties, all SEM images are taken near the center of the ribbon and for consistency all samples are aligned with the wheel side of the ribbon cropped off to the left and approximately parallel to the vertical edges of the images shown below. The samples were polished using conventional metallographic sample preparation methods. The compositional contrast in SEM images is enhanced by chemical etching that preferentially attacks the cobalt-rich regions. Magnetic properties were measured at room temperature using a commercial DC magnetometer (Quantum Design, Physical Property Measurement System). Samples were measured with the field along the long axis of the ribbons, and sample masses were measured using a microbalance (Mettler MT5). The magnetization (J ) in units of kg was determined by multiplying the magnetization (M ) in emu/cm 3 by 4π. A density of 1.7 g/cm 3 was used in the calculations [9]. III. RESULTS AND DISCUSSION The magnetic properties of the annealed Hf 2 Co 11 B ribbons measured at room temperature are summarized in Fig. 1. The panels are labeled by the annealing temperature, and the data in each panel are labeled by the time spent in the furnace at that temperature. The data were measured upon decreasing the magnetic field (H ) from 6 koe (1 koe = 1 6 /(4π) A/m). For each annealing condition, the magnetization (J ) and magnetic energy product (BH, where B = J + H is the magnetic induction) are shown. The magnetization is a measure of the magnetic moment per unit volume and is given in units of kilogauss (J = 4πM with M in units of emu/cm 3 ). In SI units, 1 kg =.1 T. The first and second quadrant of this data is shown separately for clarity. The first quadrant data [Fig. 1(a,d,g,h)] gives an estimate of the saturation magnetization (J S ), for which J (H = 6 koe) is used here, and illustrates the approach to saturation. The second quadrant [Fig. 1(b,e,h,k)] contains the demagnetization curves. From these, the remanent magnetic induction [B r = J(H = )] and the intrinsic coercive field [ i H C = H(J = )] are determined. The energy product [Fig. 1(c,f,i,l)] is a measure of magnetic potential energy, and is a product of the magnetic induction of the sample (B) and the applied magnetic field (H). It is reported in units of megagauss- Oersted. To convert to SI units, note that 1 MGOe = 7.96 kj/m 3. In addition to the maximum energy product (BH max ), the coercive field for the magnetic induction [ B H C = H(B = )] can be identified from these plots. The as-spun, amorphous starting material is a very soft ferromagnet, which reaches saturation at applied fields on about 1 Oe and has no detectable coercivity ( i H C < 1 Oe) [9]. The data in Fig. 1 shows that annealing this material results in the development of hard ferromagnetic behaviors, including curvature extending to high magnetic fields in the approach to saturation, and significant coercivity noted in the demagnetization curves. Values of the relevant magnetic properties determined from this data are summarized in Table I. The magnetic properties evolve slowly when annealed at 5 C [Fig. 1(a,b,c)]. This is not unexpected, since the differential thermal analysis (DTA) showed the first crystallization peak occurs near 565 C [9]. The DTA curve is reproduced in Fig. 2a. For annealing times less than, no coercivity is detected. For longer annealing times, both B r and i H C gradually increase and positive energy products are observed, but for the longest annealing time studied, 26 h or 11 days, i H C was still limited to less than 1 koe, and BH max < 1 MGOe. Increasing the annealing temperature to 6 C significantly accelerates the magnetic property development, as shown in Fig. 1(d,e,f). This temperature lies between the two thermal anomalies at 565 and 63 C observed in DTA [Fig. 2a]. Intrinsic coercivities near 2 koe are observed for all of the samples. BH max reaches a maximum after annealing, and decreases as the time is further increased. This decrease in performance for longer annealing times is due primarily to a decrease in the magnetization, rather than a change in the coercivity. This decrease in remanent magnetization may be related to

4 4 J (kg) J (kg) BH (MGOe) (a) 6 h 28 h 26 h 4.5 h.5 h (b) 5 C (e) 6 C (h) 7 C (k) 8 C C (d) 6 C (g) 7 C (j) 8 C 5 C (f) 6 C (i) 7 C (l) 8 C.125 h.25 h.5 h FIG. 1. Magnetization data from Hf 2 Co 11 B ribbons measured at room temperature after annealing at the indicated temperatures and times. The field was directed along the length of the ribbon segments, so demagnetization effects are neglected. the evolution of chemical phases present or coarsening of the hard magnetic crystallites with time. Figure 1(g,h,i) shows that annealing at 7 C has a similar effect on the evolution of the properties as annealing at 6 C, although the saturation and remanent magnetization very less with annealing time. The main effect here is on the coercivity, which is maximized for t =.5. In addition, the demagnetization curves after annealing at 7 C [Fig. 1(h)] are more square than those obtained after annealing at 6 C [Fig. 1(e)]. This results in a larger energy product, exceeding 3 MGOe for annealing. For the shortest annealing time,.5 h, at 7 C, the coexistence of a soft magnetic component is evidenced by the slight inflection in J(H) near H =. The highest annealing temperature investigated here is 8 C. Figure 1(j,k,l) shows that the magnetic properties evolve very rapidly at this temperature. Significantly shortened annealing times were also investigated, since the times used at lower temperatures were observed to produce relatively soft behavior in the demagnetization curves. [Fig. 1(k)]. This results in a rapid decrease in ih C with annealing time, likely due to the presence of a large fraction of Hf 6 Co 23 and little HfCo 7 (see discussion of diffraction results below). The highest coercivity and energy product were obtained for the shorted annealing time of.125 h. Shorter times were not investigated as this is already so short that the sample is likely not reaching thermal equilibrium with the furnace before being removed. The evolution of the magnetic properties shown in Fig. 1 and discussed above are directly related to the evolution of the microstructure of the ribbons, including nucleation, growth, and decomposition of chemical phases as well as their morphologies. For selected annealing conditions, these were studied by x-ray diffraction (XRD) from the surfaces of the ribbons and by scanning electron microscopy (SEM) analysis of the interiors. The XRD and SEM results for samples annealed for at 5, 6, 7, 8 C are compared in Fig. 2. A smooth background has been subtracted from the raw xrd data. Before annealing, xrd and SEM showed the material to be amorphous, but with some compositional modulations which appeared to template the nucleation of crystallite phases [9]. The data in Fig. 2 show that the thermal treatments produce fine-grained crystalline material containing multiple chemical phases. Figure 2(a,e) shows that after at 5 C, crystallization of the amorphous phase is just beginning. The micrograph shows some texture developing in the amorphous material, which shows smooth contrast modulations on the length scale of about.5 µm. The xrd pattern contains a single broad feature near 44 degrees.

5 5 (a) (b) (d) heat flow (exo. = up) T ( C) (fcc) HfCo 7 (fcc) HfCo 7 (fcc) FIG. 2. X-ray diffraction patterns (a-d) and backscattered SEM micrographs (e-h) from ribbons annealed for at the indicated temperatures. The inset in (a) shows the thermal analysis results from Ref. 9. Diffraction peaks are indexed by markers corresponding to the phases listed in the legend in. Peaks which cannot be assigned to known phases in the Hf-Co-B system are marked with *. These are attributed to HfCo 7. (e) (f) (g) (h) Annealing for at higher temperatures produces well defined Bragg peaks in the diffraction patterns shown in Fig. 2(b,c,d). By comparison with reference data, the reflections were indexed to known phases in the Hf-Co-B system. Relative intensities as well as peak positions were considered when indexing the patterns. Matches to weak reflections from a candidate phase were discarded if the stronger lines from that phase were not observed. The legend in Fig. 2 lists the phases observed and identifies the peak markers used in the xrd panels. Peaks marked with * were not indexed by known structures, and are assigned to the HfCo 7 compound. What is known about the structure of this phase and its diffraction pattern will be discussed in more detail below. Figure 2(b,f) indicates that after at 6 C, little or no amorphous component remains. The SEM image shows the presence of at least three phases with different contrasts, with anisotropic grains ranging from about 1-5 nm in size. The xrd pattern indicates the presence of HfCo 7, Co 23 B 6, and HfCo 3 B 2 primarily, with a small amount of elemental Co as well. Clearly this sample has not reached thermal equilibrium. It is important to note that the incorporation of some boron into the nominally boron-free phases can neither be confirmed nor excluded based on the present data. After 1 at 7 C, reflections from Hf 6 Co 23 are also observed, as indicated in Fig. 2. Five distinct crystalline phases are identified in this diffraction pattern. The corresponding micrograph [Fig. 2(g)] is similar in appearance to that observed in the sample annealed at 6 C, in terms of grain sizes and contrast. Results of annealing at 8 C for are shown in Fig. 2(d,h). At this temperature, the HfCo 7 and Co 23 B 6 phases are absent. The majority phases are Co and Hf 6 Co 23, with a smaller amount of HfCo 3 B 2. A reflection near 28 degrees indicates the presence of some HfO 2, from reaction with vapor in the silica tube or by direct contact with the tube wall at this high temperature. Excluding the oxide, three phases are observed at this annealing temperature, indicating thermal equilibrium may have been reached. The SEM image shows similar grain sizes as seen in the other crystalline samples. Considering these observations in combination with the magnetic properties shown in Fig. 1, the occurrence of hard ferromagnetic behavior can be correlated with the presence of a significant amount of the HfCo 7 phase, which is present after of annealing at 6 or 7 C, but not after of annealing at 5 or 8 C. It is likely that the two thermal anomalies observed in the DTA data are both crystallization events, perhaps occurring separately in the two compositionally modulated regions noted in the amorphous precursor [9]. The absence of Co 23 B 6 at the highest temperatures is expected, since this is not a thermodynamically stable phase [28, 29]. Interestingly, none of the diffraction data indicate the presence of Hf 2 Co 7. Published phase diagrams show this to be the only thermodynamically stable Hf-Co compound between Co and HfCo 2 for temperatures below 95 C [23]. If the phase diagrams are correct, then either thermodynamic equilibrium is not reached in the annealed samples studied here, or the presence of boron significantly changes alters the phase stability in the Hf-Co system. It is not clear whether boron is incorporated into the HfCo 7 phase, and there is clear evidence in the diffraction patterns for boron forming secondary phases. However, boron appears to play an important role in the magnetic properties, in particular the coercivity of the melt-spun ribbons. Measured i H C values of melt-spun, boron-free HfCo 7 and Hf 2 Co 11 ribbons have been limited to about 2 koe [21, 3], while boron-containing ribbons produced under similar conditions have values up to koe [9, 21]. The role of boron may be to modify the intrinsic properties of the hard magnetic phase, or, perhaps just as likely, to modify the microstructural development through precipitation of secondary boride phases. Similarly, it was found that small boron additions resulted in coercivity enhancements in ZrCo 5 -based alloys, and maximum performance was achieved when these alloys

6 6 (a) J J (b) 5 o C, 26 h 3 5 o C, 28 h (d) 5 o C, FIG. 3. X-ray diffraction patterns and SEM images for samples annealed at 5 C. The HfCo 7 phase is coarsened as the annealing time is increased. The images show that this phase forms as anisotropic, plate-like grains. were annealed at 5 7 [12]. Since at 6 C multiple phases are present in significant quantities, the formation of the HfCo 7 phase was explored further by examining the samples annealed for longer periods of time at a lower temperature of 5 C. The relevant xrd patterns and and micrographs are shown in Fig. 3. Results are shown for annealing times of 1, 28, and 26 h. As noted above, after little change is seen from the amorphous starting material, though some crystallization is evident in the SEM image, reproduced in Fig. 3(d). After annealing for 28 h at this temperature, the xrd patten in Fig. 3(a) indicates a significant crystalline fraction, producing sharp Bragg peaks. The HfCo 7 phase forms as a network of inter-grown, platelike crystallites, shown in Fig. 3. A thin-plate morphology has also been reported for the related Zr material [31]. Annealing for a longer period of time at 5 C allows more Bragg peaks associated with HfCo 7 to be observed. Some further coarsening of the microstructure is suggested by comparison of Figs. 3(b) and. In the materials annealed for shorter times at higher FIG. 4. Similar demagnetization curves from materials with clearly distinct phase compositions. The sample annealed at 8 C for a short time contains a large amount of Hf 6Co 23, which is absent in the sample annealed at 6 C for [see Fig. 2(b)]. Coercivity approaching 2 koe in a material containing a large fraction ( 5%) of cubic ferromagnet Hf 6 Co 23 suggests that this phase may be favorable to high energy products in these complex alloys. temperatures, a significant fraction of HfCo 7 is associated with high coercivity values. However, the samples annealed for 28 or 26 h at 5 C show large amounts of HfCo 7 but little coercivity. This is likely due in part to effects of grain size, which is significantly larger in longer annealed samples. Also, the presence of multiple phases may play a key role in the hard magnetic properties in this family of alloys. Comparison of magnetization curves in Fig. 1 shows that very similar behavior can be achieved by thermal treatments which differ substantially. This is demonstrated in Fig. 4, which compares ribbons annealed for 1 h at 6 C with those annealed for.125 h at 8 C. The J vs H behavior differ somewhat in saturation values, but the demagnetization curves are quite similar, with exception of the small kink near H = in the sample annealed at 8 C. The associated xrd patterns in Fig. 4 show that the primary difference between these two materials is the large fraction of Hf 6 Co 23 present in the sample annealed at 8 C. It is surprising to see i H C approaching 2 koe in a sample containing at about 5% of this compound, which is expected to be a soft ferromagnet due to its non-uniaxial (cubic) crystallographic symmetry. Based on the observed remanence and the shape of

7 7 AuBe 5 -type HfCo 5 F -4 3 m a = 6.72 Å θ (deg.) FIG. 5. Powder diffraction patterns from material annealed for 26 h at a temperature of 5 C, and material that was melt-spun at a lower temperature, resulting in a crystalline ribbon without post-annealing [9]. Reflections marked with * are attributed to HfCo 7 phase. The sets of tick marks below the patterns locate reflections from proposed unit cells and symmetries [19, 32 34]. The inset shows vertical bars indicating positions and intensities of reflections from hypothetical HfCo 5 in the cubic AuBe 5 structure type. the demagnetization curve, the sample annealed at 8 C does not appear to have large soft component; only a small kink is observed near H =. This means the magnetization reversal in the presumably soft Hf 6 Co 23 phase is impeded, which suggests exchange coupling with the hard HfCo 7 phase may be occurring. It is unfortunate that the crystal structure of the hard magnetic phases with compositions MCo x with x = 5 7 is not well determined. It may be likely that M = Zr and Hf phases are isostructural with small differences in lattice constants. This is supported by the similarity in reported diffraction patterns from both materials [8 11, 15 18, 21], and the fact that Zr and Hf are two of the most similar elements on the periodic table with respect to their chemical behavior. Several structural studies have been reported, some using xrd and some using electron diffraction, and are briefly summarized here. A tetragonal structure was first proposed for HfCo 7 with lattice constants a = 7.7 Å and c = Å from xrd [32]. An electron diffraction study concluded that Zr 2 Co 11 and HfCo 7 are both orthorhombic with space group P cna and lattice constants for the Zr compound of a = 4.8 Å, b = 8.2 Å, c = 36 Å and for the Hf compound of a = 4.7 Å, b = 8.3 Å, c = 38 Å [33]. An electron diffraction and xrd study on Zr 2 Co 11 and ZrCo 5.1 found evidence for polymorphism in these materials [34]. A high temperature rhombohedral phase with hexagonal lattice constants a = 4.67 Å and c = 2.42 Å and a low temperature orthorhombic phase with a = 4.71 Å, b = 16.7 Å and c = 24.2 Å were reported. A possible additional high temperature phase of hexagonal symmetry was also indicated. It was suggested that the high temperature rhombohedral phase may be related to the cubic AuBe 5 structure [8, 34]. Most recently, another orthorhombic unit cell was proposed, based on xrd from HfCo 7, with lattice constants a = Å, b = Å, and c = 8.75 Å [19]. A structural model was proposed in Ref. 35 based on close packing and the 1:7 stoichiometry. Figure 5 shows xrd patterns from samples from the present study which contain the largest fraction of the HfCo 7 phase. Features in the patterns confirmed to be associated with this phase are marked with *. These patterns are similar to those reported in the literature for these materials [8 11, 15 18, 21]. The difficulty in assigning a unit cell to this phase with confidence is illustrated in this figure. The tick marks locate Bragg reflections from the unit cells proposed in the literature and listed above. In each case, even for the smallest unit cells, there are numerous reflections which are allowed but not observed. In such a case, comparison of relative intensities of Bragg peaks from a particular phase usually aids in the identification of the correct structure. Without atomic positions for the proposed structures, they cannot be distinguished with confidence. As noted above, the AuBe 5 -type has been suggested as a possible parent structure. The inset in Fig. 5 shows the the xrd pattern from the sample annealed for 26 h at 5 C, and vertical bars marking the position and intensity of reflections from HfCo 5 using this cubic structure type with Hf replacing Au at Co replacing Be, using the lattice constant reported in Ref. 8 for the Zr compound. A few of the peak positions match, and the intensity distribution is reasonable at lower angles. Clearly the structure is more complicated than this prototype, but the data seem to indicate that this may indeed be a reasonable model with which to explore supercells of lower symmetry in future studies of these compounds, as has been suggested previously [8, 34]. The close-packed model proposed in Ref. 35 can be generated by using space group B222 with the reported unit cell of a = Å, b = Å, c = 8.7 Å by placing Hf at (1/2, 1/2, ), Co1 at (1/2, 1/2, 1/2), Co2 at 1/4,, 1/2), Co3 at (1/4, 1,/2, 1/4), and Co4 at (,,

8 (a) (b) (d) 6 C,, µ H = 9 T 6 C, 5 C, 26 h 6 C,, µ H = 9 T 6 C, 5 C, 26 h role of aligning ferromagnetic particles, strong magnetic fields can have significant effects on the thermodynamics and kinetics of phase transformations in both magnetic [36, 37] and nonmagnetic materials [38], including diamagnetic polymers [39]. Some of the results are illustrated in Fig. 6, which compares the microstructure and demagnetization curve from a sample annealed at 6 C for in a magnetic field of 9 T to those from samples annealed in zero magnetic field. The thermo-magnetically processed sample in Fig. 6(a) shows well-defined platelike crystallites which extend to length scales near. This microstructure is distinct from that observed in the sample annealed at the same temperature and for the same amount of time but with no applied magnetic field [Fig. 6(b)]. This demonstrates a large influence of the magnetic field on the microstructural evolution of the amorphous precursor material. The microstructure shown in Fig. 6(a) resembles closely that obtained by annealing at a lower temperature of 5 C for a significantly longer period of time (see Fig. 3). For comparison, Fig. 6 shows the sample annealed for 26 h at 5 C. The corresponding demagnetization curves are shown in Fig. 6(d), and confirm the strong effect of the magnetic field applied during annealing. This shows that, in addition to annealing temperature and time, magnetic field is a viable tuning parameter that can be exploited during the development of the magnetic properties that accompanies microstructural evolution. In the case demonstrated in Fig. 6, the field appears to accelerate the coarsening of the microstructure, which indicates that the properties could perhaps be improved by shortening the annealing time and/or temperature with the field applied. This technique may allow access to microstructures and corresponding permanent magnet performance that are otherwise unattainable in Zr/Hf-Co hard ferromagnets. 8 IV. SUMMARY AND CONCLUSIONS FIG. 6. Backscattered SEM images comparing the ribbons annealed at (a) 6 C for in a magnetic field of 9 T, (b) 6 C for in no magnetic field, and 5 C for 26 h in no magnetic field. The corresponding demagnetization curves are shown in (d). 1/4). The authors note that this is only an approximate model, and the true structure may be more complex. The powder diffraction pattern generated from this structural model using Cu radiation gives the strongest intensity reflection near 49.5 degrees, which is incompatible with the measured diffraction patterns. However, Rietveld fitting of the HfCo 7 phase has been reported using this as a starting model, though the atomic coordinates obtained from the fit are not reported [21]. The effect of high magnetic fields applied during the annealing process was also investigated. Beyond the In summary, we have investigated effects of thermal and thermo-magnetic treatment of amorphous Hf 2 Co 11 B alloys, demonstrating the evolution of the magnetic properties, phase fractions, and microstructure. The microstructure evolves slowly at 5 C, forming primarily the HfCo 7 phase. Long annealing times at 5 C gives mostly crystalline and nearly single phase HfCo 7 with relatively large grains that result in soft ferromagnetic behavior. At higher temperature other Hf-Co-B phases are formed, and after at 8 C only HfCo 3 B 2, Hf 6 Co 23, and Co are observed. The energy product BH max is maximized by heat treatment at 7 C for approximately 1 h, and the best energy products are found in samples containing multiple phases including HfCo 7 with a finegrained, crystalline microstructure. The resulting remanence, coercivity, and magnetic energy product are significantly lower than those obtained in optimally melt-spun, crystalline materials. The associated microstructures differ as well. This indicates that the non-equilibrium evo-

9 9 lution occurring during melt-spinning is key to development of microstructures most favorable for permanent magnets. Application of a high magnetic field during annealing dramatically changes the microstructure, and may provide a route to properties not attainable by thermal treatment alone. Comparison of the diffraction peaks associated with the HfCo 7 phase with proposed crystal structures for (Zr/Hf)Co 5 7 materials is inconclusive, indicating the need for further crystallographic studies of this family of rare-earth free permanent magnet materials. ACKNOWLEDGEMENTS Research sponsored by the U. S. Department of Energy, Office of Energy Efficiency and Renewable Energy, Vehicle Technologies Office, as part of the Propulsion Materials Program. Microscopy work supported by ORNL SHaRE, Division of Scientific User Facilities, Office of Basic Energy Sciences, U. S. Department of Energy. [1] R. C. O Handley, Modern Magnetic Materials: Principles and Applications, 1st ed. (John Wiley and Sons, Inc., New York, 2). [2] J. F. Herbst, Rev. Mod. Phys. 63, 819 (1991). [3] Y. Matsuura, IEEE Trans. Appl. Supercond. 1, 883 (2). [4] K. J. Strnat and R. Strnat, J. Magn. Magn. Meter. 1, 28 (1991). [5] A. S. Rao, Proc. of IEEE conf. on EEIC/ICWA, Chicago, 373 (1993). [6] U.S. Department of Energy, Critical Materials Strategy (211). [7] T. Ishikawa and K. Ohmori, IEEE Trans. Magn. 26, 137 (199). [8] A. M. Gabay, Y. Zhang, and G. C. Hadjipanayis, J. Magn. Magn. Meter. 236, 37 (21). [9] M. A. McGuire, O. Rios, N. J. Ghimire, and M. Koehler, Appl. Phys. Lett. 11, 2241 (212). [1] H. W. Chang, Y. H. Lin, C. W. Shih, W. C. Chang, and C. C. Shaw, J. Appl. Phys. 115, 17A724 (214). [11] B. Balamurugan, B. Das, W. Y. Zhang, R. Skomski, and D. J. Sellmyer, J. Phys.: Condens. Matter 26, 6424 (214). [12] T. Saito, Appl. Phys. Lett. 82, 235 (23). [13] L. Y. Chen, H. W. Chang, C. H. Chiu, C. W. Chang, and W. C. Chang, J. Appl. Phys. 97, 1F37 (25). [14] J.-B. Zhang, Q.-W. Sun, W.-Q. Wang, and F. Su, J. Alloys Cmpd. 474, 48 (29). [15] M.-Y. Zhang, J.-B. Zhang, C.-J. Wu, W.-Q. Wang, and F. Su, Physica B 45, 1725 (21). [16] W. Zhang, S. R. Valloppilly, X. Li, R. Skomski, J. E. Shield, and D. J. Sellmyer, IEEE Trans. Magn. 48, 363 (212). [17] I. A. Al-Omari, W. Y. Zhang, L. Yue, R. Skomski, J. E. Shield, X. Z. Li, and D. J. Sellmyer, IEEE Trans. Magn. 49, 3394 (213). [18] H. W. Chang, C. F. Tsai, C. C. Hsieh, C. W. Shih, W. C. Chang, and C. C. Shaw, J. Magn. Magn. Mater. 346, 74 (213). [19] B. Balamurugan, B. Das, V. R. Shah, R. Skomski, X. Z. Li, and D. J. Sellmyer, Appl. Phys. Lett. 11, (212). [2] B. Balasubramanian, B. Das, R. Skomski, W. Y. Zhang, and D. J. Sellmyer, Adv. Mater. 25, 69 (213). [21] H. W. Chang, M. C. Liao, C. W. Shih, W. C. Chang, C. C. Chang, C. H. Hsiao, and H. Ouyang, Appl. Phys. Lett. 15, (214). [22] E. F. Kneller and R. Hawig, IEEE Trans. Magn. 27, 3588 (1991). [23] K. Ishida and T. Nishizawa, in Binary Alloy Phase Diagrams, edited by T. B. Massalski (ASM International, 199) 2nd ed. [24] C. Kittel, Introduction to Solid State Physics, 6th ed. (John Wiley & Sons, Inc., New York, 1986). [25] S. C. Bedi and M. Forker, Phys. Rev. B 47, (1993). [26] H. Ido, M. Nanjo, and M. Yamada, J. Appl. Phys. 75, 714 (1994). [27] P. Ohodnicki, Jr., N. C. Cates, D. E. Laughlin, M. E. McHenry, and M. Widom, Phys. Rev. B 78, (28). [28] P. K. Liao and K. E. Spear, in Binary Alloy Phase Diagrams, edited by T. B. Massalski (ASM International, 199) 2nd ed. [29] D. Kotzott, M. Ade, and H. Hillebrecht, J. Solid State Chem. 182, 538 (29). [3] M. A. McGuire, Unpublished,. [31] G. V. Ivanova and N. N. Shchegoleva, Phys. of Metals and Metallography 17, 287 (29). [32] K. H. J. Buschow, J. Appl. Phys. 53, 7713 (1982). [33] B. G. Demczyk and S. F. Cheng, J. Appl. Cryst. 24, 123 (1991). [34] G. V. Ivanova, N. N. Shchegoleva, and A. M. Gabay, J. Alloys Compd. 432, 135 (27). [35] B. Das, B. Balamurugan, P. Kumar, R. Skomski, V. R. Shah, J. E. Shield, A. Kashyap, and D. J. Sellmyer, IEEE Trans. Magn. 49, 333 (213). [36] Z. H. I. Sun, M. Guo, J. Vleugels, O. Van der Biest, and B. Blanpain, Curr. Opin. Solid State Mater. Sci. 16, 254 (212). [37] S. Rivoirard, JOM-J. Min. Met. Mat. S. 65, 91 (213). [38] Z. H. I. Sun, M. Guo, J. Vleugels, O. Van der Biest, and B. Blanpain, Curr. Opin. Solid State Mater. Sci. 17, 193 (213). [39] Y. Li, O. Rios, and M. R. Kessler, ACS Appl. Mater. Interfaces 6, (214).

10 J (kg) J (kg) BH (MGOe) (a) 6 h 28 h 26 h 4.5 h.5 h (b) 5 C (e) 6 C (h) 7 C (k) 8 C C (d) 6 C (g) 7 C (j) 8 C 5 C (f) 6 C (i) 7 C (l) 8 C.125 h.25 h.5 h

11 (a) heat flow (exo. = up) (e) T ( C) (b) (f) (fcc) HfCo 7 (g) (fcc) HfCo 7 (d) (h) (fcc)

12 (a) (b) 5 o C, 26 h 5 o C, 28 h (d) 5 o C,

13 J 3 J

14 AuBe 5 -type HfCo 5 F -4 3 m a = 6.72 Å θ (deg.)

15 (a) 6 C,, µ H = 9 T (b) 6 C, 5 C, 26 h (d) 6 C,, µ H = 9 T 6 C, 5 C, 26 h