The properties of melt-spun NdFeB alloys

The properties of melt-spun NdFeB alloys 3.0

3

Introduction

Within several years of their discovery, melt-spun NdFeB magnetic powder were being used for the production of bonded Nd magnets and were being incorporated into a wide range of products, primarily motors for computer peripheral, consumer electronic, office automation, and automotive applications. This chapter discusses the basis magnetic properties, microstructure, and annealing behavior of rapidly solidified NdFeB materials that were first reported by Croat et al. (1984a,b). In particular, this chapter presents how the properties and microstructure change as the quench rate is varied and current theories on what causes the dramatic changes that are observed. Also presented is how the properties of melt-spun NdFeB change when the Nd and Fe are replaced with other rare earths and transition metals (TMs). Finally, the magnetization and demagnetization process that is believed applicable for these isotropic materials is described, including the important role that intergrain interactions are now believed to play in determining the magnetic properties. The final section in this chapter discusses nanocomposite permanent magnet, which is also melt-spun material that contains the Nd2Fe14B phase but also contains α-Fe and Fe3B. The high-volume commercial production of melt-spun NdFeB magnetic powder is covered in Chapter 4, Production of rapidly solidified NdFeB magnetic powder, and the commercial production of bonded Nd magnets by various techniques is discussed in Chapter 5, Production and properties of bonded Nd magnets.

3.1

The melt-spinning process

Melt spinning is a process in which a thin stream of molten alloy is directed against the outer surface of a cold, rapidly rotating rim. This process was also discussed in some detail in Chapter 2, The Nd2Fe14B intermetallic compound, and a photograph of the laboratory melt spinner employed to investigate NdFeB alloys is shown in Fig. 2.6. Much of the data discussed in this chapter are concerned with the properties of melt-spun NdFeB as a function of quench rate and, to the extent possible, all the variables that could change the quench rate were carefully controlled. This involved fixing the amount of sample used (10 g), the crucible orifice diameter (gauged to 1 mm), and the melt temperature. Because of the small amount of sample employed, the hydrostatic head of molten alloy was insufficient to drive the molten alloy from the crucible. Therefore, the crucible was capped and the molten alloy driven from the crucible with pressurized high-purity argon. The argon Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets. DOI: http://dx.doi.org/10.1016/B978-0-08-102225-2.00003-X Copyright © 2018 Elsevier Ltd. All rights reserved.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

ejection pressure was keep as constant as possible for each experiment. Because all rare earth alloys are easily oxidized, it was necessary to carry out the melt spinning in a vacuum chamber that was evacuated and back filled with pure argon. The quench rate was varied by simply changing the disk surface velocity (vs), which is equal to the meters of substrate surface that passes beneath the orifice in 1 second. Therefore, the properties are often shown as a function of substrate velocity (vs) in m/s. To a first approximation, vs is believed to be proportional to the quench rate. Although the quench rate is not precisely known, it is believed to be as high as 105 K/s at the highest wheel speeds. Melt-spun ribbon produced was typically 25
40 μm thick and B1.5 mm wide. As would be expected, melt-spun ribbon at higher quench rates was thinner. A number of different laboratory melt spinners were built at the General Motors Research Laboratories during the late 1970s and early 1980s to develop NdFeB magnetic powder. An example of one such device is shown in Fig. 3.1. This laboratory melt spinner was referred to as the “carrousel,” since six different crucibles could be indexed into the melt-spinning position during a single turn or set-up of the machine. The 1.0 in. diameter quartz crucible could hold up to 100 g of alloy and could, in combination, produce 600 g of a single composition or 100 g each of a series of different compositions. Again, all of the variable were kept as constant as possible and the orifice diameters were always gauged to 1 mm. Although this photo shows only one collection bin, it could also be fitted with six separate bins. This melt spinner was used to produce enough magnetic powder to produce early prototype motors for General Motors using bonded Nd magnets. One unexpected problem found during these early studies was that the quenched ribbon coming off the quench wheel was so hot that it heated the argon gas in the melt spinner to the

Figure 3.1 One of the several different laboratory melt spinners used at the General Motors Research Laboratories for the development of melt-spun NdFeB magnetic powder. Source: Courtesy Tim Trueblood.

The properties of melt-spun NdFeB alloys

67

extent that back pressure would reduce the flow of molten alloy from the crucible. This, of course, would change the quench rate and, in turn, the magnetic properties as the run proceeded. To prevent this, a sensitive pressure relief valve had to be added to the melt spinner to vent the hot argon gas. With the addition of the relief valve, more consistent magnetic properties could be obtained. Measurements were made to determine how hot the melt-spun ribbon actually was as it departed the quench rim and, in most instances, it was found to be over 700 C. As might be expected, this varied with quench rate, with ribbon quenched at the higher rate being cooler. In retrospect, it now seems naive that the temperature of the as quenched ribbon was not cooled to near room temperature. However, in the earliest stages of the research, the samples were small and, when ejected into a metal bin, were always cool to the touch. Besides, they had been quenched, and it seemed logical that they would be cool. It was only when larger amounts of material were melt spun that the hot nature of the fresh melt-spun ribbon became apparent. A pile of the melt-spun ribbon was found had a very low thermal conductivity and this presented a serious engineering challenge in trying to cool production quantities of ribbon. This issue, along with other technical problems encountered during the development of high-volume production melt spinning, is discussed in Chapter 5, Production and properties of bonded Nd magnets.

3.1.1 The characteristics of an isotropic magnetic material As was discussed in Chapter 2, The Nd2Fe14B intermetallic compound, these meltspun NdFeB materials are crystallographically and, hence, magnetically isotropic and therefore do not have a square-shaped second quadrant demagnetization curve that is characteristic of an anisotropic permanent magnet. In contrast, isotropic materials exhibit a monotonic change in M in the second quadrant as the field is varied. Fig. 2.7 displays the full hysteresis curve that is typical for these isotropic materials and Fig. 2.8 shows the second quadrant demagnetization characteristics of melt-spun NdFeB. Two important properties that are derived from the second quadrant are the remanance or residual induction Br, defined as the magnetization at zero applied field and the intrinsic coercivity Hci, defined as the value of the applied field at which the magnetization becomes zero. Also shown in Fig. 2.8 is the magnetic induction B, defined as M 2 H. The reverse field required to bring the induction to zero is called the inductive coercivity Hc. From the induction, we can calculate the maximum energy product of the magnet (BH)max, defined as the maximum product of B and H along the B curve and proportional to the amount of useful work that can be performed by the magnet. Energy product is the figure of merit almost universally used to compare various grades of permanent magnet materials. The greater the energy product, the smaller the volume and weight of the magnet required for a given application. Another characteristic of these isotropic materials is the very high magnetizing field that is required to completely develop the second quadrant magnetic properties. This is due to the very high magnetocrystalline anisotropy of the Nd2Fe14B intermetallic compound that resists rotation of the magnetic moment into the field direction. All of the magnetic data discussed in this chapter are obtained from a standard vibrating sample magnetometer (VSM) with a

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

field strength of 19 kOe. However, this field was found to be far too low to achieve full magnetization and all of the samples were premagnetized in a pulsed magnetic field of B4.5 T. It is important to note that no demagnetization correction is applied to any of the data in this chapter.

3.2

Properties of melt-spun NdFeB alloys

3.2.1 Magnetic properties This section discusses the magnetic properties, microstructure, and physical properties of melt-spun NdFeB alloys. Most of the data are from these studies can be found in the following publications: Croat et al. (1984a,b), Croat (1988), Croat and Herbst (1988), Croat (1989a,b), Herbst and Croat (1991) as well as in several of the composition and process patents that were filed on these materials (Croat, US Patent 4,802,931, issued 1989 and Croat, US Patent 5,172,751, issued 1992). Much of these data are also discussed in a general review of R2Fe14B intermetallic compounds and NdFeB magnets by Herbst (1991). Figs. 3.2
3.6 display the magnetic properties of various NdFeB alloys as a function of the substrate velocity vs or quench rate of the melt spinner. Fig. 3.2 shows a plot of the intrinsic coercivity versus vs for a series of Nd0.15(Fe12yBy)0.85 alloys. The data show a maxima in Hci for

20

y = 0.07

Hci (kOe)

15 y = 0.05 y = 0.09

10

y = 0.03 5 Nd0.15(Fe1–yBy)0.85 0

0

10

20

30

VS (m/s)

Figure 3.2 The intrinsic coercivity of Nd0.15(Fe12yBy)0.85 alloys versus vs, the substrate surface velocity of the melt spinner quench rim (Croat et al., 1984a).

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69

Figure 3.3 Cu Kα X-ray diffraction spectra of melt-spun samples for y 5 0 (A), 0.3 (B), 0.5 (C), 0.7 (D), and 0.9 (E). All of the samples were melt-spun at a quench rate of 15 m/s (Croat, US Patent 4,802,931).

substrate velocities between 15 and 20 m/s, with values sharply lower at both lower and higher quench rates. Coercivity also increases dramatically with increasing boron content. The highest coercivity levels between 17 and 20 kOe were achieved for boron levels between y 5 0.5 and 0.9, with a maximum value of B20 kOe for y 5 0.7. As discussed in more detail later, these materials consist of a two-phase microstructure with a major Nd2Fe14B intermetallic phase and a minor Nd-rich (BNd0.7Fe0.3) intergranular phase. This later phase surrounds the crystallites of the Nd2Fe14B primary phase and is the source of the coercive force in these materials. If this phase is absent, the material will exhibit almost no coercivity. The maxima in these data are believed to results from both the change in the average crystallite size of the Nd2Fe14B intermetallic phase and the thickness of the Nd-rich intergranular phase surrounding each crystallite as the quench rate is varied. The increase in Hci with increasing boron content (y) is due to an increase in the volume content of these two phases in the microstructure of the melt-spun samples at the expense of the equilibrium Nd2Fe17 and α-Fe phases.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

16 14

(BH)max MGOe

12

Nd0.135(Fe0.945B0.055)0.865

10 8 6 4 2 0 10

15

20

25 VS (m/s)

30

35

40

Figure 3.4 The remanance (Br) of a series of Nd0.135(Fe0.945B0.055)0.855 alloys versus vs, the substrate surface velocity of the melt spinner quench wheel (Croat et al., 1984b).

Nd0.135(Fe0.945B0.055)0.865

M (kG) 10

8

Optimum-quench VS = 19 m/s 14.0 MGOe

6 Under quenched VS = 14 m/s 6.9 MGOe

4

Overquenched VS = 21.7 m/s 2.2 MGOe

2

–15

–10

–5 H (kOe)

0

5

Figure 3.5 Second quadrant demagnetization curves for Nd0.135(Fe0.945B0.055)0.865 alloys as a function of substrate velocity vs. Shown here is the difference in the demagnetization characteristics for optimum, underquenched and overquenched alloys (Croat et al., 1984b).

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71

10

Nd0.135(Fe0.946B0.054)0.865 VS = 19 m/s 14.0 MG•Oe

8

VS = 20 m/s 12.1 MG•Oe

6

VS = 20.5 m/s 9.0 MG.Oe

4 VS = 35 m/s 0.0 MG•Oe 2

VS = 21.7 m/s 2.2 MG•Oe –15

–10

–5 H (kOe)

VS = 27.5 m/s 0.6 MG•Oe 0

5

Figure 3.6 Second quadrant demagnetization curves for Nd0.135(Fe0.945B0.055)0.865 alloys as a function of substrate velocity vs (Croat et al., 1984b).

The development of the Nd2Fe14B phase with increasing boron content is shown clearly in Fig. 3.3, which displays X-ray diffraction (Cu Kα) data for the melt-spun samples from Fig. 3.2 (y 5 0.3, 0.5, 0.7, and 0.9) as well as a melt-spun sample containing y 5 0. All of these samples were melt spun at a quench rate of 15 m/s, which is close to the substrate velocity (vs) giving the highest or close to the highest Hci values for all the compositions tested. The X-ray data for the y 5 0 sample reflect the spectra of the equilibrium microstructure for this composition, which consists of α-Fe and the Nd2Fe17 intermetallic phase. No data for this sample are shown in Fig. 3.2 because this alloy exhibited no measurable coercivity over the entire range of vs. With increasing boron, the spectra show a decreasing amount of the Nd2Fe17 and α-Fe phases and an increasing amount of the Nd2Fe14B phase. For values of y between 0.5 and 0.9, only the spectra of the Nd2Fe14B intermetallic phase are observed. There are no observable Bragg reflections corresponding to the Nd-rich intergranular phase. It is noted, however, that the spectra exhibit some line broadening due to the disorder introduced by the rapid solidification, which would tend to mask small amounts of secondary phases present. It is also noted here that the Nd2Fe14B phase contains 5.88 at% boron, while the sample with composition y 5 0.07 contains 5.95 at% boron. Therefore this sample would be expected to contain predominately the Nd2Fe14B phase. Fig. 3.4 displays the energy product (BH)max, for Nd0.135(Fe0.945B0.055)0.865 alloys as a function of quench rate. A maximum value of 14 MGOe is found at a quench rate corresponding to vs 5 19 m/s. The energy product drops dramatically on either side of this value and decreases to negligible values at quench rates

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

greater than 30 m/s. This pronounced variation is a reflection of the equally dramatic change in the shape of the second quadrant demagnetization curve of the material as the quench rate is varied. This feature is shown in Fig. 3.5, which displays the demagnetization curves of several samples selected from Fig. 3.4. These data were taken over a magnetic field range of 15 kOe to 215 kOe because the first-quadrant data provide valuable information about the quench condition of the samples. As with all the data shown here, the samples were premagnetized at B4.5 T prior to conducting the VSM test. Included in Fig. 3.5 are the data for the samples prepared at a quench rate of 19 m/s and referred to as “optimum,” since it was found to have the maximum energy product of 14 MGOe found for this series of samples. Also shown are the demagnetization curves of samples prepared at both a lower quench rate of 14 m/s and referred to as “underquenched” and a higher quench rate of 21.7 m/s and referred to as “overquenched.” The sample prepared at 14 m/s is very typical of an underquenched material in that the magnetization of the sample is seen to be significantly lower across the entire field range tested when compared to the optimum sample. In contrast, the overquenched sample, prepared at 21.7 m/s, exhibits both higher first-quadrant magnetization combined with much lower coercivity. Fig. 3.6 again shows the demagnetization characteristics of the samples from Fig. 3.4 but, in this case, over a wider range of quench rate, from 19 to 35 m/s. The samples prepared at 20 and 20.5 m/s are often referred to as slightly too moderately overquenched. At the highest quench rate of 35 m/s, the coercivity almost completely disappears. For comparative purposes, the data also show the demagnetization characteristics of the NdFeB ingot from which these samples were prepared. This material had a coercive force of B400 Oersted. It is noted that the dependence of the first-quadrant magnetization with quench rate seen in Figs. 3.5 and 3.6 is still used in producing commercial quantities of melt-spun NdFeB powders for bonded Nd magnets. The production of underquenched material is carefully avoided because of the lower remanance, and its presence can be best detected by it first-quadrant behavior when measured on a VSM. As discussed later, overquenched materials can be annealed to improve the magnetic properties by growing the Nd2Fe14B to a size more commensurate with the single domain size. However, underquenched material cannot be annealed, since the average Nd2Fe14B grain size is already larger than the single domain size. Fig. 3.7 summarizes the properties that are typically found for melt-spun NdFeB alloys as a function of quench rate. Displayed here is the energy product (BHmax), remanance (Br), intrinsic coercivity (Hci), and inductive coercivity (Hc) for series of melt-spun Nd12x(Fe0.95B0.05)0.87 alloys versus vs. The data shown here are for the sample providing the highest energy product (optimum with respect to vs) for each composition. For these melt-spun NdFeB alloys, the best properties are always found at a composition slightly Nd-rich relative to the composition of the Nd2Fe14B1 intermetallic phase. For this series of alloys, the highest remanance and energy product were obtained at x 5 0.87 or a Nd composition of B13 at% Nd, corresponding to a 28.68 wt.%, while the maximum Hci was found at a higher Nd content. This is not surprising since, in this two-phase microstructure, the remanance and energy product are proportional to the volume fraction of the Nd2Fe14B primary

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73

16

12

8

BHmax (MG.Oe)

Hc (kOe)

4 Br (kG) 0 Hci (kOe)

16

8

Nd1.x(Fe0.95B0.05)x

0 0.7

0.8

0.9

x

Figure 3.7 Energy product (BH)max, remanance Br, inductive coercivity Hc, and intrinsic coercivity Hci of melt-spun Ndx(Fe0.95B0.05)12x alloys prepared at the optimum-quench rate (Croat et al., 1984a).

phase while the coercivity results from a slight excess of Nd, which produces the Nd-rich intergranular phase that surrounds each Nd2Fe14B crystallite. As mentioned, this intergranular phase is closely related to the coercive force in these melt-spun materials. The Nd-rich intergranular phase is generally thought to be paramagnetic phase that should not contribute to the remanance or energy product of these melt-spun materials. However, it is noted that recent studies by Kohashi et al. (2014), Nakamura et al. (2014), and Murakami et al. (2014) have reported that the Nd-rich phase in sintered Nd magnets is actually ferromagnetic. Studies of the intergranular phase by Liu et al. (2014, 2015), using three-dimensional atomic probe analysis (3DAP), have also reported that the Nd content is much lower than found in earlier studies. It seems likely that if the phase is ferromagnetic in sintered Nd magnets, then it is likely to be the same in these fine-grained melt-spun materials. These results and the implications regarding the magnetic properties of all types of NdFeB magnets are still being debated. Fig. 3.8 shows data for a series of melt-spun Nd12x(Fe0.95B0.05)x alloys. These data reflect the change in overall magnetic properties as the Fe:Nd ratio is changed.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

10 Nd1-x(Fe0.95B0.05)x 8 x = 0.866 x = 0.9

M (KG)

6

x = 0.8

4 x = 0.75 2

x = 0.9

–15

–10

–5

0

5

H (KOe)

Figure 3.8 Room temperature demagnetization curves of Nd12x(Fe0.95B0.05)x alloys prepared at optimum-quench rate (Croat, 1988).

For each composition, samples were melt spun over a range of quench rate (vs) and the data shown are for the optimum samples with the highest energy products. With increasing Nd composition (decreasing x), there is a steady increase in Hci and a corresponding decrease in Br. These data demonstrate one of the most important properties of all NdFeB magnets, namely, the ability to change the properties of the materials by changing the Nd:Fe ratio and is the basis for the various grades of NdFeB magnetic powder that are commercially available today. These grades of powder are discussed in Chapter 4, Production of rapidly solidified NdFeB magnetic powder. Again, this change reflects the change in the two-phase microstructure of these melt-spun materials. With increasing Nd content, the thickness of the Nd-rich boundary layer is believed to increase, resulting in higher Hci. In contrast, decreasing the Nd content is believed to result in a thinner Nd-rich boundary layer and a higher volume fraction of the Nd2Fe14B intermetallic phase, resulting in higher remanance and energy product but lower Hci. The remanance Br and energy product (BH)max are related to the volume fraction of the primary Nd2Fe14B phase. As discussed in Section 3.3.5, the change exhibited in the magnetic properties of this melt-spun NdFeB material with quench rate is now believed to be due to intergrain interactions, specifically exchange interaction, which can change significantly as the amount of the Nd-rich intermetallic phase is varied. As the amount of Ndrich intergranular phase is reduced, exchange interaction increases, leading to an increase in Br and magnetization but also to a dramatic drop in the coercive force.

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3.2.2 The microstructure of melt-spun NdFeB alloys Transmission electron microscopy (TEM) studies of optimum melt-spun NdFeB ribbons were first carried out by Chen (1985) and Mishra (1986). Lorentz TEM investigations have also been carried out by Hadjipanayis et al. (1985) and Hadjipanayis and Gong (1987, 1988b). The TEM of a melt-spun ribbon with an energy product of B12 MGOe and composition of Nd0.135Fe0.810B0.055 is shown in Fig. 3.9 (Mishra, 1986). The microstructure was found to consist of extremely small polyhedron or round-shaped grains of the Nd2Fe14B intermetallic phase with an average grain size of B30 nm, or 0.03 μm. It is noted here that this grain size is more than 100 times smaller than the 2
5 μm average grain size found applicable for sintered Nd magnets. However, as is apparent, there is a wide variation in the grain size, with some grains smaller than 5 nm and others larger than 80 nm. These grains are surrounded by a thin layer (1
2 nm) of a Nd-rich phase that was reported to be probably amorphous (Mishra, 1986). One such grain boundary is shown in the TEM image in Fig. 3.10A. This study also observes domain walls in the melt-spun material and found that the domain walls in the thermally demagnetized sample examined were pinned at the grain boundary phase and found to extend around a number of grains. These domains are referred to as “extended domains” in this study but are now universally referred to as “interaction domains.” While there appears to be obvious domain wall pinning at the Nd-rich grain boundary phase in these TEM images, this does not answer the question of how new domain walls form during demagnetization of material that has been magnetically saturated. The coercivity mechanism in these fine-grained melt-spun materials is discussed in Section 3.4.3. The average grain size in these melt-spun materials is highly dependent on quench rate, with material quenched at a higher rate having a much smaller average grain size. This is clearly shown in Figs. 3.11
3.13, which show examples of scanning electron micrographs (SEMs) of melt-spun ribbons quenched over a wide range of quench condition. These micrographs were taken from fracture surfaces of the melt-spun ribbons and represent regions from the top one/third of the ribbon,

Figure 3.9 Transmission electron micrograph (TEM) of a melt-spun Nd-Fe-B ribbon having an energy product of about 12 MGOe (Mishra, 1986).

Figure 3.10 (A) TEM image of the thin noncrystalline grain boundary layer between two Nd2Fe14B crystals. (B) Lorentz TEM image showing domain walls that are pinned at the grain boundary phase and extending around the two highlighted interaction domains (Mishra, 1986).

Figure 3.11 SEM micrographs of the fracture surface of an “optimum-quenched” melt-spun ribbon near the free surface (A), middle (B), and quench surface (C) of a Nd0.135(Fe0.945B0.055)0.865 ribbon melt spun at vs 5 19 m/s (Croat et al., 1984a).

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Figure 3.12 SEM of regions of the fracture surface of an “underquenched” melt-spun ribbon near the free surface (A), middle (B), and quench surface (C) of a Nd0.135(Fe0.945B0.055)0.865 melt spun at vs 5 14 m/s (Croat et al., 1984a).

the middle one/third, and the bottom one/third nearest the quench surface. Since these ribbons are approximately 25
40 μm thick, depending on quench rate, these data represent an extremely small fraction of the total surface area of the fracture surface. The SEM for the sample prepared at 19 m/s is shown in Fig. 3.11. This is the “optimum-quench” sample, with an energy product of 14 MGOe from Fig. 3.5. Although providing less microscopic detail than a TEM, these SEM data can provide valuable insight into the microstructure of these materials. This SEM clearly shows the microstructure to consist of a fairly uniform distribution of round-shaped grains of the Nd2Fe14B1 intermetallic phase with a size ranging from 20 to 80 nm. Fig. 3.12 shows a similar SEM for the “underquenched” ribbon prepared at 14 m/s from Fig. 3.5. This sample is also shown to consist of 20
80 nm crystallites in the bottom two-thirds of the ribbon. However, the upper-thirds of the ribbon was found to have a much coarser crystallite size, with average grain size of over 1000 nm. Note that the scale on the image of the fracture surface representing the top of ribbon is larger. It

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

is believed that the large grains found in these underquenched materials near the free surface significantly exceed the single domain size for these rapidly quenched materials and would, therefore, be expected to have low or even no coercivity. However, a material with a volume fractions of hard and soft magnetic materials should exhibit a step in the demagnetization curve between the first and second quadrant and this is not observed in the magnetization data. At this time, this behavior is not understood. It is possible that this drop or loss in magnetization may also be due to intergrain interaction, but is this case magnetostatic or dipole interaction, which can become important in materials with larger grains, as is the case with sintered Nd magnets and hotdeformed NdFeB magnets produced from melt-spun ribbon. These hot-deformed magnets are discussed in Chapter 6, Hot-deformed NdFeB permanent magnets. Fig. 3.13 shows a similar SEM for the sample prepared at 35 m/s; the demagnetization curve for this sample is also shown in Fig. 3.6 and exhibits almost zero

Figure 3.13 SEM of regions of the fracture surface of an “overquenched” melt-spun ribbon near the free surface (A), middle (B), and quench surface (C) of a Nd0.135(Fe0.945B0.055)0.865 melt spun at vs 5 35 m/s (Croat et al., 1984a).

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coercivity. The microstructure of this “highly overquenched” ribbon appears to be almost amorphous or glassy across its entire thickness, although some of the ribbon samples prepared at this quench rate did show miniscule crystallites near the upper surface. The SEM shown in Fig. 3.13 is one such example. The magnification used here is 147 kX, or roughly three times the magnification used in Figs. 3.11 and 3.12. Fig. 3.14 compares X-ray diffraction patterns of the Nd0.135(Fe0.945B0.055)0.865 ingot with several samples that were melt spun from this ingot at various quench rates. Included are the melt-spun ribbons prepared at 19 and 35 m/s and whose microstructures are shown in the SEM data in Figs. 3.11 and 3.13. As would be expected, these data show increased line broadening in the Bragg reflections as the quench rate is increased, indicating a progressive reduction in the average crystallite size and increased crystalline disorder. Despite this considerable line broadening, the one-to-one comparison with the Bragg reflections of the ingot clearly shows that the Nd2Fe14B1 intermetallic compound continues to be the primary phase in these highly quenched materials. Although the spectrum of the vs 5 35 m/s sample shows extensive line broadening indicative of a glassy or highly disordered material, there are still small peaks in the diffraction pattern corresponding to those of the Nd2Fe14B phase, and which are likely due to the small crystallites found near the free surface in the SEM data for these overquenched samples as shown in Fig. 3.13A.

Figure 3.14 Cu Kα X-ray diffraction patterns for the (A) Nd0.135(Fe0.945B0.055)0.865 ingot and selected samples prepared from this ingot (B
D) (Croat et al., 1984b).

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

3.2.3 The temperature-dependent properties of melt-spun NdFeB alloys Fig. 3.15 displays demagnetization curves of a melt-spun Nd0.13(Fe0.95B0.05)0.87 alloy as a function of temperature. These data were taken by first magnetizing the sample in a pulsed field of 4.5 T and measuring on a VSM with and operating field of 19 kOe. The data show the demagnetization curve at room temperature and at temperatures up to 446 K. Also shown are the room temperature properties after returning the sample to 446 K. The loss shown here is believed to be irreversible loss due to thermal demagnetization and would be recoverable by remagnetizing the sample. From these demagnetization curves, the coefficient of Br(α) and the coefficient of Hci(β) are calculated. The coefficient values obtained are α 5 20.12%/ C and β 5 20.40%/ C, which are the same as those for an A-Grade powder. These coefficients provide the % change in Hci and Br that can be expected for a 1 degree change in temperature. Fig. 2.26 shows a plot of saturation magnetization versus temperature for all of the R2Fe14B intermetallic compounds, which form, including Nd2Fe14B (Hirosawa et al., 1986). The change in Br seen in the demagnetization curves in Fig. 3.14 is directly related to the drop in the saturation 10

Nd0.13(Fe0.95B0.05)0.87 8 295 K after 450 K 295 K

M (kG)

6 340 K

4

372 K 2 446 K

–15

–10

–5 H (kOe)

0

5

Figure 3.15 Demagnetization curves of an optimum-quenched Nd0.13(Fe0.95B0.05)0.87 alloy as a function of temperature (Croat et al., 1984b).

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81

magnetization with temperature seen in these data. However, the change in coercivity with temperature is more complicated, since the observed coercivity is dependent on microstructure and is only a fraction of the 7.3 T anisotropy field found for the Nd2Fe14B compound (see Table 2.2). The demagnetization behavior and the values of α and β are used extensively by engineers who design motors and other devices, which use magnets, since the loss in magnetic properties with temperature must be accounted for in the design. This subject as well as the temperature aging properties of bonded Nd magnets produced from the melt-spun powder is discussed in more detail in Chapter 5, Production and properties of bonded Nd magnets, and some of the design criteria used to design magnetic circuits using these materials are included in Chapter 8, Major applications for rapidly solidified NdFeB permanent magnets.

3.2.4 The annealing behavior of melt-spun NdFeB alloys An important feature of melt-spun NdFeB alloys is the observation that the magnetic properties of the material can be varied by appropriate annealing. This characteristic of these fine-grained materials is still used today in the commercial production of NdFeB magnetic powder that is used for the production of bonded Nd magnets. As might be expected, annealing the material results in the growth of the microcrystals and the improvement in properties is believed due to the growth of the grains to a size more closely approximating the single domain size of the material. This allows melt-spun ribbon that has been overquenched (see Figs. 3.5 and 3.6) to be annealed so that its magnetic properties are at or close to the properties of optimum-quenched material. A significant advantage resulting from this annealing behavior is that it does broaden the range of substrate velocity (vs) or quench wheel speed over which production powder can be produced. However, as might also be expected, the magnetic properties of a material that has been underquenched or quenched at a rate less than the optimum rate cannot be improved by annealing, because improving this type of material would require reducing the average size of microcrystalline grains. Somewhat unexpected was the observation that the annealing behavior of these melt-spun NdFeB materials was found to have a significant dependence on quench rate and Fig. 3.16 shows a typical example of this behavior. This figure shows a Nd0.135(Fe0.945B0.055)0.865 alloy that has been melt spun at a substrate velocity of 20.5 m/s (see Fig. 3.6) to produce a sample that is slightly to moderately overquenched. This sample was then annealed by heating the samples to 950, 975, and 1000 K. A maximum energy product of 13.5 MGOe was obtained at an annealing temperature of 950 K, which is very close to the optimum 14 MGOe found by directly quenching this alloy from the melt. Annealing occurs very rapidly: these samples were annealed using a differential scanning calorimeter (DSC), which increased to the anneal temperature at a rate of 160 K/min and then immediately cooled to room temperature at the same rate. The sample size was only several grams, so the response time of the sample to temperature change was quite rapid. The annealing was carried out in an atmosphere of high-purity argon gas to protect

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

10 Nd0.135(Fe0.946B0.054)0.865 VS = 20.5 m/s 8

950 K 13.7 MG.Oe 975 K 13.3 MG.Oe

M (KG)

6

1000 K 12.9 MG.Oe 4 No anneal 9.0 MG.Oe 2

–15

–10

–5 H (KOe)

0

5

Figure 3.16 Second quadrant demagnetization curves for Nd0.135(Fe0.945B0.055)0.865 alloys quenched at a substrate velocity of 20.5 m/s and then annealed at 950, 975, and 1000 K (Croat 1988).

the samples from oxidation. However, for reasons not fully understood, annealing is only effective for material that has been melt spun at close to the optimum rate, as shown in the example in Fig. 3.16. Materials that have been moderated to highly overquenched can never be annealed back to even close to the properties obtained for the optimum direct-quenched material. An example of this is shown in Fig. 3.17, which displays second quadrant demagnetization curves for samples melt spun at 19, 20.5, and 35 m/s after annealing at 950 K. Again, these samples were part of the same group of material whose demagnetization curves are shown in Fig. 3.6. The 20.5 and 35 m/s alloys were annealed using a heating and cooling rate of 160 K/min. While the demagnetization curve of the 20.5 m/s sample comes close to that the optimum 19 m/s sample, the result for the 35 m/s sample is significantly lower. The behavior shown here is typical of all the melt-spun NdFeB alloys: the more overquenched the material, the less response, or improvement in magnetic properties was observed by annealing. A summary of this is provided in Fig. 3.18, which shows samples from Fig. 3.4, prepared at quench rates ranging from 14 to 40 m/s. All of the overquenched samples, i.e., those prepared at quench rates ranging from

10

Nd0.135(Fe0.935B0.065)0.865 8

19 m/s

M (kg)

6

20.5 m/s 4

35 m/s 2

–15

–10

–5 H (KOe)

0

5

Figure 3.17 Second quadrant demagnetization curves for Nd0.135(Fe0.945B0.055)0.865 alloys quenched at a substrate velocity vs of 20.5 and 30 m/s then heated to 950 K at a heating and cooling rate of 160 K/s (Croat, US Patent 4,802,931). 16 14

(BH)max MGOe

12 10 8 Direct quench 6

950 K anneal

4 2 0 10

15

20

25 Vs (m/s)

30

35

40

Figure 3.18 The energy product (BH)max of Nd0.135(Fe0.945B0.055)0.865 alloys quenched at substrate velocities ranging from 14 to 40 m/s and then annealed at 950 K at a heating and cooling rate of 160 K/s (Croat, US Patent 4,802,931).

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

19.5 to 40 m/s were annealed at 950 K in the DSC using a heating and cooling rate of 160 K/min. These data show that slightly or moderately overquenched material can be annealed to near optimum magnetic properties. However, the more pronounced the overquenched condition, the less effective the annealing in improving magnetic properties. Although annealing was found to improve the properties of all overquenched materials, it is almost always found that the highest magnetic properties were obtained by quenching directly from the melt (direct-quench) rather than annealing slight to moderate overquenched material. This observation is still found today in the commercial production of NdFeB magnetic powder. As would be expected, the magnetic properties of the annealed samples are related to the microstructure produced, as is readily apparent from SEM examination of these materials. Annealed sample of slight to moderately overquenched samples produced a relatively uniform microstructures similar to the SEM in Fig. 3.11, which is the image from the fracture surface of an optimum melt-spun material. However, annealing samples of highly overquenched melt-spun material produced microstructures very similar to the underquenched sample (vs 5 14 m/s) shown in Fig. 3.12, which have much larger grains in the top third of the ribbon than in the middle and lower third of the fractured ribbon. The more overquenched the sample, the larger the grain structure in the top portion of the annealed sample. The magnetization of the sample is also found to have significantly lower magnetization across the entire field range tested when compared to the optimum sample, similar to the behavior of the underquenched ribbon shown in Fig. 3.5. This behavior was unexpected and is still not fully understood. While annealing has been found to be an important feature of these melt-spun material, it is also observed that the annealing process must be carried out at as low a temperature as practical and as rapidly as possible in order to prevent overannealing and degradation of the magnetic properties. The results shown in Fig. 3.16 are fairly typical of this problem. While the optimum result, defined as the highest remanance and energy product, was obtained at 950 K, the samples annealed at 975 and 1000 K exhibit lower remanance and slightly higher coercivity. Why overannealing always results in both lower Br and higher Hci is not fully understood but it is suspected that the higher anneal temperatures result in some rearrangement or compositional change to the intergranular phase. The higher Hci is believed due to the growth of average grain size to one more approximating the optimum single domain size. However, there is no good explanation why the Br of the sample should drop. Since these samples would have been premagnetized at .4.5 T prior to the VSM test, this loss in Br should not be due to thermal demagnetization, which would be recoverable by demagnetizing the sample. One possibility is that the loss of Br is due to structural loss due to oxidation of the powdered sample even though the tests were carried out in an atmosphere of high-purity argon. Fig. 3.19 shows another example of overannealing. Displayed here are several Nd0.14(Fe0.95Fe0.05)0.86 alloys, which were melt spun at 27.5 and 30 m/s and then annealed on a DSC to 950 K. However, the samples were taken to and cooled from the maximum anneal temperature at two different rates (α) 5 40 and 160 K/min. As expected, the most overquenched sample (30 m/s) gave lower magnetic properties

The properties of melt-spun NdFeB alloys

85

10

Nd0.14(Fe0.95B0.05)0.85

VS = 27.5 m/s 8

α = 40 K/min, 11.4 MGOe α = 160 K/min, 12.6 MGOe

M (kG)

6 VS = 30 m/s α = 160 K/min, 11.1 MGOe

4

α = 40 K/min, 10.4 MGOe 2

–15

–10

–5 H (kOe)

0

5

Figure 3.19 Demagnetization curves for Nd0.14(Fe0.95B0.05)0.85 alloys melt spun at 27.5 and 30 m/s, which show the effect of both the quench condition and annealing time on the properties of the annealed ribbon (Croat, US Patent 4,802,931).

under both heating rates. However, the samples heated at the slower heating rate of 40 K/s samples both exhibit lower Br and higher Hci. Again, this change is believed due to growth of the Nd2Fe14B crystallites, so that the average grain size is slightly higher than the optimum single domain size. However, as mentioned, why such a change would result in a drop in Br is not readily understood. As is discussed in Chapter 4, Production of rapidly solidified NdFeB magnetic powder, annealing is an important part of the commercial production of melt spun NdFeB magnetic powder. However, the process must be carried out as rapidly as possible and at the lowest effective annealing temperature. Grain growth in these fine-grained materials is extremely fast and during the commercial production of melt-spun magnetic powder, the material is annealed at B700 C for only 1
2 minutes. A good technique for observing the crystallization and grain growth that occurs in these melt-spun materials during the annealing process is by use of a DSC. A representative example of the results obtained is shown in Fig. 3.20, which displays plots of apparent specific heat as a function of annealing temperature for the samples prepared at 19, 20.5, and 35 m/s, from Fig. 3.6. The 35 m/s sample shows a pronounced peak centered at around 950 K, which is an exotherm corresponding to

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

(A) Nd0.135(Fe0.946B0.054)0.865

Apparent specific heat

VS = 19 m/s

(B) VS = 20.5 m/s

(C) VS = 35 m/s

400

600

800

1000

T (K)

Figure 3.20 Differential scanning calorimeter of Nd0.135(Fe0.945B0.055)0.865 alloys melt spun at varying substrate velocities. The samples were heated at a rate of 160 K/min (Croat, 1988).

the crystallization of the highly disordered to amorphous structure into a finely crystalline microstructure. This peak has largely disappeared in the 20.5 m/s sample and is barely evident in the 19 m/s sample indicating a decreasing amount of disorder as the quench rate is decreased. The other prominent feature in these DSC tracings is the endotherm at B570 K (300 C), which is due to the magnetic ordering temperature of the Nd2Fe14B intermetallic phase. One curious feature is that this endotherm is largely absent in the most highly quenched alloy prepared at 35 m/s, which may suggest that this alloy is so disordered that it also losing long-range magnetic ordering. Although not shown here, the DSC tracing of the Nd0.135(Fe0.945B0.055)0.865 ingot from which these alloys were prepared was not noticeably different than that of the sample prepared at 19 m/s, indicating that both samples consists largely of a highly crystalline microstructure without significant disorder. Fig. 3.21 displays similar DSC data for melt-spun Nd0.15Fe0.85, Nd0.15(Fe0.95B0.05)0.85, and Nd0.15(Fe0.91B0.09)0.85 alloys at 15 and 30 m/s. For comparison, the data are overlaid on top of each other. The samples were also heated and cooled at a rate of 160 K/min. The results for the boron-free Nd0.15Fe0.85 sample is somewhat surprising in that the specific heat is almost flat

The properties of melt-spun NdFeB alloys

87

VS = 15 m/s

Nd0.15Fe0.85 VS = 30 m/s

Apparent specific heat

Nd0.15(Fe0.95B0.05)0.85 VS = 30 m/s VS = 15 m/s

Nd0.15(Fe0.91B0.09)0.85 VS = 30 m/s VS = 15 m/s

400

600

800

1000

T (K)

Figure 3.21 Differential scanning calorimeter tracings of Nd0.15Fe0.85, Nd0.15(Fe0.95B0.05)0.85, and Nd0.15(Fe0.91B0.09)0.85 alloys melt spun at 15 and 30 m/s. The samples were heated at a rate of 160 K/min (Croat, 1988; Croat, US Patent 4,802,931).

and shows no peak associated with crystallization, suggesting that the sample does not become disordered at zero B content, even at a quench rate of 30 m/s. This result is not explained since it would have been expected to contain some features resulting from the crystallization of the Nd2Fe17 intermetallic phase. This is in sharp contrast to the Nd0.15(Fe0.95B0.05)0.85 and Nd0.15(Fe0.91B0.09)0.85 samples, which shows behavior very similar to data in Fig. 3.19, with a pronounced crystallization peak at B950 K and the endotherm at B570 K corresponding to the magnetic ordering temperature. In the Nd0.15(Fe0.91B0.09)0.85 sample, there is also a thermal event at B940 K, which is believed to be an exotherm associated with the melting point of the Nd-rich intergranular phase. This is in good agreement with DSC investigations of textured hot-deformed NdFeB magnets (Brown et al., 2004; Kirchner et al., 2004) that found melting point of B660 C (B933 K) for the intergranular phase. What is surprising is that this exotherm is more strongly evident in the more boron-rich Nd0.15(Fe0.91B0.09)0.85 alloy and not observed in the Nd0.15(Fe0.95B0.05)0.85, which should also contain the Nd-rich intergranular phase.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

3.2.5 The magnetic properties of melt-spun R-Fe-B alloys As was pointed out in Chapter 2, The Nd2Fe14B intermetallic compound, the R2Fe14B crystal structure is reported to form for all of the rare earth or lanthanide elements except Eu. Despite the formation of this structure across virtually the entire lanthanide series, only the Nd and Pr compounds have properties suitable for commercially development of permanent magnets. Additions of all of the other rare earth elements result in a substantial reduction in magnetic properties in melt-spun materials. Fig. 3.22 shows demagnetization curves for a series of (Nd0.8R0.2)0.135(Fe0.0935B0.065)0.0865 alloys that were all melt-spun under optimum conditions. The results are, for the most part, in general agreement with those found for the cast and annealed R2Fe14B alloys, which were tabulated in Table 2.2 in Chapter 2. The room temperature magnetization of R2Fe14B intermetallic compounds for all R 5 La, Ce, Sm, Dy, and Tb is lower than that of Nd2Fe14B and Pr2Fe14B and consequently, all of these additives result in a drop in the Br and second quadrant magnetization of these alloys. The higher coercivity found for the Dy and Tb alloys is also expected, given the higher magnetocrystalline anisotropy values (Ha) found for these alloys relative to the Nd alloy. Both La and Ce additives result in a lowering of the magnetization and a reduction in coercivity, as also would be expected from their intrinsic properties. However, the comparatively low 10 (Nd0.8R0.2)0.135(Fe0.935B0.065)0.865

8

6 M (kG)

Tb

Dy 4 Nd Sm 2

La Pr Ce

–15

–10

–5

0

5

H (kOe)

Figure 3.22 Demagnetization curves for a series of (Nd0.8R0.2)0.135(Fe0.0935B0.065)0.0865 alloys with R 5 La, Ce, Pr, Nd, Sm, Dy, and Tb that were all melt spun under optimum conditions (Croat et al., 1984b).

The properties of melt-spun NdFeB alloys

89

coercivity of the melt-spun (Nd0.8Sm0.2)0.135(Fe0.0935B0.065)0.0865 alloy is unexpectedly given its high magnetocrystalline anisotropy, but this is probably due to the planar symmetry of the Sm2Fe14B compound. Fig. 3.23 shows results for a series of R0.135(Fe0.935B0.065)0.0865 alloys where R 5 La, Ce, Pr, Nd, Sm, Gd, Dy and Tb. Again, these alloys were all melt spun over a range of quench rate or substrate velocities, and the data shown represent the alloy exhibiting the highest energy product. The negligible coercivity found for La is expected since it is not believed to form in melt-spun alloys. Although La2Fe14B has been found to form, it was found to do so only after a long anneal at 850 C (Bolzoni et al., 1987). In any event, it would not be expected to develop significant coercivity since it has no 4f electrons and, therefore, would not be expected to have any magnetic anisotropy except that generated by the Fe sublattice. The very low first-quadrant magnetization and low coercivity of the Dy and Tb alloys are the most puzzling of all the data in Fig. 3.23 since both Dy2Fe14B and Tb2Fe14B have high Ha and Tc values (see Table 2.2) and should have made a significant contribution to the coercivity of the melt-spun samples. The only plausible explanation is that the anisotropies of these alloys are so high that the magnetic properties could not be developed even at the pulsed magnetic field of B4.5 T. Support for this is provided by Pinkerton (1986), which studied melt-spun DyFeB and TbFeB alloys and found that the TbFeB alloy could not be magnetically saturated, even in an applied field of 110 kOe. As was shown in Table 2.2, the anisotropy of the Dy2Fe14B and Tb2Fe14B intermetallic phases is B150 and B220 kOe, respectively.

10 R0.135(Fe0.935B0.065)0.865 8

La

6 M (kG)

Pr Sm

4 Nd

Ce

Gd 2 Tb. Dy

–15

–10

–5

0

5

H (kOe)

Figure 3.23 Demagnetization curves for a series of R0.135(Fe0.935B0.065)0.0865 alloys that were all melt spun under optimum conditions (Croat et al., 1984b).

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

The much different behavior of the Sm alloy, specifically the high magnetization combined with relatively low coercivity, is surmised to be due to the planar anisotropy of the Sm2Fe14B compound. The nonnegligible coercivity found for both Ce and Gd is also not unexpected since crystalline R2Fe14B compounds of both compounds are found to have fairly high Ha values of 26 and 24 kOe, respectively. This was shown in a study by Herbst et al. (2012), which prepared melt-spun CeFeB alloys and found optimum properties of Br 5 4.9 kG, Hci 5 6.2 kOe and (BH)max 5 4.1 kOe. The low Br is expected since the Ce2Fe14B intermetallic phase has a much lower saturation value than Nd2Fe14B. The coercivity could be due to a substantial anisotropy generated by the Fe sublattice in these compounds. Although it is often taught that the orbital momentum in transition metals is quenched by the strong crystal field, this is obviously not always the case, as evidenced by the high magnetocrystalline anisotropy found in intermetallic alloys like YCo5 and Y2Co17. These alloys should have no magnetocrystalline anisotropy, since Y has no unpaired 4f electrons (J 5 0). At one point, Ce was also believed to have no 4f electron and, therefore, would not contribute to the magnetocrystalline anisotropy and coercivity of magnetic alloys. However, as discussed previously, the coercivity observed in the Ce0.135(Fe0.935B0.065)0.0865 compound may result from the Ce ion carrying a mixed 13/ 1 4 valence, as reported by Ro¨hler (1987) and Capehart et al. (1993), which would result in a nonzero orbital angular momentum contributing to the magnetocrystalline anisotropy. Recently, Pathak et al. (2015) reported a substantial increase in Hci in both melt-spun and sintered alloys for certain compositions, for example, (Nd0.8Ce0.2)14Fe12Co2B, and proposed that Ce may be suitable as a replacement for the much more expensive and scarce Dy currently used to enhance coercivity in sintered Nd magnets.

3.2.6 The effect of transition metal (Co, Ni, Mn, Cr, Cu) additives The effect of replacing the Fe in melt-spun NdFeB alloys with one of several common TMs is shown in Figs. 3.24 and 3.25. Fig. 3.24 displays demagnetization behavior a series of melt-spun Nd0.135(Fe0.841TM0.094B0.065)0.865 alloys, where the amount of TM additive represents 10% of the total TM in the composition. Again, the data shown here all represent the sample having optimum magnetic properties when melt spun over a range of substrate velocities (vs). Additives include the common ferromagnetic TMs Co and Ni as well as the Mn and Cr, which are ferrimagnetic. Copper was also tested because it is a common TM additive in certain types of rare earth-cobalt magnets. Of these elements, only Co was found to have no serious deleterious effect on the magnetic properties of melt-spun Nd-Fe-B. All of the other elemental additives effected a significant drop in both remanance and coercivity. These results are not unexpected since only Co has a moment equivalent to Fe and a Curie temperature higher than Fe. Although Ni is also ferromagnetic, its magnetic moment and Curie temperature are significantly lower than Fe and would be expected to reduce both the remanance and coercivity of melt-spun NdFeB alloys. Likewise, Mn and Cr are ferrimagnetic with quite low-ordering temperatures and, barring some unexpected structural or magnetic phenomena, would also be

The properties of melt-spun NdFeB alloys

91

10 Nd0.135(Fe0.841TM0.094B0.065)0.865 8

6

Fe Ni

4 Co

M (kG)

Cr

2

Mn Cu –15

–10

–5

0

5

H (kOe)

Figure 3.24 Demagnetization curves for Nd0.135(Fe0.841TM0.094B0.065)0.865 alloys that were all melt spun under optimum conditions (Croat, US Patent 4,802,931). 10

Nd0.135(Fe0.748T0.187B0.065)0.865 8

6 M (kG)

Fe

4

Nl Co

Cr

2 Mn

–15

–10

–5

0

5

H (kOe)

Figure 3.25 Demagnetization curves for Nd0.135(Fe0.748TM0.187B0.065)0.865 alloys that were all melt spun under optimum conditions (Croat, US Patent 4,802,931).

expected to effect a reduction in magnetic properties. Fig. 3.25 shows demagnetization curves for Nd0.135(Fe0.748TM0.187B0.065)0.865 alloys, where 20% of the Fe has been replaced by the other TM elements. Again, there is no reduction in properties with higher levels of Co. However, these levels of Ni, Mn, and Cr result in a further

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

dramatic drop in both remanance and coercivity. Since the Nd2Fe14B structure is not believed to form for Ni, Mn, and Cr, this drop is most likely due to the formation of additional compounds and the dilution of the Nd2Fe14B phase in these alloys. However, the microstructure or these melt-spun alloys or the nature of any additional compounds was not investigated. The effect of even higher levels of Co on the properties of melt-spun NdFeB was examined by Fuerst and Herbst (1988) who studies melt-spun Nd0.135(CoxFe12x)0.95B0.05 alloys for x 5 0, 0.05, 0.10, 0.15, 0.20, 0.35, 0.50, 0.75, and 1.0. The samples were all melt spun under optimum conditions. For x 5 0.5 to 0.35, the demagnetization curves were similar to that of the cobalt free sample (x 5 0). However, for x . 0.5, the remanance, energy product, and especially the intrinsic coercivity deteriorated rapidly. X-ray diffraction data showed that for x . 0.05, the Nd2Fe14B structure is suppressed and the sample showed increasing greater amounts of the Nd2(Co,Fe)17 compound. Similar behavior was also found for Pr2(Co,Fe)14B alloys (Fuerst, 1990). Cobalt is a technically important additive to melt-spun NdFeB materials because, as shown in Fig. 3.24, it is the only TM that does not reduce the remanance, at least in relatively low amounts. Addition of Co also increases the Curie temperature, resulting in an increase in Tc of approximately 10 C for each weight % of Co added. Cobalt was originally thought to improve the aging behavior of bonded Nd magnets, i.e., reduce the loss in magnetic properties when the magnets are aged for a period of time at a given temperature. In recent years, however, this has been found not to be the case and Co has generally been removed from commercial grades of NdFeB magnetic powder because of its comparatively high cost. Because of the rapid growth in the market for Nd magnets in the early 1990s, driven by the growth in demand for hard disk drives (HDD) in personal computers, there was a temporary but serious shortage of Nd metal. To alleviate this problem, Pr was added to make up part of the composition. However, this almost immediately resulted in customer problems, who found that their spray coatings, typically phenolic epoxies, began to fail if the Pr content exceeded a certain percentage of the total rare earth content. The surface of the magnet would form a thin oxide layer, causing the coatings to delaminate. The exact cause of this problem was never determined, but it was believed at the time to be somehow related to the mixed 13/ 1 2 valence in Pr2O6 oxide compared to the 13 valance in Pr2O3. This problem was solved by the addition of 5.0 wt% Co to the composition, which did solve the coating problem. However, with improved coating technology and techniques, the use of Co for this purpose also became unnecessary. However, the Co was never completely removed until more recently.

3.2.7 The effect of small amounts of elemental additives There is a significant body of literature on the effects of small additions of TM and post-TM additives on the magnetic properties of melt-spun NdFeB alloys. However, the results are somewhat mixed and in some case contradictory. In general, there have been two types of additives that have been investigated, which include so-called grain boundary modifiers and grain growth inhibitors.

The properties of melt-spun NdFeB alloys

93

3.2.7.1 Grain boundary modifiers These are additives which tend to alloy with the Nd-rich intergranular phase and change the properties by modifying this phase. These additives are typically lower melting point elements such as Ga, Al, Cu, and Sn and have smaller ionic radii than Nd and, therefore, alloy less easily with the Nd2Fe14B intermetallic phase. Most studies have found that small additions of these elements can lead to an increase in the coercivity, which is believed to result by changing the properties of this important intergranular phase. For example, Herbst et al., (1991) studied the effect of 0.5 at% Cu on melt-spun Nd-Fe-B ribbons and found a 30% increase in the intrinsic coercivity but, as expected, a drop in Br. TEM analyses showed the Cu preferentially segregated in the intergranular phase. Among the grain-modifying elements, Ga has probably been investigated the most extensively and best exemplifies the role that these grain boundary modifiers can have on the properties of these melt-spun materials. Panchanathan and Croat (1989) reported a large increase in Hci with addition of 0.6 wt% Ga in NdFeCoB alloys without significant loss in Br. TEM analysis did find that the Ga was concentrated in the grain boundary phase and it was postulated at the time that the Ga improved the coercivity by changing the strength of the domain wall pinning sites. Additions of Ga and Ge were investigated by Gholamipour et al. (2006) using atomic probe analysis. They also found an increase in Hci without loss of Br and that the Ga was deposited uniformly in the Nd-rich intergranular phase. Germanium, however, was found to form Ge-Nd precipitates in the grain boundaries, which lead to a drop in coercivity. This study postulated that Ga resulted in a more uniform intergranular phase, which increased the isolation between grain, resulting in less exchange interaction between the grains and an increase in coercivity. In contrast, Ge was found to produce a less uniform intergranular layer which, resulted in less intergrain isolation which, in turn, resulted in an increase in exchange interaction and a corresponding drop in coercivity. The effect of Ga additions or combinations of Ga and Co have also been studied in hot-deformed NdFeB magnets produced from melt-spun powder. Mishra et al. (1993) and Fuerst and Brewer (1993) reported that combinations of Ga and Co were effective in increasing both the coercivity and remanance of hot-deformed magnets. Gallium has also been studied by Tokunaga et al. (1989) who reported that Ga was the best of a number of additives investigated for improving the coercivity of hot-deformed magnets. Again, the Ga is believed to influence the magnetic properties by segregating in the Nd-rich intergranular phase and improving hot workability during the hot deformation process. Kirchner et al. (2004) studied the intergranular phase in hot-deformed magnets using highresolution EDX and found that the Ga was almost completely segregated in the intergranular phase. Their DSC studies showed that the addition of 0.6 wt% Ga reduced the melting point of the intergranular phase by B20 C. These hotdeformed materials NdFeB magnets are discussed in detail in Chapter 6, Hotdeformed NdFeB permanent magnets.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

3.2.7.2 Grain growth inhibitors The second type of additives that have been investigated are often referred to as grain growth inhibitors and generally include the refractory metal elements such as Nb, Mo, V, W, Zr, and Ti, although other elements such as Co and Cr are also thought to also act as grain growth inhibitors, since they are thought to result in the development of a finer and more uniform microstructure by preventing grain growth during the melt-spinning process or during subsequent annealing of the powder. It was originally hypothesized that these elements expressed themselves by alloying with the Nd2Fe14B intermetallic phase rather than segregating in the grain boundaries. However, this does not appear to be the case and many of these elements also segregate in the intergranular phase. Additions of these elements or combinations have been studied by Kohmoto et al. (1987), Pollard et al. (1988), Wecker and Schultz (1990), Kim et al. (1991), Harland and Davies (2000), Chen et al. (2004), Bao et al. (2009), and Zhang et al. (2012). As in all investigations where nonmagnetic elements are added to melt-spun NdFeB, the addition of all but small amounts these elements usually results in a reduction in remanance and energy product. Therefore, most of the studies are concerned with small amounts (,0.5 at%) of these additives. As a general observation, it has been found that small amounts of the refractory metals, for example, Nb, are found to enhance both the Br and Hci of melt-spun alloys and to improve the thermal aging properties. These improved properties have been attributed to the development of a finer grain texture and a more uniform average grain size. However, the Nb was found as a second phase in the grain boundary phase rather than alloying directly with the Nd2Fe14B compound, as originally theorized. Several of the studies (Kim et al., 1991) found that additions of Nb, in particular, resulted in an increases in both Br and Hci. Similar results were reported by Zhang et al. (2012) and found that small additions of Nd increased the coercivity and also reduced irreversible losses. Chen et al. (2004) also reported higher Br and Hci in melt-spun Nd12Fe80.5B6Nb1.5 ribbons and found, from TEM analysis, a substantially finer average grain size than in melt-spun Nd12Fe82B6 ribbon under the same conditions. However, they reported that the melt-spinning and subsequent annealing process expels the Nb from the Nd2Fe14B phase into the grain boundaries, where it formed a Nb-rich phase, which inhibited grain growth and results in a more uniform grain structure. As discussed earlier, it is believed that the highest magnetic properties, usually Br, energy product (BH)max, and second quadrant knee shape result when the microstructure is uniform across the thickness of the melt-spun ribbon. As a result of these types of studies, small amounts of Nb (0.2
0.4 wt%) are now added to the chemistry of several commercial high-performance magnetic powders.

3.2.8 Melt-spun Nd-Fe-Me (Me 5 Si, C, Al, Ge, and P) alloys Croat et al. (1984a,b) reported that boron appeared to be unique among the metalloid elements in its ability to form the Nd2Fe14B phase. A series of Nd0.135(Fe0.935Me0.065)0.865 alloys with Me 5 Si, C, Al, Ge, and P were prepared by arc melting and then melt spun over a range of substrate velocities (vs), similar to the data shown in Figs. 3.5 and 3.6. This composition was chosen because it is

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95

The magnetic properties of melt-spun Nd0.135(Fe0.935Me0.065)0.865 alloys with Me 5 Si, C, Al, Ge, and P

Table 3.1

close to that exhibiting optimum results for NdFeB alloys. The results are summarized in Table 3.1. The highest coercivity found for the Al, Si, P, and Ge samples was ,100 Oe. Only the sample containing C was found to develop an appreciable coercivity of 750 Oe. X-ray diffraction analysis of the samples showed all of the alloys to consist primarily of the Nd2Fe17 intermetallic phase, so it is not surprising that the alloys showed no appreciable magnetic hardening. To this date, there have been no reports that the Nd2Fe14B structure forms for Me 5 Si, Al, Ge, and P. However, a number of subsequent studies (Buschow et al., 1988; Coehoorn et al., 1989; Lui et al., 1990; Eisses et al., 1991; Yang et al., 2000) have reported that Nd2Fe14C can be prepared by annealing either cast ingot, melt-spun ribbon, or mechanically alloyed precursors. A Curie temperature of 532 K and a saturation magnetization of 14.1 kG have been reported for this phase. However, the carbide phase is reported to form by a very sluggish peritectic reaction from the R2Fe17C2 and RFeC phases and only by annealing below 700 C. Above this temperature, a solid state transformation occurs in which the Nd2Fe14C phase transforms into the R2Fe172xCx compound. This is because the solid state transformation temperature is quite low for the light rare earths but increases with increasing atomic weight across the lanthanide series. As a consequence, the formation of the R2Fe14C phase has been found to be most difficult for the light rare earth elements and easier for the heavy rare earth elements. As would be expected, the formation of the phase is stabilized by the addition of boron or one of the heavy rare earths.

3.3

The magnetization process in isotropic melt-spun NdFeB

3.3.1 Magnetic domains and domain walls An understanding of the magnetization and demagnetization process, as well as the applicable coercivity mechanism, in melt-spun NdFeB requires an understanding of domain walls and domain wall motion (Parker and Studders; 1962; Morrish, 1965; Cullity, 1972; Campbell, 1994). All ferromagnetic materials are composed of small regions, called magnetic domains, in which there is local magnetic saturation

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

(Msat), that is, regions in which the magnetic moments or magnetic dipoles are aligned parallel. However, the direction of magnetization in neighboring domains is not necessarily parallel and in the absence of an applied magnetic field, the domains arrange themselves such that the average magnetization is equal to zero. Magnetic domains are separated by domain walls, which are interfaces between regions in which the magnetization has different directions. Domains are small, but much larger than atomic distances. A depiction of a domain wall where the magnetization varies by 180 degrees is shown in Fig. 3.26. The figure depicts how the magnetic dipoles become canted and rotate relative to the c-axis within the wall. Magnetic domains form because a large ferromagnetic single crystal that is uniformly magnetized will have a north pole at one end and a south pole at the other. These magnetic poles generate an external field outside the magnet as shown in C-axis

Rotation axis

Domain wall

Figure 3.26 Depiction of the change in the direction of the magnetic dipole across a domain boundary. The change in direction does not occur abruptly but extends through the thickness of the boundary.

N

N

N

S

N

S

N

S

S

S

N

S

N

S

Single crystal with single domain

Single crystal with multiple domains

Figure 3.27 The multiple magnetic domain structure which forms to reduce the magnetostatic energy in ferromagnetic materials.

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Fig. 3.27. However, the magnetic material also forms a second magnetic field within the magnet but in the opposite direction. This field, which tries to demagnetize the material, is called the demagnetization field and the energy from this field is called the magentostatic energy. This magentostatic energy is also equal to the amount of work required for the magnetic poles to exist counter to the internal magnetic field at both ends of the magnetic body. A strongly magnetic body will form magnetic domains to minimize this magnetostatic energy as shown in Fig. 3.27. Formation of the domain wall brings the north (N) and south (S) charges close together, thus decreasing the spatial extent of the magnetic field. However, this subdivision into more and more domains cannot continue indefinitely because the domain walls also require energy to be produced and maintained. Domain walls have a finite width that is determined principally by a tradeoff between exchange and magnetocrystalline or anisotropy energy. The exchange energy favors dipoles or spins that are parallel and this energy can be kept small if the 180 degree rotation takes place gradually, over many atomic units. However, the magnetic dipoles within the wall are no longer aligned along an easy axis of magnetization, which in the Nd2Fe14B structure, is the c-axis direction. This leads to an increase in the magnetocrystalline energy, which is higher if the transition occurs gradually. Thus, the exchange energy tends to make the wall as wide as possible, whereas the anisotropy energy tends to make the wall as thin as possible. The balance between these two energies results in the domain walls having a certain thickness and a specific domain wall size and also results in an optimum single domain size. As the grain size decreases, a critical size will be reached where the grain can no longer accommodate a wall. Below this critical size, the grain contains a single domain. As the crystallite size increases, it becomes energetically favorable for the grains to form multiple domains. The critical single-domain particle size is the largest crystallite size in which the energy cost to form a domain wall is higher than the reduction in magnetostatic energy. Because the magnetocrystalline anisotropy is different for various ferromagnetic materials, this optimum domain size will also vary. The TEM shown in Fig. 3.9 is from a melt-spun sample Nd0.135Fe0.810B0.055 and an energy product of B12 MGOe with no demagnetization correction. The average diameter of the grains in this sample is B30 nm, which should be close to the single domain size for these materials. However, critical domain sizes have been calculated from several different studies and all have shown a single domain size much larger than this observed value. From TEM studies, Mishra (1986) reported a domain wall width (δ) averaging 4 nm, from which he calculated a domain wall energy σw 5 17 mJ/m2 and a critical single domain size of 150 nm using the anisotropy constant K 5 4.5 mJ/m3 for Nd2Fe14B. This is very close to the value of B160 nm reported by Fidler and Yang (1985). Livingston (1985) and Sagawa et al. (1985) have both reported a single domain size of 300 nm. All of these calculated values for the single domain size are at least 10 times higher than the 30 nm grain size found in melt-spun ribbons having optimum magnetic properties.

3.3.2 The magnetization and demagnetization process The TEM in Fig. 3.9 shows the microstructure of melt-spun NdFeB with optimum magnetic properties to consist of extremely fine-grained, round shaped crystallites

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of the Nd2Fe14B intermetallic phase surrounded by a thin layer of a Nd-rich intergranular phase. The magnetization process that is believed applicable for these finegrained materials is shown in Fig. 3.28. In this hypothetical model, the individual Nd2Fe14B crystallites are shown with arrows representing the easy c-axis of the tetragonal crystal structure. Each grain is surrounded by a thin layer of the Nd-rich intergranular phase that surrounds each grain. In a thermally demagnetized sample, the magnetic moments would be randomly oriented, so that for H 5 0, the magnetization M is also equal to zero. This is the situation shown in Fig. 3.28A. TEM studies (Chen, 1985; Mishra, 1986; Hadjipanayis and Gong, 1987, 1988b) have reported that the individual crystallites in thermally demagnetized samples consist predominately of single magnetic domains. Since M 5 0, exactly half of the moments are directed toward the field direction and half are directed opposite the field direction. As the field is increased to the right as shown in the full hysteresis curve, the magnetization increased along the dotted initial magnetization curve. During this phase of the magnetization process, the moments that lie in grains with their c-axis most closely aligned with the magnetic field but whose moments are opposed to the magnetic field direction, flip or reverse into the c-axis direction most closely aligned with the magnetic field. The mechanism by which these grains reverse their magnetization involves either the unpinning of an existing domain wall or the nucleation of a new domain wall from the intergranular phase. Both processes would require energy, so there is a coercive force that developed. This subject is discussed in (A)

(B)

H– 0

(C)

H> 0 Mf

(D)

H = Ha (Gauss) +0

Oc

–H

Hci 0

Hc

+H (Oersted) H1

Or

H = Hci

–0

Figure 3.28 The magnetization process in melt-spun NdFeB alloys showing the direction of the magnetic moments in the individual Nd2Fe14B crystallites as the magnetic field is first applied in a positive direction (1H) and then reversed into a negative direction (2H). The separate panels show the structure of the magnetic moments at (A) H 5 0, (B) H in a positive direction equal to Hci value, (C) H at a high field equal to Ha and (D) H in the negative direction and5 Hci.

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more detail in Section 3.4.3 but there is still no consensus opinion on the applicable coercivity mechanism in these materials. As the magnetic field continues to increase, the moments in the grains continue to reverse until all are aligned along the c-axis that is most closely aligned with the field direction. This is the situation shown in Fig. 3.28B, for H . 0. For even higher field levels, the magnetization continues to increase as the moments begin to rotate out of their c-axis direction and toward the direction of the magnetic field. This continued alignment is very difficult, because this rotation is opposed by the magnetocrystalline anisotropy (Ha), which is B7.3 T for the Nd2Fe14B intermetallic compound. This is the situation shown in Fig. 3.28C, where the applied field is equal to the magnetocrystalline anisotropy of the Nd2Fe14B compound. At this high field, the magnetization is equal to the saturation magnetization of the melt-spun ribbon materials and indicated by the horizontal dotted line in Fig. 3.28. In a single phase, crystalline alloy, this would be B16 kG. However, this value would not be as high in the melt-spun ribbons because the primary Nd2Fe14B phase diluted by the Nd-rich intergranular phase, which has until recently been assumed to a paramagnetic phase, which would not contribute to the magnetic moment. However, as mentioned earlier, a number of studies have reported that this phase may be ferromagnetic and, if so, it would likely be the same in these fine-grained melt-spun materials and would also make a contribution to the magnetization of the sample. It is noted, however, that the question of whether the intergranular phase is paramagnetic or ferromagnetic still a subject of debate at this time. As the magnetic field is reversed, but still positive, the magnetization falls as the magnetic moments relax back toward their natural position along the easy c-axis of the individual Nd2Fe14B crystallites. At H 5 0, the magnetization is equal to the remanance (Br) and the magnetic moments should once again be aligned along the crystallographic axis direction closest to the direction of the magnetic field. Thus, the magnetic alignment would appear to be similar once again to that depicted in Fig. 3.28B. However, this is not believed to be technically correct. Because of exchange interaction between the magnetic dipoles in neighboring grains, some of the moments in the grain boundary regions remain partially aligned with the moment in the neighboring grain, resulting in a Br/Msat . 0.5, the maximum allowed for an assembly of noninteracting particles. This is discussed in more detail in Section 3.4.3. When the magnetic field is reversed (2H), the direction of the moments of the individual crystallites begins to reverse. However, there is a resistance to this process because reversal occurs by the formation of a domain wall, which then sweeps across the grain to reverse the magnetization. There is an energy associated with the formation of the domain wall, thus resulting in significant coercive force. When exactly half of the moments have flipped or have reversed their direction, the net magnetization is again equal to zero. This is the point on the horizontal axis, which defines the intrinsic coercivity (Hci) of the material. At this point, the situation would be technically the same as the starting thermally demagnetized shown in Fig. 3.28A, where M is again equal to zero. However, this condition was reached by DC field demagnetization instead of thermal demagnetization. Therefore, if the sample were removed from the applied reverse field, there is likely to be some recoil or domain wall relaxation

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

and M ¼ 6 0, but would have a small value. The only way that the sample could be completely demagnetized would be to apply a reverse field slightly in excess of Hci, until the sample recoiled back to where M 5 0. Of course, as the reverse field continues to increase, the magnetic moments once again rotate toward the field direction and a situation similar to Fig. 3.28C is again produced at 2 Ha, but with the moments pointed in the opposite direction. It is clear from the shape of the second quadrant demagnetization curve for these isotropic materials that the magnetization of the various grains does not reverse at the same time, but over a wide range of field levels. As might be expected, the reverse field for which 2 H 5 Hci is not nearly as large as the magnetocrystalline anisotropy (Ha) of the Nd2Fe14B phase, which is B7.3 T. The coercive level in all permanent magnet materials, including melt-spun NdFeB, is always much lower than Ha and occurs by the formation and motion of domain walls, rather than coherent rotation of the magnetic dipoles. This is because the coercive force is not an intrinsic property of the material, but rather results from a complex relationship between the formation of domain walls and the microstructure of the material. In most of rare earth permanent magnet materials, the observed coercivity is less than 20% of the measured magnetocrystalline anisotropy, a problem referred to as Brown’s paradox. An example of experimental data showing the initial magnetization and demagnetization of melt-spun ribbon from Pinkerton and Van Wingerden (1986) is shown in Fig. 3.29. These data were obtained from a standard VSM with a maximum field level of 18 kOe. Each partial hysteresis curve was taken on separate sample taken 12

4 π M (kG) Nd-Fe-B ribbon

8

Hm (kOe) 110 17.5 16

4 14 12 10 8

–20

–15

–10

–5

6 4

0

5

10

15

20

H (kOe)

Fig. 3.29 Initial magnetization and demagnetization curves for melt-spun NdFeB ribbon. The number on each curve from 4 to 17.5 represents the maximum field that the sample was taken to on the VSM during the test. The sample labeled 110 was premagnetized at a field of 110 kOe prior to the demagnetization test on the VSM (Pinkerton and Van Wingerden, 1986).

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101

from a single batch of crushed and homogenized melt-spun ribbon. The sample labeled 110 was magnetized at this field level in a superconducting solenoid prior to being measured on the VSM. However, the remaining samples were not premagnetized in a high field as was the case for most of the data discussed or shown in this chapter. The numbers on the curves represent the maximum field level that the sample was taken on the VSM during each separate test. These data show the initial magnetization behavior that is characteristic of melt-spun NdFeB and which is marked by a low initial permeability, which increases dramatically as the magnetizing field becomes comparable to the intrinsic coercivity of the sample (B12 kOe), and then again drops for higher applied fields. This initial magnetization curve is, of course, the same as that in the hypothetical hysteresis curve in Fig. 3.28. From this behavior, it was concluded in this study that the coercivity in these melt-spun ribbons was controlled by a domain wall pinning mechanism.

3.3.3 Intergrain interaction in melt-spun NdFeB alloys There are two types on interactions that can occur between grains in NdFeB magnetic materials. These include long-range magnetostatic (dipolar) interaction and exchange interaction between neighboring grains. These two effects are shown in the drawing in Fig. 3.30 (Fidler and Schrefl, 2000). The magnetostatic interaction arises from the net magnetic moment and is thought to be large in magnets with larger grains, such as hot-deformed or sintered Nd magnets, but less important in fine-grained material such as isotropic melt-spun NdFeB. Exchange interaction between neighboring grains can cause the magnetic dipoles to deviate from the easy c-axis of one grain toward the easy axis of the neighboring grain. A simple depiction of these phenomena is shown in Fig. 3.31, which shows two Nd2Fe14B grain, both of which are surrounded by a thin layer of the Nd-rich intergranular phase. The large gray arrows represent the crystallographic c-axis of the grains and the small black arrows depict the individual magnetic moments. The moments within a grain are parallel except in the neighborhood of the grain boundaries where the moments on either side of the boundary are believed to rotate into alignment due to exchange interaction between the two grains. As might be expected, this interaction becomes much stronger as the grain size is reduced and as the amount or thickness of the Nd-rich intergranular phase becomes thinner. If the grain size Exchange interaction

J

J

Dipolar interaction

Figure 3.30 Long-range dipole (magnetostatic) interaction and short-range exchange interaction between two neighboring Nd2Fe14B grains (Fidler and Schrefl, 2000).

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

Figure 3.31 A representation of the exchange interaction that is believed exist between moments in adjacent grains in finely crystalline melt-spun NdFeB alloys.

becomes small enough (,20 nm), the exchange interaction can lead to complete alignment of the moments in adjacent grains, particularly if there is less of the intergranular phase to isolate one grain from another. However, as the grain size increases, the short-range nature of the exchange interaction is reduced and only the region of the grains next to the grain boundaries is aligned as shown in the rendering in Fig. 3.31. These phenomena are quantum mechanical in origin and result from a coupling of the electron spins, resulting in a reduction in the total magnetic energy as the moments become parallel. It was initially thought that significant amount of exchange interaction would not occur in melt-spun alloys containing significant amounts of the Nd-rich intergranular phase, because the latter would magnetically isolate the individual grains from each other. However, various researchers (Mishra, 1986; Pinkerton, 1987; McCallum et al., 1987; Clemente et al., 1988; Hadjipanayis and Gong, 1988b; Hadjipanayis and Kim, 1988a; Matsumoto et al., 1988) have studied this problem and all concluded that exchange interaction does occur in these standard isotropic melt-spun NdFeB materials. Exchange interaction was first postulated by Stoner and Wohlfarth (1948) who derived various relationships for uniform assemblies of interacting and noninteracting, single-domain particles. One of the relationships observed by Stoner and Wohlfarth is Bd(H) 5 Br(N) 2 2Br(H), where Br(H) is the remanance after magnetizing a thermally demagnetized sample in a magnetic field of relatively low strength, Bd(H) is the remanance after magnetizing the same sample in a very large field, comparable to its magnetocrystalline anisotropy, and Br(N) is the remanance of a fully magnetized sample and equal to Msat. Failure to comply with this relationship is believed to be an indication that some kind of interaction does occur between the individual particles. Various researchers (Pinkerton, 1987,1988; Hadjipanayis and Gong, 1988b) have applied this relationship to the hysteresis behavior of melt-spun NdFeB and found quite significant departure from this relationship. An example from Pinkerton (1987) is shown in Fig. 3.32, where

The properties of melt-spun NdFeB alloys

103

1.0

Bd/Br (∞)

0.5

0

–0.5

–1.0

0

0.2

0.4 0.6 Br/Br (∞)

0.8

Figure 3.32 Demagnetizing remanance Bd (H) versus magnetizing remanance Br (H) for melt-spun NdFeB. An assembly of noninteracting particles would fall on the dashed line (Pinkerton, 1987).

the dashed line is calculated from the relationship shown earlier and would indicate no exchange interaction. The significant deviation seen here suggests that there is considerable exchange interaction between the individual Nd2Fe14B crystallites. In addition, there has always been a qualitative argument for some exchange interaction in these materials. This is because Msat for the Nd2Fe14B phase is 16 kOe, and the reduced remanance Br/Msat for these melt-spun materials is often observed to exceed 8 kG, the maximum expected for an isotropic material consisting of noninteracting particles. Moreover, these materials consist of two phases, and the Nd-rich intergranular phase, which is generally believed to be paramagnetic, would contribute only weakly to the total magnetization. Finally, there has been no demagnetization correction applied to any of magnetic properties that have been displayed in the various figures in this chapter. When a reasonably demagnetization correction is applied, Br values of 9 kG or higher are easily obtained. If we assume that just 5% of the total volume consists of the Nd-rich intergranular phase, then the remanance provided by the Nd2Fe14B phase would be B9.5 kG, leading to a Br/Msat of almost 0.6. This would imply that exchange interaction enhances the remanance by as much as 15%.

3.3.4 The coercivity mechanism in melt-spun NdFeB alloys From Lorentz TEM studies, both Mishra (1986) and Hadjipanayis and Gong (1987) concluded that the individual grains in these isotopic materials consist of single

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

1.2

1.0

Br/Br (∞)

0.8

0.6

0.4

Ribbons Hot pressed Die upset Sintered

0.2 0 0

0.5

Hm/Hci (∞)

1.0

1.5

Figure 3.33 Room temperature intrinsic coercivity of various NdFeB magnets versus magnetizing field. The quantities are both normalized to the coercivity of a fully magnetized sample, expressed here as Hci(N) (Pinkerton and Van Wingerdon, 1986).

magnetic domains, or at most, small cluster of grains comprising a single interaction domain. In these fine-grained materials, there would be no domains walls within the grains, since the grains are much smaller than the single domain size and the formation of a domain would be energetically unfavorable. From the Lorentz TEM images shown in Fig. 3.10B, it would appear that the magnetization process of a thermally demagnetized material occurs by the pinning of domain walls by the Nd-rich intergranular phase, which surround each grain. If this is the case, then the initial magnetization curve would represent the unpinning of domain walls, which move across the grains until they are again pinned at the opposing grain boundary phase. Magnetization is believed to occur on a grain by grain basis, or a most, a few grains at a time until the material is completely magnetized in the direction of the applied field. As was mentioned earlier, a pining model was used by Pinkerton and Van Wingerden (1986) to explain the initial magnetization data in Fig. 3.29. A domain wall pinning model was given added support by comparing the development of the initial magnetization in melt-spun ribbon to that of other NdFeB magnetic materials, specifically hot-deformed melt-spun ribbon and sintered Nd. This comparison from Pinkerton and Van Wingerden (1986) is shown in Fig. 3.33. To show a better comparison, the Br of the sample has been normalized to that of a fully magnetized samples Br(N) and the applied field Hm normalized to the intrinsic coercivity Hci. It is observed that the magnetization develops much more slowly in melt-spun NdFeB ribbon than in either hot-deformed NdFeB or in sintered Nd magnets, and that the maximum value does not developed until HcHci. Sintered Nd magnets show a much more rapid initial increase in the coercivity followed by a much slower increase as the magnet approaches saturation. The phenomenological

The properties of melt-spun NdFeB alloys

105

explanation for this is that the individual grains in a thermally demagnetized sintered Nd magnet, which measure 2
5 μm in diameter, consist of multiple domains, which move easily in an applied magnetic field, resulting in a rapid increase in susceptibility. This issue is also discussed in Chapter 7, The production and properties of sintered Nd permanent magnets. In contrast, the hot-deformed magnet shows a two-step process with a rapid initial increase followed by a much slower increase. The explanation for this is that thermally demagnetized samples also contain domain walls that move easily until they are pinned, at which point the susceptibility decreases. This is consistent with Lorentz TEM studies (Mishra and Lee, 1986; Mishra, 1987a,b, Mishra et al., 1988; Liu et al., 2014, 2015) and which are discussed in more detail in Chapter 6, Hot-deformed NdFeB permanent magnets. In contrast, the individual grains in melt-spun NdFeB are believed to consist largely of single domains and magnetization requires the unpinning of a domain wall or the nucleation of a new wall in order for magnetization of a single grain to reverse. Therefore, the initial susceptibility is lower as shown in Fig. 3.29. This same characteristic has been observed in many other studies (Hadjipanayis et al., 1985a, 1986; Hilsher et al., 1986; Hadjipanayis and Kim, 1988a; Pinkerton and Fuerst, 1990), which also concluded that the initial magnetization in thermally demagnetized melt-spun ribbons was most likely controlled by domain wall pinning. However, Durst et al. (1987b) asserted that the initial susceptibility is not an accurate guide to determining the coercivity mechanism in these fine-grained melt-spun material. To prove their point, they prepared isotropic sintered magnets and found that the same initial magnetization behavior as for the melt-spun ribbon, that is, with maximum values of Br/Br(N) was not reached until HcHci. Since the coercivity in sintered Nd magnets is widely believed to be based on the nucleation of new domains, this would seem to be a strong argument that a nucleation model is also applicable in these fine-grained melt-spun materials and that the low initial susceptibility is simply an artifact of any isotropically aligned magnetic material. Girt et al. (2001) also studied the initial susceptibility in melt-spun samples having a range of Nd composition and also concluded that the initial susceptibility is not a good indicator of the coercivity mechanism if the average grain size is at or below the single domain size. This is because, in these very fine-grained materials, the same energy is required to unpin an existing domain wall as would be required to nucleate a new domain wall. If this is the case, then the coercivity in these finegrained melt-spun materials could equally be controlled by nucleation in which new domains are formed from the grain boundary phase, sweep across the grain, and then are annihilated when the domain reaches the opposite side of the grain. The initial susceptibility data in Figs. 3.29 and 3.33 show magnetization of a thermally demagnetized sample. As such, it does not address the question of the coercivity mechanism in a fully magnetized sample and how new domain walls form during the demagnetization process. Because the Nd2Fe14B intermetallic compound has very high magnetocrystalline anisotropy, the magnetization and demagnetization process must occur by the motion of domain walls. The Lorentz TEM studies, for example, Fig. 3.11B, of thermally demagnetized melt-spun ribbon seem to show that magnetization occurs by the motion and pinning of domain walls and

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

it is tempting to ascribe the same mechanism to demagnetization of a fully saturated sample. However, a number of studies have been carried out to determine the dominant coercivity mechanism during the demagnetization process. Folks et al. (1994b) and Crew et al. (1999) carried our reversible magnetization and magnetic viscosity studies of melt-spun ribbons. The goal of these studies was to determine the relative amounts of reversible versus irreversible processes during both the magnetization and demagnetization process, as well as the relative amounts of exchange and magnetostatic interactions. The viscosity studies involve magnetizing the sample in a large magnetic field, quickly applying a negative field and observing the time dependence of the magnetization as the field is held constant. The time dependence is measured for the same time period at different fields along the hysteresis curve. From these curves, information about the coercivity mechanism can be determined. These studies concluded that the coercivity mechanism between the magnetization and demagnetization process was different and that the mechanism during the demagnetization process was not necessarily different than that of sintered magnets, that is, that the coercivity mechanism was most likely based on a nucleation model. The various micromagnetic modeling studies have also generally supported a nucleation model based on a change in the anisotropy of the Nd2Fe14B phase at the grain boundaries owing to increased exchange interaction. This results in a reduction in the energy needed to nucleate reverse domains and a reduction in coercivity (Schrefl and Fidler, 1999; Fidler and Schrefl, 2000). However, others studies have tended to support a pinning model, and explain the drop in coercivity as due to a drop in the strength of domain wall pinning at the grain boundaries become of exchange interaction. Girt et al. (2000, 2001), Woodward et al. (2001), and Crew et al. (2002) carried out a comprehensive joint study, which included reversible magnetization and magnetic viscosity studies of both initial magnetization and demagnetization, as well as TEM, DSC, and micromagnetic modeling. They examined samples, which were prepared by annealing melt-spun ribbon, which had microstructures ranging from slightly to highly enriched in Nd and which were intended to change the amount of intergrain interactions as the amount of intergranular phase separating the Nd2Fe14B grains was changed. These study did find a significant difference in the magnetic behavior as the Nd2Fe14B grains became more isolated by the increasing amount of the Nd phase and attributed this difference to changes in intergranular exchange interaction. Magnetization reversal in the most Nd-rich alloy postulated that magnetization reversal occurs by the nucleation of domains at the edges of the isolated grains. In contrast, the studies concluded that the coercivity mechanism in the slightly Ndrich sample, comparable to the standard melt-spun materials discussed in this chapter, is probably controlled by pinning and that the pinning centers are created by strong intergrain exchange interaction. Consequently, the coercivity mechanism in these isotropic melt-spun NdFeB materials still is a subject of debate. Compounding this uncertainly is the fact that it is very difficult to distinguish the difference between unpinning of a domain wall from a defect and nucleation of a reverse domain, as has been pointed out by both Livingston (1985) and Hadjipanayis and Kim (1988a).

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There have also been a number of micromagnetic modeling studies (Fukunaga and Inoue, 1992; Schrefl et al., 1993; Schrefl et al., 1994c; Fidler and Schrefl 1996, 2000; Girt et al., 2001; Crew et al., 2002) that have been carried out with the aim of better understanding the magnetic properties of these fine-grained melt-spun materials. These studies have generally concluded that the magnetic properties are largely controlled by intergrain interactions between the individual Nd2Fe14B grains. Although most of these studied were carried out to better understand the magnetic properties of nanocomposite materials these modeling studies have also shed light on the magnetic properties of standard melt-spun NdFeB materials. Specifically, these studies conclude that exchange interaction can cause the magnetic moments near the grain boundary between two grains to become aligned, leading to a significant enhancement of Br and the observation that Br . Msat/2 is observed in these materials. However, another unfortunate conclusion is that exchange interaction can lead to a substantial drop in coercivity owing to a decrease in magnetocrystalline anisotropy at the grain boundaries, leading to a reduction in the field needed to nucleate a reverse domain or, alternatively, a reduction in the pinning strength of existing domain walls. These regions of low anisotropy at the grain boundaries result in a sharp drop in coercivity. Anything that reduces the amount of exchange interaction, such as the Nd-rich intergranular phase, would, therefore, increase the coercivity. This would explain why the coercivity of these melt-spun materials correlates strongly with the Nd content, provided the materials have been produced at the same quench rate. These various modeling studies have also considered the magnetization reversal process in NdFeB magnets as a function of grain size and grain alignment and all conclude that both are important factors that affect coercivity. Various modeling studies, including Fidler and Schrefl (2000), Thielsch et al. (2013), Sepehri-Amin et al. (2014) and Bance et al. (2014), have also carried out modeling studies on large-grain NdFeB materials with different particle sizes and shapes and all concluded that the coercive field in these materials will decrease with increasing grain size and increasing grain misalignment, as is observed experimentally for hotdeformed and sintered Nd magnets. Moreover, these studies have concluded that domain wall reversal occurs when the total reverse field, which is the sum of the local demagnetizing field and the applied field, reaches the field needed to nucleate a grain boundary. However, magnetostatic interactions are believed to dominate in these larger grain magnets and strongly affect the demagnetization field. These issues are discussed in more detail in Section 6.6. In contrast, exchange interactions are believed to dominate in fine-grained melt-spun materials and modeling is complicated by the fact that tfor exmaplehe microstructure consists of random isotropic alignment and a wide range in grain size. Also, as the grain size decreases for a fixed Nd content, there is a significant decrease in coercivity, which is believed to result from an increase in exchange interaction. However, the coercivity of any given grain is also believed to be affected by the demagnetization field as is the case in large-grained materials. The total demagnetizing field also changes as the angle that the total demagnetizing field makes with the c-axis of a given grain changes. Thus the coercivity has an angular component with the most misaligned

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grains relative to the applied field, having the highest coercivity. Since in an isotropic materials, the angle between the c-axis and the applied field vary by a full 180 degrees, this would suggest a demagnetization curve in which the coercive field in the individual grains vary from very low to very high levels, as observed experimentally. This would nicely explain the data in Fig. 3.29, where it is observed that the initial magnetization curve has an almost sinusoidal appearance, with the field level required for magnetization reversal increasing from initialing quite low to very difficult.

3.3.5 An interpretation of observed magnetic properties Any satisfactory explanation of the magnetic properties of these melt-spun NdFeB materials must explain the characteristic initial magnetization that is seen in Fig. 3.29 but also the dramatic change in magnetic properties of the melt-spun material as the quench rate and Nd content are varied. Fig. 3.34 shows a plot of the intrinsic coercivity of a Nd0.15(Fe12yBy)0.85 alloys versus vs, the substrate surface velocity of the melt spinner. These are the same data that are shown in Fig. 3.2. There is an obvious peak in the coercivity at a quench rate of B20 m/s, with the coercivity lower at both higher and lower quench rates. The peak in the coercive force is believed to occur at the optimum domain size which, based on TEM

20

y = 0.07

Hci(kOe)

15 y = 0.05 y = 0.09

10

y = 0.03 5

Nd0.15(Fe1-yBy)0.85 0 0

10

Vs(m/s)

20

30

Increasing quench rate Decreasing grain size Increased exchange interaction

Figure 3.34 The change in the coercivity for a melt-spun Nd0.15(Fe0.95B0.05)0.85 alloy as the grain size of the isotropic melt-spun ribbon is varied by changing the quench rate. Source: Adapted from Croat et al., 1984a. Appl. Phys. Lett. 44, 148.

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studies, appears to be B30 nm in these standard isotropic melt-spun ribbon. As the quench rate is reduced (lower vs) below the optimum value, the drop in intrinsic coercivity is thought to results from the change in the average microstructure to one having an average grain size that is larger than the single domain size and in which multiple domains are more easily formed. This drop in coercivity is readily explained because the these domains can move easily as a reverse field is applied, resulting in lower coercivity. However, the even more dramatic drop in intrinsic coercivity at higher quench rate cannot be so easily explained. In theory, the intrinsic coercivity should not drop nearly so dramatically as the average grain size is reduced below the single domain size and coherent rotation of the magnetic moments would not be expected to occur until a much smaller average grain size. The most satisfactory explanation for the behavior is that the decrease in the average grain size results in a significant increase in exchange interaction which, in turn, results in a drop in coercivity. The drop in coercivity is also believed to occur because the thickness of the intergranular phase surrounding each grain is significantly reduced as the quench rate is increased and the crystallite size decreases. Simple geometry would seem to require that as the grain size is reduced for a fixed composition, the increase in crystallite surface area will result in a thinner layer of the intergranular phase. Another example is shown in Fig. 3.35, which displays data

10

NdI-x(Fe0.95B0.05)x 8 x = 0.866 x = 0.85

x = 0.8

4

M (KG)

6

x = 0.75 x = 0.9

–15

–10

H (KOe)

–5

2

0

5

Decreasing Nd content Increased exchange interaction Decreasing coercivity

Figure 3.35 Data showing the dramatic change in the demagnetization behavior of meltspun Nd12x(Fe0.95B0.05)x alloys as the Nd content is varied. Both alloys were quenched at the optimum rate giving highest energy product. Source: Adapted from Croat, 1988. J. Mater. Eng. 10, 7
13.

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for a series of Nd12x(Fe0.95B0.05)x alloys, all of which were prepared at the optimum-quench rate. All of these samples were melt spun at close to the same optimum-quench rate and, therefore, should have approximately the same grain size. The dramatic decrease in coercivity is believed, therefore, to result from a decrease in the thickness of the layer of Nd-rich intergranular phase and a corresponding increase in exchange interaction between the individual Nd2Fe14B grains. This increase in exchange interaction is also believed to account for the large increase in first-quadrant magnetization that is observed for the x 5 0.9 alloy. Another example is shown in Fig. 3.8, which shows the demagnetization curves of a series of Nd12x(Fe0.95B0.5)x alloys that have all been melt spun at the same quench rate. As the Nd content is decreased (increasing x), there is a substantial reduction in coercivity, which is now believed due to a thinning of the layer of intergranular phase surrounding each Nd2Fe14B crystallite.

3.4

Nanocomposite or spring-exchange NdFeB magnets

The previous sections in this chapter discussed the properties of standard melt-spun NdFeB materials, which generally have a Nd content between roughly 26 and 28.0 wt% and a microstructure consisting of a majority phase of very small (B30 nm) Nd2Fe14B crystallites surrounded by a thin (1
2 nm) layer of a Nd-rich intergranular phase. The coercivity of these materials correlates closely with the Nd content and commercial powders today have coercivity values ranging from roughly 7 to 17 kOe. This section discusses the properties of nanocomposite or springexchange NdFeB materials, which are also typically prepared by crystallization of a melt-spun precursor. The major difference between these materials and standard powders is that the Nd content is much lower (10
12 wt%) and the microstructure is much different, consisting of a finely crystalline mixture of the magnetically hard Nd2Fe14B intermetallic phase and magnetically soft phases, which include Fe3B and α-Fe, usually combination of both. There is typically no Nd-rich intergranular phase separating the grains as is the case for standard powders. The original theory behind these materials is that exchange interaction between the hard and soft grains results in enhanced remanance values, which are much higher than that predicted for an isotropic material. Initially it was thought that nanocomposite magnets could be developed, which combined both high remanance and coercivity In practice, however, coercivity values have been disappointing, typically no higher than 4.0 kOe. It is now believed that the large exchange interaction, which provides the enhanced remanance, also reduces the coercive force of these material. Therefore, the reason for the lack of high coercivity in NdFeB nanocomposite materials is the same phenomena that cause a reduction in the coercivity of the standard powders as the Nd content is reduced, as was discussed in Section 3.4.3 through 3.3.5. Although the properties of these materials have not reached the point of commercial viability, they remain an interesting and promising family of permanent magnet materials. As mentioned several times in previous chapters, many of the end users for bonded Nd magnets, for example,

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micromotor manufacturers, would like magnets with the highest possible Br combined with a coercivity that is just high enough to prevent demagnetization during the life of the application. If even moderately higher coercivity levels could be developed, these magnets would become possible candidates for these applications. The lower Nd content would also result a substantial reduction in direct material cost, since the Nd is by far the most expensive component.

3.4.1 The theory of spring-exchange magnets Spring-exchange or nanocomposite magnets are combinations of hard and soft magnetic materials, typically combination of Nd2Fe14B, Fe3B, and α-Fe, and usually contain some of all three phases. These magnetic materials were first reported by Coehoorn et al. (1988), who found that crystallization of a melt-spun Nd4Fe78B18 alloy resulted in a fine-grained magnet consisting of 15 vol% Nd2Fe14B, 73 vol% Fe3B, and 12 vol% α-Fe. This material was found to have a Br 5 12 kG, Hci 5 4 kOe, and an energy product (BH)max 5 11.9 MGOe. The Br value was substantially higher than the reduced remanance value of Br/Msat 5 0.5 predicted for an isotropic magnet, leading to the speculation that the enhanced remanance was due to exchange interaction between neighboring grains. One simple way to visualize the principle behind these materials is seen in Fig. 3.36, which shows an overlay of the full hysteresis curves of a hard magnet and a soft magnet. The theory behind these materials is that the hard material helps retain the soft material’s anisotropy, which increases its coercivity, and the combination takes on a shape resembling that of the sum of its hard and soft magnetic components, represented by the dashed hysteresis curve, and resulting in an energy product that is higher than those of its M

H

Figure 3.36 Hysteresis curves of the hard and soft magnetic material whose sum is the hysteresis curve of an idealized spring-exchange magnet.

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

components. The scientific principle behind these magnets is that exchange interaction between the magnetic moments in the soft and hard magnetic phases causes the magnetic moments within the soft grains to arrange themselves parallel to the average direction of the magnetization in the hard magnetic grains. This exchange interaction results in a higher remanance, which exceeds the theoretical limit Br/ Msat 5 0.5 for an isotropic material and is driven by a reduction in the total energy of the magnet when the magnetic moments become more parallel. If the grains of the magnetically soft material are sufficiently small, then the moment of the entire particle will align parallel with the moment of the hard magnetic phase. If the grains are larger, then it is believed likely that only the moment near the perimeter of the soft phase becomes aligned by the exchange interaction. The volume fraction of the soft phase needs to be as large as possible in order to achieve high magnetization. In theory, Br values as high as 21.5 kG and energy products as high as 115 MGOe are possible (Skomski and Coey, 1993). In practice, however, magnetic properties obtained have been substantially below these values and, in particular, the highest coercivity values achieved have generally been .5 kOe. Fig. 3.37 shows a drawing of a hypothetical nanocomposite microstructure. Here the white grains are the Nd2Fe14B phase and the hashed grains are the soft magnetic phase, either Fe3B or α-Fe. This microstructure is B1/4 hard magnetic phase and B3/4 soft magnetic phase. The arrows represent the magnetic moments in both the soft and hard phases. This drawing shows a situation where the sample has been magnetized and then the magnetic field has been returned to zero. In the magnetized state, all of the moments in the soft magnetic phase would become aligned with the magnetic field and the moments in the Nd2Fe14B grains would lie along the c-axis direction most closely aligned with the applied magnetic field. When the field is removed, the moments in the Nd2Fe14B grains would continue to lie along

Figure 3.37 Representation of a nanocomposite magnet consisting of particles of Nd2Fe14B (white grains) and a soft magnetic phase such as Fe3B or α-Fe (hatched grains). This drawing represents the nanocomposite material after magnetizing and then removing the magnetic field.

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the c-axis direction closest to the applied magnetic field. However, exchange interaction causes the magnetic moments of the soft magnetic grains to align themselves parallel to the average direction of the magnetization in the neighboring hard magnetic grains. This is the situation that the drawing in Fig. 3.37 is meant to represent. Fig. 3.38 is a representation of the magnetic moments in this material, where the moments of the hard phase are the gray arrows and soft material are black. The situation seen in Fig. 3.38A is the same as the drawing in Fig. 3.37, in which all of the moments in the Nd2Fe14B grains are aligned along their c-axis direction as close to the field direction as possible and the moments in the soft magnetic grains are aligned parallel to the average moments of the hard magnetic phase. The force that causes this is exchange interaction between the soft and hard magnetic grains combined with the anisotropy of the Nd2Fe14B intermetallic phase which, in order to minimize the total magnetic energy, wants to hold the moments of the soft phase parallel to those of the hard magnetic phase. The moment of the soft phase is held in alignment with those of the hard phase as a reverse field is applied, resulting in a significant coercive force. Because of this coercive force, only part of the moments of the soft magnetic phase reverses in an applied field H
Hci, as depicted in Fig. 3.38B. After subsequently removing the external field, the rotated moments in the soft phase will rotate or spring back into alignment once again with the average moment of the hard magnetic phase, as shown in Fig. 3.38C. Because of this behavior, the hysteresis curves of these spring-exchange magnets have a very sharp recoil curve, as is shown in the comparison of the hysteresis behavior of a spring-exchange magnet with a traditional permanent magnet in Fig. 3.39. This sharp recoil behaves prompted Kneller and Hawig (1991) to coin the expression of spring-exchange magnets.

H=0 (A)

(B)

(C)

H = Hci

H=0

Figure 3.38 A representation of the magnetic moment in a spring-exchange material as a field is applied and then removed. The gray arrows are the moments in the Nd2Fe14B grains and the black arrows represent the moments in the soft magnetic grains. The figure shows the direction of the hard and soft magnetic moments for (A) H5 0, (B) after applying a reverse field H5 Hci and (C) after returning the field level back to H5 0.

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Exchange spring magnet

Conventional magnet

M

M

H

H

Reversible Irreversible

Figure 3.39 Comparison of the hysteresis curves of an exchange spring magnet and a conventional magnet. The demagnetization curves of the exchange spring magnets are highly reversible and when the reverse field is removed, the magnetization recoils or springs back to close to the original Br value (Kneller and Hawig, 1991).

3.4.2 Experimental studies of nanocomposite magnets Following the first report of nanocomposite materials by Coehoorn et al. (1988), these materials became the subject of intense investigation in the early 1990s to better understand their underlying physical principles and to determine if better magnetic properties, notably higher coercivity, could be achieved (Coehoorn et al., 1989; Schneider et al., 1990; Eckert et al., 1990; Manaf et al., 1991, 1993; Davies, 1996). The standard process by which these materials are produced is by melt spinning followed by annealing or crystallization to produce a microstructure consisting of microcrystals of a magnetically hard phase and a magnetically soft phase. Alternatively, these materials have been produced by mechanical alloying, a technique which involves the use of a high energy ball mill to blend the elements into an amorphous state, which is then annealed to form the desired two-phase microstructure. Neu et al. (1996) used this technique to prepare Nd2Fe14B/α-Fe-type alloys with varying amounts of Co and achieved properties of Br 5 11.5 kG, Hci 5 5.2 kOe, and (BH)max 5 18 MGOe, some of the highest properties reported for these spring-exchange materials. Similar mechanical alloying studies were carried out by McCormick et al. (1998). Eckert et al. (1990), Schneider et al. (1990), and Kneller and Hawig (1991) were the first to propose that the properties of these materials resulted from exchange interaction between the soft and hard magnetic phases and to explain the steep recoil in the hysteresis curves. Micromagnetic modeling studies of these nanocomposite materials (Kneller and Hawig, 1991; Schrefl et al., 1993, 1994a,b; Fisher et al., 1996, 1998; Fidler and Schrefl, 1996, 2000; Fisher and Kronmuller, 1996; Bauer et al., 1996; Schrefl et al., 1997; Schrefl and Fidler, 1999; Schrefl, 1999; Fullerton et al., 1999; Fukunaga and Mukaino, 2005) were carried out with the aim of predicting the microstructural conditions required to achieve optimum combinations of high

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remanance and high coercivity, including the ideal size and distribution of the soft and hard magnetic grains. These studies all concluded that a uniform, fine grain structure was necessary to achieve optimum properties. These modeling studies also addressed the mechanism of the magnetization reversal process and have provided guidelines for the possible optimization of these materials, including the reasons why high coercivity has not been achieved to date in these materials. All of the early micromagnetic modeling studies recognized the importance of the grain size on the exchange interaction. Therefore, many of the experimental studies aimed at improving the coercivity of these materials were directed at ways by which uniform, very fine-grained microstructures could be produced. Refinement of grain size has usually been accomplished by introducing elements that act as grain growth inhibitors and, not surprisingly, the additives tested are the same grain growth inhibitors that were used in studies related to standard isotropic powders. These studies have included the effect of Al, Si, Ga, Ag, and Au (Kanekiyo et al., 1993; Hirosawa et al.,1993), Cr (Hinomura et al., 1997; Uehara et al., 1998), Co, Dy, and Ga (Mishra and Panchanathan, 1994), Co (Neu et al., 1996), Ga and Co (Ping et al., 1998), Cu and Nb (Sano et al., 1999; Ping et al., 1999a,b), Cu, Nb, and Zr (Hirosawa et al., 2002), Ti and C (Hirosawa et al., 2004), Cu, Ti, and Cu 1 Ti (Yang et al., 2003), Zr and Ti (Jianu et al., 2004), and Co and Nb (Dospial et al., 2012). Most of these studies were carried out by crystallizing melt-spun alloys. For almost all of the additives, a finer grain size was achieved relative to the alloy without the additive. However, the coercivities obtained remained stubbornly low, generally between 3 and 5 kOe. For example, Ping et al. (1999a,b) studied the effect of Cu and Nb additions on Nd2Fe14B/Fe3B alloys and found optimum results of Br 5 12.5 kG, Hci 5 3.43, and (BH)max 5 15.7 MGOe in a melt-spun Nd4.5Fe75.8B18.5Cu0.2Nb1 ribbon after crystallization at 600 C for 6 minutes. Their 10 min @ 660iC

Nd4.5Fe77B18.5 Nd4.5Fe76.8B18.5Cu0.2 Nd4.5Fe75.8B18.5Nb1Cu0.2

–0.4 –0.3 –0.2 –0.1

0

0.1

0.2

0.3

0.4

H (MA/m)

Figure 3.40 Demagnetization properties of Nd0.45Fe77B18.5, Nd4.5Fe76.8B18.5Cu0.2, and Nd4.5Fe75.8B18.5Nb1Cu0.2 melt-spun ribbons annealed for 10 minutes at 660 C (Ping et al., 1999a,b).

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results, shown in Fig. 3.40, are fairly typical of the demagnetization behavior of these materials to date. They observed that both Cu and Nb additions resulted in a finer grain size and an increase in coercivity over ternary NdFeB alloys. They also observed clusters of Cu, which they postulated served as nucleation sites for the crystallization of the Fe3B phase with a finer grain size. Ping et al. (1998) studied the effect of Ga and Co additives on melt-spun Nd2Fe14B/Fe3B-type alloys with a composition of Nd4.5Fe73B18.5Co3Ga1. The alloy was melt spun on a Cu rim at a rim speed of 20 m/s and the ribbon material then annealed for 10 minutes at varying temperatures. Fig. 3.41A shows a bright field TEM image after annealing at 580 C and shows a fully crystalline microstructure consisting of Nd2Fe14B, Nd2Fe23B3, and Fe3B grains with an size of B20
30 nm. One interesting observation was that the Co was distributed in both of the ternary phases but Ga was only found in the Nd2Fe14B phase. No Co or Ga was found in the Fe3B phase. Fig. 3.41B shows the same material after annealing at 700 C. In this alloy, the Nd2Fe23B3 had completely disappeared and the average grain size had increase to between 25 and 50 nm. Unfortunately, the highest Hci observed was only 3.9 kOe. These images shown here are, again, fairly typical of the type of fine-grained microstructure that is obtained for all of these spring-exchange magnets. However, in both of these studies, the highest coercivities were not obtained at the smallest grain size as had been predicted. Various studies have also investigated the use of high anisotropy rare earths to increase the Hci of the nanocomposite magnets. For example, Bernardi et al. (2000) prepared samples of Nd3.25Tb1Fe72.15Co5B18 by several different techniques including melt spinning, splat cooling, and mechanical alloying. Remanance values between 10.5 and 11.0 kG were obtained with Hci values as high as 6.28 kOe. The best results were found for samples that were melt spun and then rapidly annealed. They reported this sample to consist of B50% soft phase (both α-Fe and Fe3B) and 50% (Nd,Tb)2Fe14B intermetallic phase with a uniform microstructure. There have also been a number of studies in which the Nd or Nd/Pr content was more midway between the typical nanocomposite composition and that of the

Figure 3.41 Bright field TEM of melt-spun Nd4.5Fe73B18.5Co3Ga1 alloys after annealing for 10 minutes at (A) 580 C and (B) 700 C (Ping et al., 1998).

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standard powders. For example, Berra-Barrera et al. (2006) prepared Pr5Fe772xCrxB18 alloys with x 5 0, 1, 2, 2.5, 3, and 4. They reported that the addition of Cr increased Hci by B50% over the Cr-free alloy. Here the 5 at% Nd equals B13.8 wt% for an alloy with x 5 3. They did, however, report Hci values as high as 7 kOe, which are among the highest reported for these nanocomposite materials. In a similar study, Tao et al. (2012) studied Fe65B22Nd9Mo4-type alloys and obtained Hci values as high as 11.55 kOe but Br values were, again, no higher than 5.6 kG. Here the 9 at% Nd converts to 24 wt%, almost as high as in standard melt-spun powders.

3.4.3 Coercivity and intergrain interaction in nanocomposite magnets The coercivity of spring-exchange magnets arises entirely from the anisotropy of the Nd2Fe14B intermetallic phase. Both the Fe3B and α-Fe phases, while having high saturation magnetization, have tetragonal and cubic crystal structures, respectively, with very low planar anisotropy. However, in these fine-grained material, the anisotropy of the magnetically hard Nd2Fe14B phase holds the moments of the soft phase into alignment as the field is reversed. The question is why the exchange interaction between the hard Nd2Fe14B phase and the soft magnetic phases does not produce higher coercivity? Micromagnetic modeling studies have provided a quantitative description of the correlation between the microstructure and magnetic properties and provide several reasons for the lower than expected coercivity. First, although exchange interaction can align the moments of the magnetically soft grains with the average moment of the hard grains leading to enhanced remanance, this also leads to a drop in coercivity because the exchange interaction between grains also lowers the energy barrier for domain wall reversal by reducing the anisotropy of the Nd2Fe14B grains at the grain boundaries. This is similar to that observed for standard NdFeB magnetic powders as the Nd content is lowered and exchange interaction increased. This is also thought to be the reason that higher coercivity is not observed as the grain size of the magnetically hard and soft phases is reduced, since reduction in the average grain size results in increased exchange interaction, leading to higher remanance but also to lower coercivity. Consequently, the strong exchange interaction and enhanced remanance, which made these materials so initially attractive, are also now believed to be at least partially responsible for the low coercivity that is observed. The inability to achieve higher Hci and energy densities has also been attributed to the formation of large interaction domains (Shield et al., 2006), which are believed to form because of strong exchange interaction between the grains. These large interaction domains can lead to cooperative magnetization reversal, resulting in lower coercivity. Thus the formation of interaction domains in these materials conflicts with achieving higher coercivity. A third reason proposed for the low coercivity is due to magnetostatic effects, which are also believed to be important in these nanocomposite materials despite the very small average grain size (Fidler and Schrefl, 1996, 2000;

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Rapidly Solidified Neodymium-Iron-Boron Permanent Magnets

Schrefl and Fidler, 1999; Gabay and Hadjipanayis, 2007). It is proposed that magnetostatic interactions can produce flux closure structures in regions of the magnetically soft Fe3B or α-Fe grains in which local magnetization is zero. Upon application of a reverse field, the magnetization within these regions can reverse easily, leading to a decrease in the coercive field. In combination, these problems pose a daunting task for further research into these material. As pointed out by Shield et al. (2006), there is no good solution to the problem, since any attempt to reduce intergrain exchange interaction and the formation of large interaction grains also leads to a reduction in the enhanced remanance. Despite these problems, these nanocomposite materials remain an interesting family of permanent magnets materials and research continues to be carried out. As pointed out before, magnets with higher Br and just enough coercivity to prevent demagnetization under dynamic conditions are what much of the micromotor industry would like to see developed. Coercivity levels would have to increase only marginally for these materials to be good candidates for these micromotor applications. In addition, price is always important, and the lower Nd content of these nanocomposite type materials would provide magnets with significantly lower direct material cost.

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Selected Readings Coehoorn, R., de Mooij, D.B., de Waard, C., 1989. Melt spun permanent magnet materials containing Fe3B as the main phase. J. Magn. Magn. Mater. 80, 101. Cullity, B.D., 1972. Introduction to Magnetic Materials. Addison-Wesley Publishing Co, Reading MA. Fidler, J., Schrefl, T., 2000. Micromagnetic modeling-the current state of the art. J. Phys. D: Appl. Phys. 33, R-135. Mishra, R.K.J., 1986. Microstructure of melt spun Nd-Fe-B Magnequench magnets. Magn. Magn. Mater. 54, 450. Pinkerton, F.E., Van Wingerdon, D.J., 1986. Magnetization process in rapidly solidified neodymium-iron-boron permanent magnet materials. J. Appl. Phys. 60, 3685.

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