The 2:14:1 phase has a P4₂/mnm tetragonal space group with four formula units or 68 atoms per unit cell. Nd₂Fe₁₄B is the most common studied material with 2:14:1 phase structure. The microstructure of single-phase Nd₂Fe₁₄B melt-spun materials is fairly uniform, consisting of randomly oriented grains of the magnetically hard phase as shown in Fig. 11(a). In the single-phase stoichiometric Nd₂Fe₁₄B magnets, the magnetically hard grains are directly connected without grain boundary phase and therefore they are magnetically coupled by exchange interaction, which is symbolized by a spring in Fig. 12(a).^[73,74] High coercivity can be easily obtained in the single-phase Nd₂Fe₁₄B magnets with grain sizes below 80 nm. This grain size is far below the single domain particle size, which is a fraction of for most of the high anisotropy materials. Figure 11(b) shows the high-resolution TEM micrograph of a stoichiometric Nd₂Fe₁₄B magnet with an average grain size of 20 nm.^[75] No grain boundary phase is observed. Since the magnetically hard grains are directly contacted in the stoichiometric magnets, a significantly enhanced remanence (m_r is larger than 0.5) and a slightly reduced coercivity are observed as a consequence of exchange coupling between the grains, as shown in Fig. 12(d), which will be discussed in details in Section 4. As the exchange interaction is a short-range interaction, which is in the order of ∼4 nm for Nd₂Fe₁₄B, the grain size should be smaller than 50 nm in order to notice the remanence enhancing effect macroscopically.^[17]

Fig. 11.

	Figure Option ViewDownloadNew Window
	Fig. 11. (a) Typical microstructure of nanocrystalline Nd2Fe14B;[76] high-resolution TEM of (b) stoichiometric and (c) decoupled 2:14:1 magnets.[75]

Fig. 12.

	Figure Option ViewDownloadNew Window
	Fig. 12. Schematic of microstructure for the remanent state of (a) stoichiometric and (b) decoupled 2:14:1 based nanocrystalline magnets, (c) nanocomposite magnet, and (d) their hysteresis loops at room temperature.

To achieve even higher coercivity in a permanent magnet, the magnetically hard grains need to be decoupled. Although an amorphous grain boundary can be formed by optimum the cooling conditions in the as-spun ribbons, most commercial alloys contain excess Nd. The resulting microstructure obtained by melt-spinning qualitatively agrees with the situation in sintered permanent magnets where the single crystallites of the magnetically hard phase are more or less magnetically decoupled by a paramagnetic RE-rich boundary phase. However, the scale of the melt-spun microstructure is about a factor of 100 smaller than sintered magnet and the easy axes of the nanograins are isotopically distributed instead of aligned. The high-resolution TEM micrograph in Fig. 11(c) presents the intergranular film with a thickness of about 2 nm as a bright region between the nanograins.^[75] Figure 12(b) schematizes the decoupling RE-rich grain boundaries. The grain size in such decoupled magnet ranges from 50 to 200 nm, which guarantees that mainly single-domain particles exist. Therefore, a coercivity higher than 2.5 T can be achieved, which is larger than the stoichiometric single-phase magnet, as schematized in Fig. 12(d). However, since the easy-axis direction of the grains are isotopically distributed and magnetically decoupled, the maximum normalized remanence of 0.5 can only be achieved in such magnets according to Stoner–Wohlfarth model.^[10]

If the alloy composition is chosen near to the two-phase line Fe–T₁ on the left region bordering the stoichiometric T₁ phase (as shown in Fig. 3), a so-called composite magnet can be obtained, showing a further significant remanence enhancement, accompanied by a reduction of the coercivity as can be seen from Fig. 12(d). In such two-phase magnets, nanocrystalline soft magnetic α-Fe grains arise in addition to the magnetically hard 2:14:1 grains from the melt-spinning process as a consequence of the over-stoichiometric α-Fe, as schematized in Fig. 12(c). Thus, the remanence enhancement is due to the exchange coupling among the grains but also to the very large spontaneous polarization of the α-Fe grains. The composite magnets will be discussed in Section 3.2 in details.

The Nd₂Fe₁₄B phase is the most widely studied permanent magnetic materials by melt spinning as well as other techniques due to the excellent magnetic properties, i.e., high anisotropy field ( ), and high saturated magnetization (∼1.6 T). Croat et al. reported the single-phase Nd₂Fe₁₄B melt-spun materials having coercivity of about 1.5 T and energy product of about 14 MGOe under an optimum wheel speed of 16–20 m/s as early as 1984.^[8] However, the Nd₂Fe₁₄B magnets have two major disadvantages compared to other permanent magnets. The Curie temperature of Nd₂Fe₁₄B magnet is only about 315 °C, which results in large negative temperature coefficients of remanence and coercivity, and therefore limits their application to 80 °C or lower. The material cost of Nd–Fe–B magnet is still far higher than that of ferrite magnets. Therefore, many 2:14:1-type magnets derived from Nd₂Fe₁₄B are developed through elemental additions or substitutions. They can be classified into the following categories based on the type of element substitutions:

(i) RE substitution for Nd

Replacing Nd by HRE such as Dy or Tb is an effective to raise the coercivity for higher temperature application since Dy₂Fe₁₄B and Tb₂Fe₁₄B have much higher anisotropy field comparing to Nd₂Fe₁₄B magnets, as seen in Table 5. Larger coercivity of up to 64 kOe at room temperature are reported for Dy₂Fe₁₄B and Tb₂Fe₁₄B alloys magnetically hardened by melt spinning technique,^[77] which is over three times larger than that of Nd₂Fe₁₄B ribbons. Even though they may have similar negative temperature coefficient of coercivity as Nd₂Fe₁₄B, they can provide relatively large coercivity to satisfy the higher temperature applications. However, it has to be borne in mind that there is normally a trade-off between higher coercivity and higher costs for the HRE substitutions. In addition, the antiferromagnetic HRE sublattice may lead to a strong magnetization reduction, as also given in Table 5.

Table 5.

Table 5.

Table 5. Intrinsic properties theoretical (BH)max limits of 2:14:1 compounds.[83–86] . RE Tc/°C μ0Ms/T K1/ μ0HA/T Theo. (BH)max/MGOe RE2Fe14B Y 298 1.38 1.06 1.9 48 La 243 1.38 1.10 2.0 48 Ce 260 1.17 1.7 3.7 34 Pr 292 1.56 5.6 9.0 61 Nd 315 1.60 5 7.6 64 Sm 345 1.50 −12 −20.1 56 Gd 387 0.84 0.67 2.0 18 Tb 356 0.70 5.9 21.2 12 Dy 320 0.71 4.5 15.9 13 Ho 301 0.81 2.5 7.8 16 Er 284 0.90 −0.03 −0.1 20 Tm 267 1.15 −0.03 −0.1 33 Lu 266 1.17 1.17 2.6 34 RE2Fe14C Pr 239 1.67 9.8 14.8 70 Nd 262 1.65 6.2 9.5 68

Table 5. Intrinsic properties theoretical (BH)max limits of 2:14:1 compounds.[83–86] .

Besides high-temperature application, element substitution is also useful for low-temperature application. For example, Pr₂Fe₁₄B magnets can be used at low temperature because there is no spin reorientation, while they have similar magnetic properties to those of Nd₂Fe₁₄B. Hadjipanayis et al. reported a coercivity of 2 T and energy product of 12 MGOe in the melt-spun single-phase Pr₂Fe₁₄B ribbons.^[78]

Although RE elements are not truly rare geologically, they are expensive due to the extraction process, which makes the rare-earth magnets much more expensive than those of ferrite magnets. In order to reduce the material cost while utilize the still high magnetocrystalline structure of 2:14:1 phase, La/Ce based RE₂Fe₁₄B magnets attract the interests recently. Significant cost reduction can be achieved due to the much cheaper price of La/Ce by slightly sacrificing magnetic properties. An energy product of 8.6 MGOe is achieved in the Ce–Fe–B melt-spun ribbons as reported by Zhou et al.^[79] If Nd is only half replaced by Ce, the energy product of 14.3 MGOe can be obtained through melt spinning as reported by Li et al.^[80] The magnetic properties of the Ce-substituted Nd₂Fe₁₄B ribbons are summarized in Table 6. More commonly, the so-called mischmetals (MM) which contains ∼50 wt% Ce and ∼25 wt% La, are used to prepare 2:14:1-type magnets in order to further reduce cost by avoiding the extraction of Ce/La from each other. Zhang et al. reported an energy product of about 8.3 MGOe in the single-phase (La_0.35Ce_0.65)₂Fe₁₄B ribbons prepared at a wheel speed of 20 m/s.^[81] Hussain et al. recently reported a La/Ce/La_0.5Ce_0.5 substituted Fe-rich 2:14:1 ribbons with improved magnetic properties (μ₀M_r = 1.07 T, μ₀H_c = 0.59 T, and (BH)_max = 18.4 MGOe), although it is a nanocomposite instead of single-phase magnet.^[82] The relatively high coercivity in these ribbons may be attributed to the phase segregation with certain La or LaCe content.

Table 6.

Table 6.

Table 6. Magnetic properties of melt-spun Nd12−xCexFe82B6 ribbons.[80] . x Tc/°C μ0Mr/T μ0Hc/T (BH)max/MGOe 0 366 0.98 1.09 20.9 1 357 0.93 1.04 18.0 2 348 0.91 0.96 17.2 4 331 0.87 0.84 16.0 6 313 0.84 0.77 14.3 8 298 0.77 0.68 12.1 10 279 0.74 0.52 10.8 11 268 0.68 0.41 8.04 12 253 0.64 0.32 5.54

Table 6. Magnetic properties of melt-spun Nd12−xCexFe82B6 ribbons.[80] .

(ii) Co substitution for Fe

Replacing Fe by Co substitution is an effective way to raise Curie temperature and potentially reduces temperature coefficients of coercivity and remanence. It is known that Nd₂(Co_xFe_1−x)₁₄B pseudo-ternary systems form the 2:14:1-type crystal structure for all levels of cobalt substitution, with T_c increasing monotonically to ∼722 °C for Nd₂Co₁₄B.^[87] It is found that the desirable enhancement of T_c is accompanied by improved temperature stability of the remanence in the Nd₂(Co_xFe_1−x)₁₄B melt-spun ribbons if the Co content is less than 30%.^[88] However, the coercivity deteriorates dramatically as the Co content increases above 40% due to the easy formation of Nd₂Co₁₇ rather than Nd₂Co₁₄B. By systematically varying the composition of Nd–Co–B alloys, Fuerst and Herbst successfully developed melt-spun Nd₂Co₁₄B ribbons with a coercivity of 1.02 T and energy product of 5.2 MGOe.^[89] The authors also fabricated the ribbons comprised principally of Pr₂Co₁₄B with a coercivity of 2.5 T and energy product of 5.5 MGOe.^[90] It should be mentioned that RE₂Co₁₄B compounds do not form for all rare earths but only for R = Y, La, Pr, Nd, Sm, Gd, and Tb. Unfortunately, none of them has attracted much attention due to their unsatisfied properties for permanent magnetic applications. In advanced Nd–Fe–B permanent magnets, Fe is only partially substituted by Co.

(iii) C substitution for B

The magnetic properties of the carbide and boride 2:14:1 compounds are almost same. Typically, carbides have slightly lower Curie temperatures and saturation magnetizations, slightly higher magnetocrystalline anisotropies, as shown in Table 5. The interest of RE₂Fe₁₄C phases is not only because of their intrinsic properties but also the presence of solid-state transformation, which offers the possibility of generating small-particle microstructures by controlling transformation of the high-temperature phase into the low-temperature phase. In another words, the nucleation and grain growth of the 2:14:1 phase in the RE–Fe–C systems is lower than that of RE–Fe–B systems, which should lead to finer microstructures. It has been found that Nd₂Fe₁₄C fabricated by melt spinning could easily provide a relatively high coercivity (∼10 kOe) without the aid of other additives.^[91] With a small portion of Co, Dy adoption, Zhang et al. reported an energy product of 11.5 MGOe with a coercivity of 2.7 T in the Nd₁₃Dy₂Fe₆₇Co₁₀C₆B₂ melt-spun ribbons.^[92] The energy product was further improved to 14.4 MGOe through optimizing post-annealing conditions.^[93] Zhang et al. also studied the magnetic properties of Pr₁₅Fe₇₇C_8−xB_x (x = 0–8) ribbons and found the high coercivity of 2.4 T in the Pr₁₅Fe₇₇C₄B₄ ribbons without any heat treatment.^[94]

(iv) Other additives

With no substantial cost increase, researchers found that the magnetic properties are improved by means of small amounts of additives in the 2:14:1 base materials. The most common additives are Ti, V, Cr, Mn, Co, Ni, Cu, Zr, Nb, Mo, Ru, W, Al, Ga, Si, etc. According to the summary in Ref. [83], the additions of Co, Ni, Cu, Ga, and Si increase the Curie temperature, while the additions of Nb, Al, and Si increase the magnetocrystalline anisotropy of Nd₂Fe₁₄B phase. More importantly, the additives can modify the microstructure in order to achieve desirable magnetic properties. For example, Herbst et al. found that the addition of 0.5 at.% Cu enhances the intrinsic coercivity as much as 30% over that of Cu-free materials, which is attributed to the Cu preferential segregation in the intergranular regions.^[95] El-Moneim et al. observed that partial substitution of Fe with Co and Ga leads to an improvement in corrosion resistance.^[96] Chen et al. found that a small substitution of Nb for Fe results in an uniform microstructure with average grain size of only 10–20 nm, which is much smaller than the Nb-free ribbons.^[97]

It should be emphasized again that the melt-spun materials based on 2:14:1 phase are typically isotropic which limits the maximum achievable energy products far lower than the theoretical value given in Table 5. Therefore, 2:14:1 phase usually exists as the magnetically hard phase in their nanocomposite magnets, which are discussed in section 3.2.

The CaCu₅-type 1:5 phase can be formed in the RE–Co and RE–Ni compounds with the composition of RE(Co, Ni)₅. However, the highest Curie temperature of −240 °C is found in the RENi₅ series for GdNi₅, which makes them not practical for permanent magnet applications. On the other hand, Curie temperatures of most RECo₅ compounds are higher than 500 °C, as shown in Table 7. In addition, the extra-large magnetocrystalline anisotropy and the large magnetic moment of Co in the RECo₅ compounds make them good candidates for permanent magnets.

Table 7.

Table 7.

Table 7. Intrinsic magnetic properties and theoretical (BH)max limits of 1:5, 2:17 and other phases.[19,98,99] . Compound Tc/°C μ0Ms/T K1/ μ0HA/T Theo. (BH)max/MGOe YCo5 630 1.06 5.5 13.0 28 LaCo5 567 0.91 6.3 17.5 20 CeCo5 380 0.77 6.4 21.0 15 PrCo5 620 1.20 8.1 17.0 36 NdCo5 637 1.22 0.24 5.0 37 SmCo5 727 1.14 11-20 25.0-44.0 32 SmCo4B 237 1.25 3.2 6.5 39 Sm2Co17 920 1.25 3.2 6.5 39 Sm2(Co0.7Co0.3)17 840 1.45 3.0 5.2 52 Sm2Fe17N2.94 473 1.51 8.4 21.0 57 Sm2Fe17C 279 1.24 2.6 5.3 38 SmFe11Ti 312 1.16 3.4 7.36 34 SmFe10V2 337 1.10 2.1 4.80 30 Sm3(Fe, Ti)29N 475 1.18 5.6 12 35 Sm5Fe17 280 0.95 23

Table 7. Intrinsic magnetic properties and theoretical (BH)max limits of 1:5, 2:17 and other phases.[19,98,99] .

3.1.2.1. SmCo₅

Among RECo₅ compounds, SmCo₅ compounds are obvious candidates for developing coercivity, owing to their extraordinary magnetocrystalline anisotropy. However, the CaCu₅-type 1:5 phase is stable only at temperature above 805 °C according to Sm–Co phase diagram,^[100] below which it decomposes into Sm₂Co₇ and Sm₂Co₁₇ phases. The magnetocrystalline anisotropies of the 2:7 and 2:17 phases (∼20 T and ∼6–11 T, respectively) are less than that of 1:5 phase (25–44 T). In addition, the 2:7 phase can act as nucleation site of reversing domain wall, which decreases the coercivity of the SmCo₅ magnets. Therefore, fast cooling rate is necessary to obtain a pure single phase SmCo₅ magnet by suppressing the formation of the 2:7 and 2:17 phases, which makes the melt-spinning technique very suitable to fabricate the 1:5-type magnets.

Takahashi et al. prepared the melt-spun SmCo₅ ribbons in 1985,^[101] where they found that the columnar structure with c-axis aligned parallel to the ejecting direction of the ribbons at relatively low wheel speed below 18 m/s. This is the so-called preferred crystallographic orientation or texture. In the ribbons fabricated at higher wheel speed, only homogenous spherical fine particles with random orientation of c-axis were observed. Since the Co-rich part of the Sm–Co phase diagram does not have a deep eutectic, the melt spinning does not lead to an amorphous structure or even a very fine crystallized microstructure, which is the requirement for high coercivity. The coarse microstructure in typical Sm–Co melt-spun ribbons leads to a coercivity lower than that expected for a material with such a large anisotropy field. Therefore, it understandable that the above work only achieves a coercivity up to 1.5 T,^[101] which is even lower than of the melt-spun Nd₂Fe₁₄B ribbons, although SmCo₅ has much higher magnetocrystalline anisotropy. Ding et al. optimized the melt spinning process for SmCo₅ ribbons and confirmed the dendritic structure with needle-like crystallites aligning preferentially in the ribbon plane in the low-speed prepared ribbons (v_ws –30 m/s).^[102] The coercivity of ribbons with such structure only have a low value of 0.96 T. Increasing wheel speed to leads to significantly increase of coercivity to 3.1 T due to the decrease of crystallite size, while with the sacrifice of texture in the ribbons. Suresh et al. reported a high coercivity of 5.3 T in the SmCo₅ ribbon melt-spun at 50 m/s, which has a small grain size of about 55 nm.^[103]

Between the texture and the coercivity of melt-spun SmCo₅ ribbons, texture attracts more attention due to the demand of higher energy products. Therefore, most of the recent efforts were focused on solidification at relatively low speeds. Actually, Ding et al. reported an energy product of 18 MGOe in the SmCo₅ ribbons prepared at wheel speed of 6 m/s.^[102] Yan et al. also reported the formation of crystallographic texture during melt-spinning process using low wheel speed.^[104] The degree of the preferential alignment decreases at high wheel speed, while the ribbons prepared at are completely isotropic. Figure 13 shows the typical x-ray diffraction (XRD) patterns of the melt-spun SmCo₅ ribbons prepared at low and high wheel speeds. The low-speed prepared ribbons clearly give very strong reflections of (200) and (110), while those of (111) and (002) are greatly weakened, comparing to the isotropic sample. This implies that the c-axis of crystallites prefers to be parallel to the longitudinal axis of the ribbon, which is different from the results reported for the Nd₂Fe₁₄B ribbon with the easy magnetization oriented normal to the ribbon plane.^[105] However, a higher wheel speed leads to an XRD pattern with isotropic characteristic, as shown in Fig. 13(b). The authors also found that post-annealing significantly improves the magnetic properties. Owing to this texture, the performance of the ribbons is comparable with that of the anisotropic sintered SmCo₅ magnets with μ₀M_r = 0.91 T, μ₀H_c = 1.62 T, and (BH)_max = 18.2 MGOe.

Fig. 13.

	Figure Option ViewDownloadNew Window
	Fig. 13. XRD patterns of SmCo5 ribbons prepared at (a) 5 m/s and (b) 25 m/s.[104]

The element substitutions in SmCo₅ ribbons influence both the microstructure and magnetic properties. Zhang et al. reported that Gd substation of Sm could effectively suppresses the formation of 2:7 phase,^[106] which usually exists in the anisotropic SmCo₅ ribbons without element substitution.^[107] The improved microstructure leads to the increase of both coercivity and squareness of the M–H curve significantly. A high remanence ratio of 0.95 has been attained in Sm_0.5Gd_0.5Co₅ ribbon with an energy product of 16 MGOe at room temperature. In addition, the temperature coefficient of remanence increases from −0.09% per °C for SmCo₅ to +0.02% per °C for Sm_0.5Gd_0.5Co₅ ribbons. This behavior could be attributed to the ferromagnetic coupling between HRE and Co atoms, which results in an increase of magnetization with increasing temperature in a given temperature range.

Enhanced saturated magnetization can be expected through Fe substation to Co, as Fe has larger magnetic moment than Co. Tung et al. found that M_s of as-spun SmCo₅ ribbons is only 0.42 T, while that of as-spun SmCo₃Fe₂ ribbons is 1.20 T.^[108] This huge enhancement cannot be only attributed the large magnetic moment of Fe. The TEM and energy dispersive analysis (EDX) analyses shows that a large portion of SmCo₃ grain exists in the SmCo₅ ribbons, which leads to low saturated magnetization in the ribbons. For the SmCo₃Fe₂ ribbons, the majority phase is Sm(Co, Fe)₅ with a small portion of Sm₂(Co, Fe)₇ and Sm(Co, Fe)₃ phases at the grain boundaries. In addition, XRD suggests the existence of α-Fe(Co) grains although it is not detected by TEM. Combining the grain refinement induced by the Fe substitution, a coercivity of 0.82 T and an energy product of 16 MGOe are obtained in the as-spun Sm(CoFe)₅ ribbons without post-annealing.

The nonmagnetic element substitution for Co has also been studied in the SmCo₅ ribbons. Yan et al. studied the effects of Cu and Sm doping on the magnetic properties of the oriented melt-spun ribbons with composition Sm_1+y(Co_1−xCu_x5, fabricated with a wheel speed of 5 m/s, followed by an optimized post-annealing.^[109,110] The room temperature coercivity is significantly enhanced to 2.1 and 1.24 T by Sm additions and Cu substitution, respectively. Interestingly, it is found that the coercivity mechanism of the Sm-doped and Cu-substituted ribbons are completely different, i.e., the Sm-doped ribbon has a nucleation-type while Cu-substituted ribbon has a domain-wall pinning type coercivity mechanism. In addition, compared with the Sm-doped ribbons, Cu-doped ribbons exhibit a better thermal stability of coercivity and a high coercivity of 0.52 T even at 300 °C. Fukuzaki et al. investigated the effect of nonmagnetic element M (M=Zr), V, Nb, Mo, and Ta) substitutions on magnetic properties and microstructure of Sm(Co_0.7Cu_0.3)₅Fe_0.54−xM_x ribbons.^[111] It is observed that the coercivity increases initially with increasing the substitution content, reaches a maximum value at x = 0.2–0.3, and then decreases for all studied element except Zr, as shown in Fig. 14. The high melting point of Mo and Ta prevents the diffusion of these substitution elements from grain boundaries to SmCo₅ grains, resulting to a high coercivity of 2.45 T in the Mo and Ta substituted ribbons.

Fig. 14.

	Figure Option ViewDownloadNew Window
	Fig. 14. Coercivity of Sm(Co0.7Cu0.3)5Fe0.54−xMx (M = Zr, Nb, Mo, Ta) ribbons as a function of substitution element content.[111]

Boron substitution of Co not only promotes amorphization in rapidly solidified Sm–Co alloys but also the formation of Sm–Co boride structures. The SmCo₄B compound, in which the B atoms replacing Co in the parent SmCo₅ structure, has a magnetocrystalline anisotropy even higher than SmCo₅.^[19] This makes SmCo₄B a promising candidate for nanocrystalline hard magnet.

Unfortunately, the Curie temperature of SmCo₄B is far lower than that of SmCo₅. Two methods are utilized in order to improve the Curie temperature. Saito et al. found that melt spinning can stabilize the disordered SmCo_5−xB_x structure even with .^[112] The value of T_c for the disordered compounds rapidly increases with decreasing B content from 197 °C at x = 1.0 to 557 °C at x = 0.6. Besides the adjustment of B content, Satio et al.^[112] and Gong et al.^[113] both observed the increase of Curie temperature with Fe substation for Co. Particularly, T_c of 472 °C is obtained in the SmCo₂Fe₂B ribbons. Therefore, there is a possibility of optimum hard-magnetic properties in metastable Sm(FeCo)_5−xB_x nanocrystalline alloys, though the realization of high-energy production is not likely to happen in these materials.

3.1.2.2. PrCo₅

Hexagonal PrCo₅ has comparable anisotropy field and saturated magnetization to those of SmCo₅, as shown in Table 7. It is the second choice among the RECo₅ compounds for permanent magnetic applications because of it large theoretical energy product of 36 MGOe and moderately high Curie temperature of 620 °C. However, the coercivity of PrCo₅ materials is rather low in comparison with the anisotropy field. A coercivity of only 0.6 T is achievable in the melt-spun PrCo₅ ribbons, as reported by Fuerst et al.^[114] This mainly arises from the existence of impurity Pr₂Co₁₇ and Pr₅Co₁₉ phases with low anisotropy field, which are generated by the eutectoid reaction at 846 °C, according to the Pr–Co phase diagram.^[115] Carbon addition could retard the formation of impurity phases and thus significantly improves coercivity, but decreases remanence. The highest coercivity obtained in the PrCo₅ and its carbon-doped ribbons is only 1.65 T with a rather low energy product of 7.4 MGOe so far.^[114] To understand the reason for the low coercivity and investigate the possibility to improve it, Zhang et al. designed an melt spinning experiment with a wheel speed up to 70 m/s.^[116] Figure 15 compares the TEM images of melt-spun ribbons prepared by 30 m/s and 70 m/s and the corresponding statistical grain size distributions. The 30 m/s prepared ribbons give two characteristic maxima in the grain-size distribution where the first maximum represents impurity phases Pr₂Co₁₇/Pr₅Co₁₉ with average grain size of about 110 nm and the second maximum originates from PrCo₅ phase with average grain size of about 290 nm. With increasing wheel speed to 70 m/s, only one maximum is observed in the grain size distribution, implying that the ribbons only consist of PrCo₅ phase with an average grain size of 19 nm. This indicates that the large increase of wheel speed fully impedes the formation of impurity phases and refines the nanostructure. Therefore the energy product of the 70 m/s ribbons increases almost twice compared to that of 30 m/s ribbons. A coercivity of 1.03 T and an energy product of 13.7 MGOe are achieved, which are the highest among the isotropic PrCo₅ ribbons. The coercivity value is close to that of nanocrystalline Pr₂Fe₁₄B, while the energy product is lower due to the lower saturated magnetization of PrCo₅ phase. However, PrCo₅ has much higher T_c than that of Pr₂Fe₁₄B, making melt-spun PrCo₅ ribbon a potentially good candidate for high-temperature application.

Fig. 15.

	Figure Option ViewDownloadNew Window
	Fig. 15. (color online) TEM images of nanocrystalline PrCo5 prepared at (a) 30 m/s, (b) 70 m/s, and statistical grain size distribution of (c) 30 m/s, (d) 70 m/s.[116]

It is interesting to know that the texture could be developed through die-upsetting method for the PrCo₅ ribbons, which is similar to that of the Nd–Fe–B ribbons through plastic hot deformation procedure. Fuerst and Brewer observed this behavior in both Pr–Co and Pr–Co–C nanocrystalline ribbons with PrCo₅ as the major phase.^[117,118] Although the texture is not as perfect as that in Nd–Fe–B magnets, it induces marked increase of remanence. The greatest alignment was obtained in the die-upset magnets with low-carbon compositions; in particular, Pr₁₈Co₈₁C produced a remanence of 0.87 T and an energy product of 16.9 MGOe.^[117] Through further partially substituting Sm with Pr, coercivities near ∼2.0 T were achieved in thermally quenched die-upset (Pr_0.6Sm_0.4)₁₈Co₈₁C magnets, while maintaining remanences above 0.8 T.^[118] However, for unknown reasons, the Sm–Co alloys with much higher coercivity do not develop texture during die upsetting.

1:7-type RECo₇ is a metastable phase in the RE–Co binary systems such as Sm–Co and Pr–Co. It has a TbCu₇ crystal structure which can be derived from RECo₅ when RE atoms are randomly substituted by Co pairs, confirmed by TEM study performed by Yu et al.^[119] It should be noted that RE₂Co₁₇ can also have the 1:7-type structure when they are prepared appropriately. A third doping element such as Ti, Zr, Cu, is helpful to stabilize the 1:7-type phase if they are used to substitute Co. Among the 1:7-type RE–Co alloys, SmCo₇ is the most studied materials. The SmCo₇ compound has saturation magnetization and Curie temperature lower than a Sm₂Co₁₇ but higher than a SmCo₅ compound. The anisotropy field of TbCu₇-type Sm–Co alloys is reported to be 10.5–14.0 T, which is about more than 1.6 times of the Sm₂Co₁₇ phase.^[120] In fact, the metastable SmCo₇ phase can be obtained by various methods, such as arc melting,^[121] mechanical alloying,^[122] and melt spinning which is summarized here.

Various additions were tried to stabilize the 1:7 phase and refine the microstructure of SmCo₇ ribbons. Guo et al. investigated the effect of transition metals (TM), such as Nb, Ta, Cr, and Mo, on the crystal structure and magnetic properties of melt-spun SmCo_7−xTM_x.^[123,124] It was found that a pure TbCu₇-type structure is formed in the SmCo_7−xNb_x and SmCo_7−xTa_x ribbons, while minor amount of impurity 2:17 phase appears for the that of Cr or Mo substituted ribbons, implying that the elements in VIB group (Cr and Mo) are inferior to the element in VB group (Nb and Ta) in stabilizing metastable 1:7 phase. The impact of phase impurity is also reflected by the coercivity in the TM-substituted ribbons. The Nb/Ta substitution leads to the coercivity enhancement of 0.83–0.94 T, while Cr/Mo only leads to an enhancement of 0.47–0.56 T. The effect of Si and Al substitution in SmCo₇ alloys has been investigated but the coercivity enhancement is much lower than that of Nb, Ta, Cr, and Mo.^[125] Recent study shows that the coercivity can be highly enhanced in the SmCo_7−xZr_x ribbons only if the wheel speed is high enough. Particularly, a high coercivity of 1.72 T is achieved in the SmCo_6.7Zr_0.3 ribbon with a fine grain structure around if the wheel speed is 60 m/s and above.^[126]

Carbon in combination with some transition metal element can produce fine 1:7 microstructures with further enhancement of coercivity. Du et al. found that carbon addition helps the formation of homogeneous microstructure with grain size around 60–70 nm in the Sm(Co_0.9Fe_0.16.8Zr_0.2C_x ribbons, which results in a coercivity of 1.43 T and an energy product of 11 MGOe.^[127] With an optimized Fe content, the energy product is improved to 13.7 MGOe in the Sm(Co_0.75Fe_0.256.7Zr_0.3C_0.06 ribbon, while coercivity is slightly dropped.^[128] Similar grain refinement and coercivity enhancement are observed in the NbC-doped ribbons as reported by Hsiao et al., confirming the role of carbon in the melt-spun 1:7-type alloys.^[129] In fact, Li et al. found that the grain size of the melt-spun SmCo_7−x(Cr₃C_2x ribbons decreases with the increase of Cr₃C₂ content.^[130] The amorphous state is achieved in the ribbons with when wheel speed is larger than 40 m/s. Magnetic properties can be further optimized by post-annealing of the amorphous ribbons.

Alternatively, boron can be doped into 1:7-type ribbons and significantly improves the coercivity. Makridis et al. reported that the as-spun Sm(Co_0.74−xFe_0.1Cu_0.12Zr_0.04B_x7.5 ribbons contains mainly TbCu₇-type phase with an average grain size of 60–100 nm.^[131] A record high coercivity of 3.85 T at room temperature is achieved in these ribbons. It should be noted that excessive amount of B may lead to deterioration of the hard-magnetic properties. Kim et al. found that the magnetically soft Sm₂Fe₁₄B phase with a planar anisotropy can be formed in the Sm₂(CoFeMn)₁₇B_x ribbons if B is over-doped.^[132]

TbCu₇ structure is observed in the Pr–Co alloys as well. Zhang et al. reported the stabilized 1:7 phase in both compound through arc melting^[133] and nanocrystalline powders with an average grain size of 20 nm through mechanical alloying of PrCo_7−xTi_x alloys with an energy product of 10.6 MGOe.^[134] With the assistance of C addition, Kostogorova et al. successfully fabricated the (Pr_xCo_1−x94Ti₃C₃ ribbons consisting of 1:7 phase with a coercivity around 1.0 T through melt spinning and post-annealing.^[135]

Besides microstructure refinement, the improvement of crystallographic texture in the melt-spun ribbons is also an effective way to promote the properties of 1:7-type magnets. Yan et al. observed a high degree of texture in the melt-spun Sm(CoFeCuZr)_7.4 ribbons prepared at low wheel speed such as 5 m/s.^[136] The formation of the crystallographic texture, with c-axis parallel to the longitudinal direction of ribbon, is attributed to a directional solidification process resulting from a temperature gradient. Maximum energy product of 12.3 MGOe is obtained in such highly-textured ribbons. Yan et al. also found that the C addition to Sm(CoZr)₇ alloy not only significantly enhances the coercivity of the ribbons but also affects the crystallographic texture. The preferred orientation changes from parallel to the ribbon plane for carbon-free ribbon, to normal to the ribbon plane for the SmCo_6.5Zr_0.5C_0.5 ribbon.^[137] Figure 16 shows the scanning electron micrograph (SEM) images of ribbons with and without carbon. The analysis shows that a dendrite structure was presented in the SmCoZr ribbon surface, while C addition leads to the diminishing of the dendrite and the grains become equiaxial. A higher energy product of 8.6 MGOe is observed in the C-added ribbon prepared at 5 m/s.

Fig. 16.

	Figure Option ViewDownloadNew Window
	Fig. 16. SEM micrographs on the surface of (a) SmCo6.5Zr0.5 and (b) SmCo6.5Zr0.5C0.5 ribbons.[137]

Since the Sm₂Co₁₇ and especially Sm₂(CoFe)₁₇ have saturated magnetization substantially higher than that of SmCo₅, it seems natural to increase (BH)_max in the Sm₂Co₁₇ magnets. Another advantage of Sm₂Co₁₇ is the higher Curie temperature compared to SmCo₅ phase, as seen in Table 7. Unfortunately, Chen et al. found that the coercivity of Sm₂(Co, Fe)₁₇ ribbons is less than 0.3 T, although they found that the coercivity increase with increasing wheel speed and a crystallographic texture formed during the melt-spinning process,^[138] which is similar to the behaviors observed in SmCo₅ ribbons. The low coercivity might be expected, since the anisotropy field of Sm₂Co₁₇ is only about one-fourth of that of SmCo₅. However, the major reason for the low coercivity appears to be related to a coarse dendritic microstructure with their long axis parallel to the longitudinal direction of the ribbons.^[139] The size and alignment of the dendritic grains can be significantly modified through Cu and Zr substation for Co. With an appropriate annealing, a coercivity of over 2.5 T has been obtained in the Cu and Zr substituted Sm₂(CoFe)₁₇ ribbons. However, these high coercivities cannot be attributed to the 2:17 solely. Instead, it should be attributed to the two-phase cellular structure, similar to that of conventional precipitation-type Sm(Co, Cu, Fe, Zr)_z magnets, which leads to the domain-wall pinning mechanism in the ribbons.^[139] On the other hand, Crabbe et al. studied the melt-spun Sm(Co_0.704Cu_0.061Fe_0.209Zr_0.025)_7.61 ribbons with a single-phase 2:17 microstructure and a grain size of about 30 nm, through controlling wheel speed and nozzle diameter during melt-spinning process. Unfortunately, their magnetic properties are still at the low end with μ₀M_r = 0.52–0.75 T, μ₀H_c = 0.3–0.5 T, (BH)_max = 4.4–10 MGOe.^[140] There are not many researches on the melt spinning of Sm₂(CoFe)₁₇ alloys with single-phase microstructure due to their poor hard magnetic properties and thus lack of potential applications. In fact, the melt-spun 2:17 ribbons with two-phase cellular microstructure have attracted interests, which will be discussed in Section 3.3.

Sm₂Fe₁₇ compound has high saturation magnetization but shows low Curie temperature and planar anisotropy likely due to the strong Fe sublattice anisotropy. However, interstitial modification of the Sm₂Fe₁₇ compounds can leads to uniaxial anisotropy and dramatically change their magnetic properties.^[99,141] Owing to a volume expansion of the atomic lattice caused by the absorbed N, C, or H atoms, the Curie temperature increases by more than 300 °C. In addition, both magnetocrystalline anisotropy and saturation magnetization also increase significantly, as shown in Table 7. For Sm₂Fe₁₇N_2.94, a value for the anisotropy field of 21 T at room temperature has been reported by Katter et al.^[99] These excellent magnetic properties were immediately realized in high-performance nanocrystalline magnets. However, these nitrides or carbides prepared by gas-phase interstitial modification are metastable, and decompose into Sm–N or Sm–C and α-Fe upon heating to about 700 °C and cannot be used for the development of anisotropic sintered magnets.

3.1.5.2. Sm₂Fe₁₇ carbide

The effect of interstitial carbon on Sm₂Fe₁₇ is similar to that of nitrogen. However, the number of the C atoms per formula unit after a solid–gas carburization usually does not exceed 2. This means a lesser volume expansion of the 2:17 lattice and thus lesser gain in the magnetic properties. As one can see in Table 7, both Curie temperature and anisotropy field of Sm–Fe–C are lower than those of Sm–Fe–N. However, there are several advantages of Sm–Fe–C comparing to Sm–Fe–N: i) the Sm₂Fe₁₇C_x system can be directly produced by the simple melt spinning without the two-step procedure in Sm–Fe–N system. ii) The melt-spun Sm–Fe–C ribbons have better phase thermal stability compared to the alloys prepared by solid–gas reaction. iii) The interstitial carbides Sm₂Fe₁₇C_x with higher carbon concentration can be formed by melt spinning, which leads to higher Curie temperature and larger magnetocrystalline anisotropy.

Shen and his group have systematically studied the structure and magnetic properties of the RE₂Fe₁₇C_x (RE=Dy, Er, Ho, Er, Ge, Tm, Y) ribbons prepared by melt spinning and other approaches.^[147–157] The Curie temperatures of RE₂Fe₁₇C_x are summarized in Fig. 17. A considerable enhancement of T_c is observed when x increases from 0 to 2.0. For small x-values, the Curie temperature increases with increasing carbon concentration more drastically than for larger x-values. In addition, there is a linear relation between T_c and the corresponding unit cell volumes, as shown in Fig. 17(b). It demonstrates the dependence of the exchange interaction on unit cell volume in the 2:17 structure. Figure 17(c) summarizes the dependence of room-temperature saturated magnetization on carbon concentration of RE₂Fe₁₇C_x ribbons. It was found that M_s increases with increasing x for low carbon concentration, while it becomes flat for high carbon concentration. The increase of M_s at lower carbon concentration is due to the enhancement of the Curie temperature which shifts the thermomagnetization curve to higher temperature. It should be noted that the substitution of Ga in Sm₂Fe₁₇ also increases Curie temperatures and develops a room temperature uniaxial anisotropy without introducing interstitial carbon or nitrogen atom.^[158] Unfortunately, no considerable coercivity is reported in the above mentioned Sm–Fe based ribbons.

Fig. 17.

	Figure Option ViewDownloadNew Window
	Fig. 17. (a) Curie temperatures of R2Fe17Cx ribbons (x = 0, 1.0, 1.5, and 2.0); (b) Curie temperatures versus unit cell volume of Er2Fe17Cx and Tm2Fe17Cx; (c) dependence of Ms on carbon content of R2Fe17Cx ribbons.[147–157]

Among all RE₂Fe₁₇C_x alloys, Sm₂Fe₁₇C_x is the best candidate for permanent magnets due to their high saturation magnetization and relatively high Curie temperature. It was found later by Shen et al. that substituting Fe partially with Ga, Al, Si, or Cr in Sm₂Fe₁₇C_x helps the formation of Sm–Fe compounds of even higher carbon content with the 2:17-type structure.^[159–161] The excellent magnetic properties of these 2:17-type carbides results from not only the interstitial carbon atoms but also the substitutional Ga, Al, or Si atoms, which leads to a significant increase in the uniaxial anisotropy field.

Figure 18 summarizes the intrinsic magnetic properties of Sm₂Fe₁₅M₂C_x (M = Ga, Al, Si) alloys, which were reported in Refs. [161–163] Similar to the contribution of carbon addition to Sm₂Fe₁₇C_x ribbons, carbon leads to the increases of Curie temperature in all Sm₂Fe₁₅M₂C_x alloys and reaches the maximum around x = 1.5∼2.0. The Ga, Al, and Si substitutions themselves also enhance the Curie temperatures, as shown in Fig. 18(a). Figure 18(b) presents the dependence of anisotropy field H_A on carbon content. It shows that carbon addition leads to dramatic increase of H_A as , and then increases slowly or even deteriorates with excess carbon addition. Unfortunately, carbon addition leads to the monotonic decrease of saturated magnetization for all Ga, Al and Si substitution, although it benefits the Sm₂Fe₁₇C_x alloys, as shown in Fig. 18(c). For this reason, the carbon content was usually chosen between x = 1.0–2.0 to gain the benefits of T_c and H_A but not scarify too much saturated magnetization. Comparing with Nd₂Fe₁₄B, the Sm₂Fe₁₅Ga₂C_x materials have higher T_c and H_A, but lower saturated magnetization.

Fig. 18.

	Figure Option ViewDownloadNew Window
	Fig. 18. Effect of carbon addition on (a) Tc, (b) HA, and (c) Ms of Sm2Fe17Cx and Sm2Fe15M2Cx alloys.[161–163]

With an optimized carbon content, Shen and his group further optimized the Ga, Al, and Si substitutions on intrinsic magnetic properties of Sm₂Fe₁₅M_xC_1.0-1.5 alloys,^[164,165] as shown in Fig. 19. It was found that maximum Curie temperature and anisotropy field are obtained in the alloys with x around 2, while they drop quickly with excess Ga/Al/Si substitutions. The saturated magnetization decrease almost linearly with the amount of substitution in a fast manner. Further substitution for Fe with transition metals such as Cu, Nb, V, Mo, Zr, Ti in Sm₂(Fe, TM)₁₅Ga₂C ribbons shows that they all deteriorates the Curie temperature, anisotropy field, and saturated magnetization.^[166] It should be noted that the Co substitution for Fe leads to monotonic increases of both Curie temperature and saturated magnetization in Sm₂Fe_15−xCo_xSi₂C ribbons if , while anisotropy field does not change much.^[167] Additional studies have also been conducted on the effect of Ga substitution on the intrinsic magnetic properties of R₂Fe_15−xGa_xC₂ (R = Gd, Y, Ho, Er) alloys, while similar behaviors as those in Sm–Fe–Ga–C system are observed.^{[164,168–171]}

Fig. 19.

	Figure Option ViewDownloadNew Window
	Fig. 19. The effect of Ga, Si, and Al substitutions on (a) Curie temperature, (b) anisotropy field, and (c) saturated magnetization of Sm2Fe15MxC1.0−1.5 alloys.[172,173]

Compared with Sm₂Fe₁₇C_x, the further advantage of Sm₂Fe₁₅M₂C materials is the developed coercivity and energy products. Zhang et al. studied the effects of wheel speed of melt spinning as well as the carbon content on the coercivity of the Sm₂Fe_14.5Cu_0.5Ga₂C_x ribbons,^[174] as shown in Fig. 20. The coercivity increases significantly with increasing carbon content and reaches a maximum at x = 1.0–1.5 for an optimized wheel speed of about 12.5–15 m/s. The maximum coercivity at room temperature is 2.2 T for Sm₂Fe_14.5Cu_0.5Ga₂C_1.5 ribbon. It should be noted that the coercivity shown in Fig. 20 is measured by a vibrating sample magnetometer (VSM) with a maximum applied field of 1.9 T, without reaching the saturation magnetization of the samples. The coercivity could be further improved to 2.6 T if the ribbon was annealed in a Sm atmosphere.^[175] With a small portion of Sm substituted by Pr, energy product of about 8.9 MGOe was obtained in the single-phase Sm–Fe–Ga–C ribbons.^[176] For Sm–Fe–Si–C ribbons, a coercivity about 1.5 T and an energy product of about 7.2 MGOe are achieved through Cu or Co substitution with optimized melt-spinning process.^[167,177] Such melt-spun Sm–Fe–(Ga, Si)–C ribbons usually have a fine microstructure with grain size of about 45–60 nm while the magnetic domain size is about 100–500 nm, indicating the effective exchange coupling between grains.

Fig. 20.

	Figure Option ViewDownloadNew Window
	Fig. 20. (color online) Dependence of coercivity on wheel speed of Sm2Fe15−xCuxGa2Cx (x = 0–2.0) and Sm2Fe15Ga2C ribbons.[174]

Sm(Fe, M)₁₂ (M = Ti, V, Mo, Si) compounds with the tetragonal ThMn₁₂ structure possess a uniaxial anisotropy at room temperature, which is large enough for the development of permanent magnet materials. Both melt spinning and ball milling were utilized to produce nanocrystalline Sm(Fe, M)₁₂ magnets. The SmFe_12−xM_x magnets with the desired 1:12 structure can be obtained if for Ti-substitution and if for V-substitution magnets. SmFe_12−xM_x is metastable phase which makes melt spinning suitable for development of such materials with nanostructures. Wang et al. studied the effects of V substitution on the magnetic properties of Sm–Fe–Ti–V melt ribbons. A coercivity of about 1.07 T is obtained in the melt-spun Sm₈Fe_75.5Ga_0.5Ti₈V₈ with a ThMn₁₂-type structure after post annealing, which is far higher than that of bulk magnets prepared by the powder metallurgy technique.^[178] A higher coercivity is reported in the SmFe₁₀TiV ribbons after annealing in Sm atmosphere, where such annealing technique is specially aimed to prevent the evaporation losses of Sm.^[179] Pinkerton et al. also reported a bulk Sm_0.89Fe₁₀V₂ magnet prepared by hot compaction of the overquenched ribbons with bulk magnetic properties: μ₀M_r = 0.54 T, μ₀H_c = 0.78 T, (BH)_max = 4 MGOe. With partial Co substitution for Fe, an energy product of about 8.4 MGOe is obtained in the Sm(Fe, Co, Ti)₁₂ ribbons.^[180]

Interstitial modification of the Nd(Fe, M)₁₂ compounds changes their intrinsic magnetic properties in a way very similar to that of Sm₂Fe₁₇. The room-temperature coercivity of about 0.6–0.9 T are obtained by nitriding the melt-spun Nd(Fe, Mo)₁₂ ribbons which only have coercivity less than 0.05 T prior to nitriding.^[181] The magnetic properties and interstitial atom effects in the R(Fe, M)₁₂ compounds have been reviewed in Ref. [182].

R₃(Fe, Ti)₂₉ is a novel intermetallic which crystallizes into a monoclinic structure with a space group of A2/m. The unit cell can be described as an alternate stacking of the tetragonal ThMn₁₂-type and the rhombohedral Th₂Zn₁₇-type structures. There is a strong enhancement in saturated magnetization and Curie temperature ( in the nitride) upon nitrogenation while the easy magnetization direction changes from planar to uniaxial. The anisotropy field about 12 T at room temperature is reported in the Sm₃(Fe, Ti)₂₉ nitrides.^[183] Chen et al. fabricated the Sm₃Fe_27.5Ti_1.5 ribbons via melt spinning with controlled cooled rate.^[184] It was found that 3:29 only forms at low wheel speed of 5–10 m/s, while 1:7 phase was developed at higher wheel speed of 20–25 m/s and amorphous phase became dominant with increased speed to 40 m/s. The best magnetic properties Sm₃Fe_27.5Ti_1.5 ribbons are μ₀M_r = 0.72 T, μ₀H_c = 0.15 T, (BH)_max = 2.9 MGOe, which is not very promising, even after optimized post-annealing and nitrogenation.

Nd₅Fe₁₇ phase was found to be a stable intermetallic compound with a hexagonal structure in the binary Nd–Fe system. Although the Nd₅Fe₁₇ phase has a high saturated magnetization of 1.52 T at 4 K and a Curie temperature of 230 °C, it does not possess c-axis anisotropic, which is essential for permanent magnet materials. On the other hand, Sm₅Fe₁₇ phase is a metastable phase with a Curie temperature around 280 °C and an anisotropy field . The single-phase Sm₅Fe₁₇ with two elements was first observed in the sputtered films with a coercivity of 1.47 T at room temperature.^[185] Although it was found later that Sm₅Fe₁₇ phase can be obtained through melt spinning,^[186] partial replacement of Fe by Ti and/or V could greatly favor the formation of the 5:17 phase. The ternaries around Sm₂₀Fe₇₀Ti₁₀ are the most studied alloys with the 5:17 structure. Saito et al. reported a coercivity of 4.0 T in the Sm₅Fe₁₇ ribbons consisted of grains around 50–100 nm.^[187] Unfortunately, the remanence of these ribbons is only about 0.48 T, which does not favor the permanent magnet applications.

Recent constraints on availability and cost increase of RE elements have stimulated a strong interest in the discovery of RE-free materials with excellent permanent-magnet properties. Many RE-free candidates with high magnetic anisotropy and high Curie temperature have been proposed and studied in the past years. Here we only briefly reviewed some of them such as MnBi, Co-rich Hf–Co/Zr–Co, Fe–Pt–B, Mn–Al–C and Alnico, which are prepared by melt spinning.

The low temperature phase (LTP) MnBi compound is of interest not only because of its RE-free feature, but also because of its magnetic properties. The magnetic anisotropy constant of MnBi is negative below −170 °C but increases with increasing temperature and attains its maximum value of 2.2 × 10⁶ J/m³ at 220 °C. Actually, the high temperature anisotropy constant of MnBi is much higher than that of Nd₂Fe₁₄B phase, although its room temperature value is lower. This made MnBi a good candidate for high temperature applications. However, the maximum energy product of MnBi at room temperature is much lower than most other permanent magnet materials, due to its low saturated magnetization. Guo et al. prepared the LTP MnBi by melt spinning and found a maximum anisotropy field of about 9 T at 257 °C and a Curie temperature of about 500 °C.^[188] Following the above method, Saha et al. produced the MnBi powders with a coercivity of about 2.6 T at 280 °C by a ball milling process following the melt-spinning.^[189] The saturated magnetization and remanence of these ribbons or powders are relative low. Yang et al. optimized melt spinning and annealing processes of the MnBi ribbons.^[190] They found that the grain size of 200 nm in the 20 m/s spun ribbons can be reduced to 20–30 nm with increasing wheel speed to 65 m/s. With a grinding process followed by alignment under magnetic field, an anisotropic nanocrystalline MnBi magnet is obtained with an energy product of 7.1 MGOe. The hysteresis loops of the anisotropic MnBi magnets are given in Fig. 21(a) along the directions parallel and perpendicular to the c-axis. The temperature dependence of the coercivity is shown in Fig. 21(b), which exhibits a positive temperature coefficient. The coercivity reaches a maximum of 2.5 T at 267 °C. The positive temperature coefficient of coercivity can be explained by the temperature dependence of the magnetocrystalline anisotropy field. It should be noted that MnBi compound is sensitive to oxygen and temperature. Cui et al. found that oxygen is the driving force for MnBi decomposition when subjected to a temperature higher than 200 °C. Once decomposed, the Mn reacts with oxygen forming MnO, causing irreversible loss of magnetic properties.^[191] The authors also reported an energy product of 11.9 MGOe for the MnBi magnet, although it is prepared by ball milling, not melt spinning. To improve the room temperature coercivity, Kharel et al. studied the effect of Pr substitution and found that the coercivity increases by a factor of 9 from 0.26 for Pr-free to 2.3 T for 9 at.% Pr substitution.^[192]

Fig. 21.

	Figure Option ViewDownloadNew Window
	Fig. 21. (color online) The magnetic properties of nanocrystalline MnBi. (a) The hysteresis loops of MnBi at room temperature along the directions parallel and perpendicular to the c-axis; (a) the temperature dependence of coercivity.[190]

Recently, the Co-rich transition-metal alloys such as HfCo₇ and Zr₂Co₁₁ have attracted interests to develop new RE-free magnetic materials with high-energy products. The metastable nature and/or the requirement of high formation temperature of about 1050–1230 °C for HfCo₇ phase indicate that melt spinning process is suitable to make such alloy with good phase purity. Although Zr₂Co₁₁ phase has a wider temperature range of about 300–1254 °C, a high-temperature growth or annealing process followed by melt spinning was generally used to obtain high-anisotropy crystal structures in the alloys. Balasubramanian et al. fabricated the nanocrystalline Co_100−xTM_x alloys (TM = Zr, Hf, and ) by melt spinning with varied wheel speed from 25 to 70 m/s.^[193] For the Zr–Co alloys, the grain size of the melt-spun ribbons is reduced from 120 to 50 nm with increasing wheel speed from 45 to 70 m/s. The ribbons majorly consist of both rhombohedral and orthorhombic Zr₂Co₁₁, with the minor cubic Zr₆Co₂₃ and Co phases. Although there are both magnetically hard and soft phases in the ribbons, the hysteresis loop of the 70 m/s melt-spun ribbon exhibits nearly a single-phase behavior and shows a coercivity of 0.24 T and an energy product of 4.3 MGOe. These properties have been slightly improved to 0.30 T and 4.6 MGOe by the same research group for coercivity and energy product, respectively.^[194] For the Hf–Co alloys, an energy product of 5.2 MGOe was achieved in the HfCo₇ ribbons. Besides the HfCo₇ phase, McGuire et al. reported a Hf₂Co₁₁B-type nanocrystalline by melt spinning with the best magnetic properties of μ₀M_r = 0.62 T, μ₀H_c = 0.45 T, and (BH)_max = 6.7 MGOe.^[195]

Several RE-free magnets have L1₀-type structure, such as Fe–Pt and Mn–Al. Zhang et al. studied the phase evolution, thermal stability and magnetic properties of the melt-spun Fe_70−xPt_xB₃₀ ribbons.^[196] It was found that the amorphous forming ability increases with small amount addition of Pt into the Fe₇₀B₃₀ alloy and a complete amorphous phase is obtained in the alloys with x = 5 and 10 and wheel speed of 33 m/s. However, further increase of Pt content results in the formation of fcc-FePt, Fe₂B with or without L1₀-FePt phases. A homogenous microstructure, consisting of magnetically hard L1₀-FePt and soft Fe₂B phases with an average grain size of 30–50 nm, is formed in the ribbons with x = 10–20 after appropriate annealing. The coercivity of Fe_70−xPt_xB₃₀ ribbons increases while remanence decreases with increasing Pt content. Good magnetic properties were obtained in the Fe_70−xPt_xB₃₀ ribbons with μ₀M_r = 0.81–1.20 T, μ₀H_c = 0.22–0.65 T, and (BH)_max = 8.0–11.0 MGOe.

Liu et al. studied the effects of C, B, and RE doping in phase transitions and magnetic properties of Mn–Al alloy.^[197] The as-spun Mn–Al, Mn–Al–C, and Mn–Al–C–RE ribbons possessed a hexagonal ε phase. The ε to L1₀ phase transition occurs at ∼500 °C, which is of interest as it could develop hard magnetic properties in the ribbons. Moderate carbon addition promotes the formation of the desired magnetically hard L1₀ phase and improves the hard magnetic properties. The best magnetic properties, μ₀M_r = 0.3 T, μ₀H_c = 0.15 T, and (BH)_max = 1.5 MGOe, were obtained in Mn–Al–C ribbon with C content of 1.7 at.% annealed at 650 °C, though they are quite low compared to other permanent magnets.

Löwe et al. studied the magnetic properties of melt-spun Alnico ribbons.^[198] The grain size of the ribbons decreases from to about by increasing the wheel speed from 5 m/s to 60 m/s, and further decreases by addition of 1 at.% B to less than . The size of the spinodal features was about the same in the ribbons without B, which means that grain size above a certain threshold does not influence the spinodal decomposition anymore. The coercivities obtained in all samples are about 0.04 T, only half of those in anisotropic bulk samples, even though the aspect ratio of spinodals is the same as in bulk.

Coehoorn et al. first obtained this novel nanocomposite magnet consisting mainly of Fe₃B with Nd₂Fe₁₄B as a secondary phase.^[199] The amorphous alloy with composition of Nd₄Fe₇₈B₁₈ was produced by melt spinning, followed by post annealing at 670 °C for 30 minutes. The annealing process consists of two steps: the metastable phase Fe₃B is crystallized with a small amount of α-Fe phase at a lower temperature at first, while the hard-magnetic phase Nd₂Fe₁₄B is formed at a higher temperature. The annealed ribbons consist of 73% Fe₃B, 12% α-Fe, and only 15% magnetically hard phase Nd₂Fe₁₄B. The saturated magnetization of Fe₃B and Nd₂Fe₁₄B are both 1.6 T at room temperature, while Fe₃B has a higher Curie temperature of about 513 °C, nearly 200 °C higher than that of Nd₂Fe₁₄B. The annealed ribbon gives a coercivity of 0.3 T, which may be due to the presence of Nd₂Fe₁₄B since Fe₃B has much lower magnetocrystalline anisotropy. The most remarkable property in this material is the normalized remanence, which is as high as 0.75 and much higher than 0.5 predicted by Stoner–Wohlfarth model. The large remanence enhancement is attributed to the fine microstructure with Fe₃B grains of 30 nm and Nd₂Fe₁₄B grains of 10 nm plus the exchange coupling between them. Owing to the large remanence, the energy product of 11.9 MGOe is obtained.^[200] With partial substitution of Nd by Tb or Dy, the coercivity of the melt-spun ribbons can be improved to 0.45 T.^[201] Simultaneously, Shen et al. studied the amorphization and crystallization of the rapidly quenched M_1−x(Nd₂Fe₁₄B)_x ribbons (M = Fe₃B, Fe₂B; ).^[203] By the analyses of phase components and magnetic properties, it was found that the Fe₃B-based materials have superior properties than the Fe₂B-based materials. A coercivity of 0.3 T and an energy product of 13 MGOe in Nd₄Fe_77.5B_18.5 alloys are achieved at room temperature by annealing at 670 °C for a short time.

Kim et al. studied the ferromagnetic phase diagram of the Nd_xFe_yB_z ( , , ) melt-spun alloys.^[206] Most of the as-spun ribbons prepared at 35 m/s are found to contain amorphous phase. Depending on the composition, the main phases of the optimally annealed ribbons are found to possess one or more the following phase: α-Fe, Fe₃B, Nd₂Fe₂₃B₃, and Nd₂Fe₁₄B phases. It was found that the Fe₃B phase precipitates only appears in the alloys with high boron content and low Nd content. Figure 23(a) summarizes the ferromagnetic phase diagram of the optimally annealed Nd_xFe_yB_z melt-spun alloys: (i) for Nd-rich alloys with , Nd₂Fe₁₄B appears to be the main phase; (ii) for Fe-rich alloys with and , the major phase is α-Fe with Nd₂Fe₁₄B as a secondary phase; (iii) for B-rich alloys with , , and , the major phase is Fe₃B with α-Fe and Nd₂Fe₁₄B as minor phases; (iv) for the Nd₈Fe₇₇B₁₅ alloy, the major phase is Nd₂Fe₂₃B₃ with minor phases of Fe₃B and α-Fe, showing no evidence of the presence of Nd₂Fe₁₄B; (v) for the Nd₁₁Fe₇₇B₁₂ alloy, the major phase is Nd₂Fe₁₄B with precipitation of Fe₂₃B₆ and α-Fe. The coercivity of alloys with different phase components has a wide distribution, strongly depending on the Nd content, as shown in Fig. 23(b). A high coercivity of about 2.9 T was observed in the Nd₁₃Fe₇₇B₁₀ alloy, while it dramatically dropped to 0.01 T due to the formation of magnetically soft phase Nd₂Fe₂₃B₃ instead of Nd₂Fe₁₄B phase in the Nd₈Fe₇₇B₁₅ alloy.

Fig. 23.

	Figure Option ViewDownloadNew Window
	Fig. 23. (a) Ferromagnetic phase diagram and (b) coercivity contour of the optimally annealed NdxFeyBz ( , , ) melt-spun alloys. It is noted that the unit of coercivity is kOe, where 1 kOe = 0.1 T.[206]

The substitutions for either rare earth or transitional metal elements could affect the magnetic properties of melt-spun ribbons. Cheng et al. reported that the substitution of R (R=Y, Pr, Gd, or Dy) for Nd in R_xNd_4−xFe_77.5B_18.5 alloys results in the decrease of coercivity with Y, Pr, Gd substitution.^[172,207] A small amount of Dy addition improves the coercivity but full substitution of Dy also decrease coercivity due to the presence of magnetically soft phase Dy₃Fe₆₂B₁₄. It was interesting that the melt-spun Nd–Fe–B alloys annealed at 670 °C for 2 minutes consist of only body-centered-tetragonal (bct) Fe₃B and a few percent of α-Fe but no Nd₂Fe₁₄B phase. In addition, about 5 at.% Fe atoms in the Fe₃B phase are replaced by Nd in the annealed ribbons, which enhanced the hard-magnetic properties of Fe₃B. It was further found that the substitutions of Pr, Gd, and Dy do not influence the hyperfine field for ⁵⁷Fe in α-Fe, but increase the hyperfine field for ¹¹B in bct-Fe₃B. This leads to a conclusion that the hard-magnetic properties of melt-spun Nd–Fe–B result from bct-Fe₃B containing Nd atoms, but not from the presence of the magnetically hard Nd₂Fe_l4B phase.^[207] Kanekiyo et al. further improved the magnetic properties of the Fe₃B/Nd₂Fe₁₄B nanocomposite by Dy addition with remanence of 1.15 T, coercivity of 0.5 T and energy product of 16.5 MGOe.^[208]

Xiao et al. studied the effects of grain size and anisotropy of magnetically hard phase on the magnetic properties of B-rich Nd_4−xSm_xFe_77.5B_18.5 (x = 0–0.5) alloys through annealing control and Sm substitution for Nd.^[209] It was found that the grain sizes of the soft phase Fe₃B is about 16 nm and 23 nm for 670 °C and 760 °C annealed ribbons, respectively. These grain sizes are larger than the critical size of about 10 nm for magnetically soft phase, predicted by the Kneller–Hawig model.^[13] However, effective exchange coupling was observed in the Sm-substituted ribbons, which indicates that the Sm substitution reduced the anisotropy constant of magnetically hard phase and thus caused the extension of critical length for magnetically soft phase.^[209]

Shen et al. studied the Co substitution for Fe in the Nd₄(Fe_1−xCo_x)_77.5B_18.5 alloys and found that the Curie temperatures of both the as-quenched amorphous and annealed ribbons increase with Co content.^[204] Interestingly, it was observed that the magnetic moment of Fe atom increases with increasing Co content from x = 0 to x = 0.4, while the magnetic moment of Co atom keeps a constant value over the whole concentration range. Unfortunately, the coercivity decreases monotonically with increase of Co concentration. A still good magnetic property was observed in the ribbons with , which have coercivity of 0.23–0.30 T and energy products of 11.6–13.5 MGOe. The benefits of Co addition on the magnetic properties of the Fe₃B-based nanocomposite were confirmed in Nd₄Fe_77.5−xCo_x(HfGa)_yB_18.5 ( , y = 0, 0.5) ribbons.^[210] The Co addition retarded the crystallization of both α-Fe and Fe₃B phases, but accelerated that of Nd₂Fe₁₄B phase. As a result, the coercivity increased slightly with the addition of Co and an energy product of about 14.4 MGOe was obtained.

Kanekiyo et al. have reported that Al, Si, Cu, Ga, Ag, Au additives considerably reduce grain size of (Fe,Co)₃B/Nd₂Fe₁₄B nanocomposite based on Nd₅Fe_71.5Co₅B_18.5M alloy, where M is the above-mentioned elements.^[211] For example, the optimally heat-treated ribbons of Nd₅Fe_76.5B_18.5 have grains around 40–50 nm, while the Ga or Al combined with Co addition could reduce the grain size down to 10–20 nm in the Nd₅Fe_71.5Co₅B_18.5M ribbon. The grain refinement increases the remanence from 1.0 T for Nd₅Fe_76.5B_18.5 ribbon to 1.08–1.20 T for the Nd₅Fe_71.5Co₅B_18.5M ribbons, where Si additive gives the highest value. The Cu additive provides the lowest remanence among the M-added samples, but it generates the largest coercivity around 0.45 T. On the other hand, Ga-added ribbons have the largest energy product of about 16 MGOe. Ping et al. found that Cu and Nb can refine the microstructure of the Fe₃B/Nd₂Fe₁₄B using TEM and atom probe field ion microscopy (APFIM).^[212] Figure 24 shows the TEM micrographs of alloys with and without Cu and Nb additions. This indicates that Cu addition to the ternary alloy causes significant reduction of the grain size of the optimum nanocomposite microstructure. The combination of Cu and Nb can reduce the grain size down to 12 nm, where Cu is dissolved in the Nd₂Fe₁₄B phase. However, Nb mainly exists in the Fe₂₃B₆ phase. The crystallization process in Nd_4.5Fe_75.8B_18.5Cu_0.2Nb₁ ribbons are: amorphous Fe₃B + amorphous Fe₃B + Nd₂Fe₁₄B + Fe₂₃B₆. An even higher remanence about 1.25 T is obtained in these ribbons with an energy product of about 15.6 MGOe. Rajasekhar et al. studied the Mn substitution of Fe in the Nd_4.5Fe_77−xMn_xB_18.5 (x = 0–2) nanocomposite magnets and found that 1 at.% Mn addition could improve both coercivity and the normalized remanence.^[213]

Fig. 24.

	Figure Option ViewDownloadNew Window
	Fig. 24. TEM bright field micrographs of (a) Nd4.5Fe77B18.5, (b) Nd4.5Fe76.8B18.5Cu0.2, (c) Nd4.5Fe B18.5Cu0.2Nb1, and (d) Nd4.5Fe76B18.5Nb1 alloys annealed at 660 °C for 10 min. The average grain size is about 30, 17, 12, and 43 nm, respectively.[212]

Cheng et al. reported that the carbon substitution for B reduces the crystallization temperature of the as-spun Nd₄Fe_77.5B_18.5 amorphous alloy.^[214] Unfortunately, the coercivity of the crystallized ribbons decreases with increasing carbon content, with a simultaneous development of an orthorhombic Fe₃B (o-Fe₃B) phase, instead of the bct-Fe₃B phase.

Although researches still occasionally occur recently, it is concluded that the energy product of Fe₃B/Nd₂Fe₁₄B nanocomposite magnets is relatively low due to the relatively low saturated magnetization of Fe₃B phase. Their coercivity is relative low comparing to other nanocomposites, so it is not critical to understand where is the coercivity from, either the bct-Fe₃B phase or Nd₂Fe₁₄B phase. The magnetic properties of some typical Fe₃B/Nd₂Fe₁₄B nanocomposite magnets prepared by melt spinning are summarized in Table 8.

Table 8.

Table 8.

Table 8. μ0Mr, μ0Hc, and (BH)max of Fe3B/Nd2Fe14B nanocomposite magnets with different compositions. . Composition μ0Mr/T μ0Hc/T (BH)max/MGOe Reference Nd4Fe78B18 1.20 0.40 11.9 [200] Nd4Fe77.5B18.5 1.20 0.30 13.0 [203] Nd4(Fe0.94Co0.06)77.5B18.5 1.20 0.3 13.5 [204] Nd4Fe77.5−xCoxHf0.5Ga0.5B18.5 1.15 0.35 14.4 [210] Nd4.5Fe73Co3Ga1B18.5 1.20 0.42 16.0 [211] Nd3.5Dy1Fe73Co3Ga1B18.5 1.15 0.50 16.5 [208] Nd4.5Fe75.8B18.5Cu0.2Nb1 1.25 0.34 15.6 [212] Nd4.5Fe75.8B18.5Cu0.2Nb1 1.25 0.34 15.6 [212]

Table 8. μ0Mr, μ0Hc, and (BH)max of Fe3B/Nd2Fe14B nanocomposite magnets with different compositions. .

The α-Fe phase is actually often observed in the Fe₃B/Nd₂Fe₁₄B nanocomposite as mentioned in Section 3.2.1. Since α-Fe has a higher saturated magnetization (about 2.2 T) than those of Fe₃B and Nd₂Fe₁₄B phases (both about 1.6 T), it is expected that Nd₂Fe₁₄B/α-Fe magnets have higher remanence and thus higher energy product than those of Nd₂Fe₁₄B/Fe₃B magnets. Moreover, Nd₂Fe₁₄B/α-Fe magnets also have higher coercivity than Nd₂Fe₁₄B/Fe₃B. Therefore, Nd₂Fe₁₄B/α-Fe magnet is superior to Nd₂Fe₁₄B/Fe₃B for hard magnetic applications. It is true that most of investigations on exchange-coupled nanocomposites are based on Nd₂Fe₁₄B/α-Fe magnets recently. To form the nanocomposite magnet with α-Fe as the magnetically soft phase instead of Fe₃B, the B content is usually controlled to be about or less than 10 at.%, which is much lower than 18–20 at.% of the Nd₂Fe₁₄B/Fe₃B magnets. However, the low B content causes one shortcoming, i.e., the difficult in controlling the microstructure of the Nd₂Fe₁₄B/α-Fe because of the low glass forming ability in such alloys. Because of this, the production of Nd₂Fe₁₄B/α-Fe nanocomposite materials in an industrial scale with good reproducibility is challenging. The elemental additives are usually used in the production in order to control the microstructure, which is a major topic of this type of magnets.

The microstructure control, such as grain size, is one of the key factors to achieve high remanence and coercivity in the nanocrystalline and nanocomposite magnets. Davies et al. studied the effects of grain size on the remanence and coercivity of the single-phase nanocrystalline Nd_13.2Fe_79.6B₆Si_1.2 magnets.^[215] It is observed that remanence enhancement begins when grain size falls below ∼30 nm for both surfaces of the melt-spun ribbons, i.e., roller and free surfaces. The phenomenon is believed to result from exchange coupling between grains when they are free of non-magnetic intergranular phase. This condition correspondingly leads to reduction of coercivity since it lowers the resistance to reverse magnetization, as shown in Fig. 25(a). Based on the controlled grain growth, Davis et al. first reported the formation of Nd₂Fe₁₄B/α-Fe nanocomposites from the Nd_8+xFe_86−xB₆ alloys by varying the alloy composition and processing conditions of the melt spinning.^[215] It is found that the remanence is consistently above 0.8 T even for the stoichiometric composition Nd₁₂Fe₈₂B₆ and increases monotonously to for Nd₈Fe₈₆B₆ with decreasing Nd content. The two-phase structure forms in the alloys of Nd content less than 11 at% and Nd₂Fe₁₄B grains smaller than 30 nm interspersed with smaller isolated α-Fe grains typically ∼10 nm. Thus, the high remanence for these alloys is believed to arise from a combination of the remanence enhancement observed in single-phase Nd₂Fe₁₄B nanocrystalline and an increase in saturated magnetization due to the presence of α-Fe, whose volume fraction increases with decreasing Nd content. As remanence increases, the coercivity decreases. However, it should be noted that this decrease is only gradual instead of collapse because of the exchange coupling between the magnetically hard and soft phases. Therefore, high energy products around 20 MGOe with coercivity of about 0.6 T are obtained for the isotropic nanocrystalline ribbons with Nd less than 9.5 at%. The dependences of remanence, coercivity and energy product on Nd content are given in Fig. 25(b). Similar results are also reported in the work from same research group.^[216]

Fig. 25.

	Figure Option ViewDownloadNew Window
	Fig. 25. (color online) (a) Dependence of remanence and coercivity on grain size of melt-spun Nd13.2Fe79.6B6Si1.2 ribbons. (b) Dependence of magnetic properties on Nd content of melt-spun Nd8+xFe86−xB6 ribbons. The figures are reproduced from Ref. [215].

Besides the grain size control, the soft-phase component is another key factor to achieve good magnetic properties in nanocomposite Nd₂Fe₁₄B/α-Fe magnets. Bauer et al. studied the effect of α-Fe phase content (up to 40 vol.%) on magnetic properties of the composite magnets, starting from nearly single-phase Nd₂Fe₁₄B magnets.^[73] The maximum remanence μ₀M_r = 1.25 T with μ₀M_c = 0.53 T and (BH)_max = 23.2 MGOe were achieved in nanocomposite magnets containing 30 vol.% α-Fe (as shown in Fig. 26). It should be noted that demagnetization factor N of the ribbons, which is on the order of 0.01, was calculated based on the dimensions of the flakes. A representative TEM micrograph of a nanocomposite Nd₁₀Fe₈₄B₆ magnet with 14.2 vol% α-Fe is given in Fig. 27. The microstructure analysis reveals two characteristic maxima in the grain distribution at about 15 nm and 25 nm corresponding to the α-Fe phase and to the Nd₂Fe₁₄B phase, respectively. Further high-resolution TEM analysis confirmed that smaller and larger grains represent the α-Fe and Nd₂Fe₁₄B phases, respectively.

Fig. 26.

	Figure Option ViewDownloadNew Window
	Fig. 26. (color online) Dependence of magnetic properties of NdFeB-based nanocomposite magnets on α-Fe content at room temperature. The figure is reproduced from Ref. [73].

Fig. 27.

	Figure Option ViewDownloadNew Window
	Fig. 27. (a) TEM micrograph and (b) grain size distribution of melt-spun Nd10Fe84B6 ribbon with 14.2 vol.% α-Fe.[73]

The element substitutions in the Nd₂Fe₁₄B/α-Fe affect the microstructure and thus magnetic properties, which are categorized into RE substitution, refractory element substitution and other element substitution, as summarized below. The magnetic properties of the element substituted melt-spun ribbons are also summarized in Table 9.

Table 9.

Table 9.

Table 9. μ0Mr, μ0Hc, and (BH)max of Nd2Fe14B/α-Fe nanocomposite magnets with different compositions and different element substitutions. . Composition μ0Mr/T μ0Hc/T (BH)max/MGOe Reference Nd9.5Fe84.5B6 1.10 0.6 20 [215] Nd12Fe82B6 (stoichiometric) 1.02 1.05 21.2 [73] Nd8Fe87B5 1.25 0.53 23.2 (Pr0.75Nd0.25)2Fe14B 1.05 1.01 23 [217] (Nd0.4Pr0.6)9Fe86B5 1.14 0.62 19 [218] Nd8.16Dy1Fe85.26Nb1B4.58 1.02 0.88 16.8 [219] Nd10Fe77.5TiCB10.5 0.92 1.21 16.2 Nd10Fe77.5NbCB10.5 0.85 1.50 1.45 Nd10Fe77.5VCB10.5 0.87 1.04 14.3 [221] Nd10Fe77.5CrCB10.5 0.92 0.99 13.2 Nd10Fe77.5MoCB10.5 0.81 0.99 11.6 Nd10Nb2Fe82B6 0.99 1.18 20 [223] Nd9Dy0.4Fe76.24Ti5Nb0.36Co3B6 0.91 1.20 21.9 [224] Nd9.5Fe82.5B6.5Ta1.5 0.94 0.94 19.8 [225] Nd10Fe80Zr3B6C1 1.00 0.94 18.1 [226] Nd10Fe83B6In1 0.83 0.58 18 [228] (NdDy)11.5Fe79.5Nb1Zn2B6 1.14 0.83 20.6 [229]

Table 9. μ0Mr, μ0Hc, and (BH)max of Nd2Fe14B/α-Fe nanocomposite magnets with different compositions and different element substitutions. .

RE Psubstitution. Betancourt et al. studied the Nd: Pr ratio on the magnetic properties of (NdPr)_xFe_94−xB₆ ribbons.^[217] Remanence enhancement was observed for all compositions studied, arising from intergrain exchange coupling between the (NdPr)₂Fe₁₄B crystallites with an average grain size of ∼40 nm and α-Fe grains. An excellent combination of (BH)_max about 23 MGOe and coercivity about 1.0 T was observed for the (Pr_0.75Nd_0.25)₂Fe₁₄B alloy. Zhang et al. reported a similar enhancement of magnetic properties through Pr substitution for Nd in the (Nd_1−xPr_x)₉Fe₈₆B₅ ribbons.^[218] The best magnetic properties with remanence of 1.14 T, coercivity of 0.62 T, and energy product of 19 MGOe are achieved in the sample with x = 0.6. Wang et al. reported that 1 at.% Dy substitution for Nd could improve coercivity of Nd₂Fe₁₄B/α-Fe nanocomposite magnets, which is understandable because Dy₂Fe₁₄B has higher magnetocrystalline anisotropy than that of Nd₂Fe₁₄B.^[219] With optimized composition and annealing conditions, the magnetic properties of remanence, coercivity, and energy product are 1.02 T, 0.88 T, and 16.8 MGOe in Nd_8.16Dy₁Fe_85.26Nb₁B_4.58 alloy. Because it is believed that La is helpful to improve the glass forming ability due to its large atomic size and thus to reduce the optimum wheel speed required to prepare Nd₂Fe₁₄B/α-Fe nanocomposites, Change et al. studied the effect of La substitution for Nd on the microstructure structure and magnetic properties of Nd₂Fe₁₄B/α-Fe magnets.^[220] Unfortunately, it was found that the grain size increases from 20–30 nm for both phases of the N_9.5Fe_80.5B₁₀ ribbons to 70–80 and 30 nm of the (Nd_0.95La_0.05)_9.5Fe_78.5Cr₂B₁₀, respectively. Interestingly, a further increase in total rare earth content to (Nd_0.95La_0.05)₁₁Fe₇₇Cr₂B₁₀ resulted in fine grains with average grain size around 20–30 nm again. A remanence of 0.92–0.97 T, coercivity of 1.11–1.32 T, and energy product of 16.5–18.0 were obtained in the La substitution ribbons.

Refractory elements The melting points of the refractory metals (Cr, Ti, Nb, V, Mo, Zr, Hf, Ta, Mn, and W) are the highest for all element except C, Os, and Ir. It was found that the refractory element substitution for Fe can effectively reduce the grain size of the Nd₂Fe₁₄B/α-Fe nanocomposite as they are not easy to be dissolved in both Nd₂Fe₁₄B and α-Fe phase. In other words, these elements usually stay at the grain boundaries and thus improve the coercivity of the magnets. The effect of Nb on the microstructure of Nd₂Fe₁₄B/α-Fe nanocomposites has been widely studied and it has been approved as one of the effective elements to improve the magnetic properties. Chang et al. have confirmed that Nb is the best element among Ti, Nb, V, Cr, and Mo to improve the coercivity of Nd₁₀Fe_balMCB_10.5 ribbons.^[221] The Nb-substituted ribbons have a coercivity of 1.5 T, which is much larger than the second best of Ti-substituted ribbons with a coercivity of 1.2 T. The coercivity enhancement through Nb-substitution is again confirmed by Chen et al. in the Nd₁₂Fe_82−xNb_xB₆ (x = 0, 1.5, and 3) ribbons.^[222] A coercivity of 1.24 T in ribbons with 1.5 at.% Nb is much larger than that of Nb-free ribbons with a coercivity of 0.86 T. In addition, the thermal stability is significantly improved with Nb substitution. As compared with the ternary Nd₁₂Fe₈₂B₆, the temperature coefficient of coercivity is reduced from −0.36%/°C to −0.32%/°C. TEM and energy dispersive spectroscopy (EDS) analyses show that Nb atoms are expelled during melt-spinning and annealing processes, resulting in the enrichment of Nb atoms in the grain boundary region and the generation of Nb-rich phase along grain boundary, which inhibits the grain growth and finally refines the microstructure. The enrichment of Nb atoms at the grain boundary may weaken the exchange coupling between the phases in the above magnets, which causes a higher coercivity but lower remanence. As a result, the Nb-substituted nanocomposite magnets do not have best energy products. Recently, Zhang et al. reported that enhanced remanence and energy product could be achieved in Nb-substituted magnets if the microstructure are fine controlled in the melt-spun ribbons.^[223] Figure 28 compares the TEM images of the Nd₁₂Fe₈₂B₆ and Nd₁₀Nb₂Fe₈₂B₆ ribbons spun at 15 m/s. The grain size decreases from 120 nm for Nd₁₂Fe₈₂B₆ to 20 nm for Nd₁₀Nb₂Fe₈₂B₆. The grain refinement promotes intergrain exchange coupling and leads to the remanence enhancement from 0.84 T to 0.99 T and energy products increase from 15 MGOe to 20 MGOe. It should be noted that Ti substitution has similar function as Nb. Chang et al. reported that Ti substitution in the Nd₂Fe₁₄B/α-Fe nanocomposite ribbons could result in an even finer grain structure than that of Nb substation, i.e., average grain size of about 10–20 nm is achievable.^[221] Therefore, a higher energy product is achieved in the Ti-substitution nanocomposites, although it accompanies a slightly lower coercivity. Zhang et al. investigated the effect of Ti content on magnetic properties of Nd₉Dy_0.4Fe_81.24−xTi_xNb_0.36Co₃B₆ ribbons and found that the coercivity and energy product increases from 0.74 T and 15.8 MGOe to 1.20 T and 21.9 MGOe, respectively.^[224] Besides Nb and Ti, Ta is another element proved to be an effective element to improve microstructure and magnetic properties of the nanocomposite magnets. Zeng et al. found that the average grain size of 30–40 nm of Ta-free Nd_9.5Fe₈₄B_6.5 ribbons can be reduced to 20–25 nm with 1.5 at.% Ta substitution.^[225] Although the remanence is reduced from 1.07 T to 0.94, the coercivity increases from 0.64 T to 0.94 T. The refined microstructure improves the intergrain exchange coupling and leads to an energy product about 19.8 MGOe. Hashino et al. have found that Zr addition improves the glass forming ability and retards the formation of α-Fe phase in the Nd₁₀Fe_83−xZr_xB₆C₁ ribbons, which results in a finer microstructure in the melt-spun ribbons.^[226] Wang et al. confirmed that a small amount of Zr addition can inhibit the prior precipitation of the magnetically soft α-phase and refine the microstructure.^[36] A combination of Zr and Nb could further improve the maximum energy product.

Fig. 28.

	Figure Option ViewDownloadNew Window
	Fig. 28. TEM images for the (a) Nd12Fe82B6 and (b) Nd10Nb2Fe82B6 ribbons spun at 15 m/s.[223]

Other element substitutions Liu and Davis studied the substitution of Fe by Co and found that Co improves remanence, coercivity, and Curie temperature in certain concentration range. The 10–20 at.% Co substitution could lead to the best magnetic properties with μ₀M_r = 0.97 T and μ₀H_c = 0.55 T in the nanocomposite Nd₂Fe₁₄B/α-Fe alloys produced by melt spinning.^[227] Ji et al. found that 1 at% indium (In) addition in the Nd₁₀Fe_84−xB₆In_x nanocomposites significantly enhances the hard magnetic properties.^[228] The coercivity increases from 0.5 T to 0.6 T, while remanence ratio also increases from 0.71 to 0.83, accompanying improvement of squareness of hysteresis loop, indicating a strong inter-phase exchange coupling in the nanocomposite magnets. Therefore, the maximum magnetic energy product increases remarkably from 12 to 18 MGOe. The origin for these enhancements is mainly attributed to the magnetically softened grain boundaries and enhanced crystallographic coherency by indium addition. Ma et al. found Zn addition could help on developing texture in the hot-deformed (NdDy)_11.5Fe_82.5−xNb_xB₆ magnets, because the combination of Nb and Zn promotes the volume fraction of the intergranular phase where Zn prefers to diffuse into.^[229] The combination of Nb and Zn addition increases the coercivity of alloys up to 22% and a maximum energy product of 20.6 MGOe is obtained.

It is interesting to note that the magnetic field during melt spinning process may also reduce the grain size in the ribbons. Nguyen et al. customized a melt spinner which could provide a magnetic field up to 0.32 T parallel or perpendicular to the ribbons plane during the melt-spinning process.^[230,231] The grain refinement effect was proved by TEM observations in the single-phase Nd₁₅Fe₇₇B₈ ribbons, as shown in Fig. 29. The average grain size of 22 nm for the ribbons prepared under zero field is reduced to 19 nm for the one prepared under a magnetic field of 0.15 T with the field direction perpendicular to the ribbons plane. In addition, the morphologies of the two ribbons are also different. It is observed that the separated small grains are embedded in a matrix with dislocations for the regular ribbons, while only homogenous grains exist in the field-assisted melt-spun ribbons. The reason for the grain size reduction may be explained by the magnetic field induced seed-size reduction effect. Since the iron-rich seeds are formed on the wheel surface where the temperature is far below the Curie temperature of Fe (770 °C), they can be magnetized and absorb energy from an external magnetic field, which could result in the changes in solidification process.^[230] Similar procedures have been applied to the field-assisted melt-spun Nd₂Fe₁₄B/Fe–Co ribbons with nominal composition Nd₁₆Fe₇₆B₈ + 40 wt.% Fe₆₅Co₃₅.^[231] A higher magnetic field of about 0.32 T perpendicular to the wheel surface reduces the grain size, increases the texture, improves the exchange coupling, and in sequence, increases the energy product (BH)_max from 16 to 17.5 MGOe.

Fig. 29.

	Figure Option ViewDownloadNew Window
	Fig. 29. TEM images of Nd15Fe77B8 ribbons melt-spun at the wheel speed 25 m/s (a) without a field and (b) in the field of 0.15 T which is perpendicular to ribbon plane.[230]

The intrinsic magnetic properties of Pr₂Fe₁₄B are similar to those of Nd₂Fe₁₄B, but Pr₂Fe₁₄B does not have spin reorientation behavior, so it can be used at low temperature. In addition, Pr₂Fe₁₄B has slightly higher anisotropy field than Nd₂Fe₁₄B, as shown in Table 5. Therefore, Pr₂Fe₁₄B is a good candidate for the magnetically hard phase in nanocomposite magnets. To gain best magnetic properties, the optimizations on microstructure such as grain size of both magnetically hard and soft phase, soft phase distribution, grain boundary control are all very important, which is very similar to those observed in Nd₂Fe₁₄B/α-Fe nanocomposites. The magnetic properties of melt-spun Pr₂Fe₁₄B/α-Fe nanocomposites with different element substitutions are summarized in Table 10.

Table 10.

Table 10.

Table 10. μ0Mr, μ0Hc, and (BH)max of Pr2Fe14B/α-Fe nanocomposite magnets with different compositions and different element substitutions. . Composition μ0Mr/T μ0Hc/T (BH)max/MGOe Reference Pr12Fe82B6 (stoichiometric) 0.95 1.2 19.5 [74] Pr8Fe87B5 (30.4 vol% α-Fe) 1.17 0.59 22.6 Pr10Fe84B6 1.16 0.81 21.3 [240] Pr7Dy1Fe86B6 1.14 0.66 16.9 [234] (Pr0.75Nd0.25)2Fe14B 1.06 1.01 23.1 [217] Pr9Fe76Co10B5 1.15 0.63 17.3 [235] Pr9.5Fe78.5Cr2B10 0.89 1.16 14.7 Pr9.5Fe78.5Nb2B10 0.90 1.32 16.9 Pr9.5Fe78.5Ti2B10 0.88 1.28 16.3 [236] Pr9.5Fe78.5V2B10 0.89 1.11 15.5 Pr9.5Fe78.5Zr2B10 0.88 1.01 14.5 Pr8Fe84Nb2B6 1.08 0.65 17.9 [97] Pr9Fe74Co12GaB5 1.22 0.65 22.2 [239] Pr8Fe86.4B5Mn0.6 1.33 0.58 23.1 [237] Pr8Fe85VB6 1.25 0.62 23.4 [238] Pr8Fe85VB5C 1.22 0.69 21.7 Pr8Dy1Fe74.5Co10Nb0.5B6 (125 MPa) 1.03 0.87 17.1 [233] Pr8Dy1Fe74.5Co10Nb0.5B6 (5 GPa) 1.11 1.02 23.6

Table 10. μ0Mr, μ0Hc, and (BH)max of Pr2Fe14B/α-Fe nanocomposite magnets with different compositions and different element substitutions. .

Goll et al. studied the effects of composition on microstructure and magnetic properties of Pr₂Fe₁₄B/α-Fe two-phase magnets from the stoichiometric Pr₂Fe₁₄B to a nominal composition with 54.2 vol% α-Fe.^[74] The room temperature magnetic properties of the Pr₂Fe₁₄B/α-Fe ribbons are summarized as a function of the α-Fe concentration, as shown in Fig. 30. With an increasing amount of α-Fe phase, an enhancement of the remanence is achieved, but a decrease of the coercivity is observed. For the energy product, a maximum value of 22.6 MGOe is found at 30.4 vol% α-Fe where the balanced remanence and coercivity are approached. At low α-Fe concentrations ( ) the coercivity is higher than one half of the remanence, (BH)_max sensitively depends on the remanence according to the theoretical upper limit of isotropic sample, i.e., . On the other hand, if at high α-Fe concentrations ( ), the coercivity is too small so (BH)_max is limited by irreversible demagnetization processes, leading to a drastic decrease energy product. The microstructure of Pr₈Fe₈₇B₅ magnet containing 30.4 vol% α-Fe is very fine grain, as shown in Fig. 31(a). Even with high-resolution TEM one cannot observe any boundary phase. This microstructure is clearly different from the decoupled single-phase Pr₁₅Fe₇₈B₇ magnet. Two characteristic maxima in the grain size distribution are observed in the grain size distribution of Pr₁₀Fe₈₄B₆ ribbon containing 14.0% α-Fe, as shown in Fig. 31(b), which correspond to Pr₂Fe₁₄B (20–30 nm) and α-Fe (∼15 nm) phases, respectively. This behavior is very similar to that of exchange coupled Nd₂Fe₁₄B/α-Fe composite magnets.^[73] This fine microstructure is the key factor to gain such a high energy product in the melt-spun ribbons.

Fig. 30.

	Figure Option ViewDownloadNew Window
	Fig. 30. (color online) Magnetic properties as a function of the α-Fe phase content of nanocomposite Pr2Fe14B/α-Fe magnets.[74]

Fig. 31.

	Figure Option ViewDownloadNew Window
	Fig. 31. (a) TEM image of a melt-spun Pr8Fe87B5 ribbon containing 30.4 vol% α-Fe; (b) grain size distribution of a melt-spun Pr10Fe84B6 ribbon containing 14.0 vol% α-Fe.[74]

The effects of the grain size on magnetic properties of a series of nanocrystalline melt-spun Pr_xFe_94−xB₆ alloys (6 ) have been investigated by Mendoza–Suárez and Davies.^[232] It is found that the remanence increases with decreasing grain size up to a maximum at which vitrification is initiated for all cases with different compositions. The intrinsic coercivity also rises with decreasing grain size except for the single-phase near-stoichiometric alloys. The significantly better combinations of coercivity and maximum energy product are obtained for the Pr–Fe–B than for the Nd–Fe–B resulting from the higher anisotropy field of the Pr₂Fe₁₄B phase. (BH)_max increases with remanence up to 30 vol% α-Fe content, beyond this point it falls, which agrees with the conclusion reported by Goll et al.^[74]

It is interesting to note that high pressure leads to grain refinement and thus more efficient exchange coupling in the Pr₂Fe₁₄B/α-Fe nanocomposites. Wang et al. reported this high pressure effects in the Pr₈Dy₁Fe_74.5Co₁₀Nb_0.5B₆ ribbons.^[233] The grain size of Pr₂Fe₁₄B phase decreases from 22.5 to 8.7 nm, while the grain size of α-(FeCo) phase also drops from 30.3 to 10.1 nm, when increasing pressure from 125 MPa to 5 GPa. As a result, the magnetic properties are improved from 1.03 T, 0.87 T, and 17.1 MGOe to 1.11, 1.02 T, and 23.6 MGOe for remanence, coercivity, and energy product, respectively. However, when hot pressing under even higher pressure, the crystallization is constrained and the hot-pressed magnets retains a certain amount of amorphous phase besides the Pr₂Fe₁₄B and α-(FeCo) phases, which results in the deterioration of the magnetic properties.

Similar to Nd₂Fe₁₄B/α-Fe, element substitutions, such as RE, refractory elements, and other elements, can help on grain refinement and grain boundary modifications of the Pr₂Fe₁₄B/α-Fe magnets, leading to improved magnetic properties. Chen et al. found that Dy substitution for Pr not only increases the coercivity because of the much higher magnetocrystalline anisotropy of Dy₂Fe₁₄B alloys, but also leads to the microstructure refinement in the Pr₂Fe₁₄B/α-Fe nanocomposite magnets.^[234] An optimum coercivity of 0.66 T is obtained in the Pr₇Dy₁Fe₈₆B₆ magnet as compared to 0.43 T in the Pr₈Fe₈₆B₆ magnet. As a result, the energy product increases from 9.1 MGOe in the Dy-free magnet to 16.9 MGOe in the 1 at.% Dy magnets, due to the enhanced exchange coupling between the magnetically hard and soft phases. As mentioned in Section 3.2.2, a Pr:Nd ratio of 3:1 in the RE₂Fe₁₄B magnets generated the best magnetic properties with an energy product of 23.1 MGOe.^[217]

Zhang et al. investigated the Co substitution for Fe in the melt-spun Pr₉Fe_86−xCo_xB₅ (x = 0–16) nanocomposites.^[235] It was found that Co obviously raises the saturated magnetization from 1.50 T for Pr₉Fe₈₆B₅ to 1.80 T for Pr₉Fe₇₆Co₁₀B₅ (measured under an applied field of 6.5 T), which may be attributed to the formation of α-(Fe, Co) phase. At the same time, the remanence, coercivity, and maximum energy product changes from 0.96 T, 0.57 T, and 11.8 MGOe to 1.15 T, 0.63 T, and 17.3 MGOe, respectively. Further substitution of Co for Fe results in a decrease of magnetic properties because of the presence of minor amounts of 1:5 and 2:17 phases. It was also observed that the Curie temperature of the magnetically hard phase increases linearly by increasing Co-substitution at the rate of ΔT_c = 9 °C/at%.

Chang et al. found that both Pr₂Fe₂₃B₃ and Fe₃B phases have been completely suppressed and the volume fraction of magnetically hard phase Pr₂Fe₁₄B increases in the Pr_9.5Fe_78.5M₂B₁₀ (M = Cr, Nb, Ti, V, and Zr) ribbons with all the selected refractory element substitutions.^[236] Accordingly, the coercivity of the ribbons is improved remarkably from 0.83 T for ternary Pr_9.5Fe_80.5B₁₀ ribbons to 1.01–1.32 T for all chosen refractory element-substituted nanocomposites. The optimal magnetic properties with an energy product of 16.9 MGOe is obtained in Nb-substitution Pr_9.5Fe_78.5Nb₂B₁₀ ribbons. Chen et al. confirmed that the largest property enhancement is obtained in Nb-substituted Pr₈Fe₈₄Nb₂B₆ magnets with a coercivity of 0.65 T and a maximum energy product of 17.9 MGOe, as compared to the coercivity of 0.43 T and the energy product of 9.1 MGOe in the Pr₈Fe₈₆B₆ magnets.^[97] The gain of the magnetic properties is attributed to the most uniform microstructure with the smallest average grain size of 10–20 nm in the Nb-substituted magnets. Ti substitution has a very similar role as Nb substitution in the Pr₂Fe₁₄B/α-Fe but with a slightly poorer performance, which is also similar to their function in the Nd₂Fe₁₄B/α-Fe nanocomposite magnets.

Yang et al. reported significant improvements of microstructure and magnetic properties of Pr₂Fe₁₄B/α-Fe nanocomposite magnets by Cu and Mn substitutions for Fe.^[237] It was found that Mn is more effective on the enhancement of coercivity than Cu. The remanence has maximum values when the content of Cu and Mn reaches 0.2 and 0.6, respectively. The remanence of the ribbons with Mn substitution again has higher value than that with Cu substitution, which may be attributed to magnetic dilute effect of nonmagnetic element Cu. The saturated magnetization decreases from 1.68 T for Pr₈Fe₈₇B₅ to 1.61 T for Pr₈Fe_86.4B₅Cu_0.6 ribbons, while that of Pr₈Fe_86.4B₅Mn_0.6 only deceases to 1.65 T. In comparison with Cu substitution, the Mn substitution significantly improves the (BH)_max, which mainly results from the stronger exchange-coupling effect due to fineness of grain in the melt-spun Pr₈Fe_87−xB₅Mn_x ribbons. Excellent magnetic properties with remanence of 1.33 T, coercivity of 0.58 T, and energy product of 23.1 MGOe are obtained in the Pr₈Fe_86.4B₅Mn_0.6 ribbons. The authors also studied the effects of addition of V or C or the combined addition on the structure and magnetic properties of melt-spun Pr₈Fe_86−xV_xB_6−yC_y ribbons.^[238] Compared with addition-free ribbons, 1 at.% V addition is found to reduce the grain sizes and improve their magnetic properties due to a strong exchange coupling between the magnetically hard and soft phases. A normalized remanence of 0.82, a coercivity of 0.62 T, and a maximum energy product of 23.4 MGOe in melt-spun Pr₈Fe₈₅VB₆ ribbons are obtained at room temperature. The combined addition of V and C leads to the formation of an intermediate phase of VC at grain boundaries, which appears as a pinning barrier during magnetization and results in a slight increase of the coercivity value to 0.69 T for melt-spun Pr₈Fe₈₅VB₅C ribbons. However, energy product drops to 21.7 MGOe for the ribbons with C substation. Zhang et al. studied the effect of Ga additive on the microstructure and magnetic properties of nanocomposite Pr₂(Fe, Co)₁₄B/α–(Fe, Co) ribbons.^[239] The microstructure refinement by introducing 1 at.% Ga leads to a stronger exchange coupling between the magnetically hard and soft phases in comparison with that in Ga-free sample. Therefore, the remanence and energy product increase from 1.14 T, 17 MGOe for Ga-free ribbons to 1.22 T, 22.2 MGOe for the Ga-added ribbons.

The compound Sm₂Fe₁₅Ga₂C has an anisotropy field exceeding 12 T at room temperature, which is much larger than those of Nd₂Fe₁₄B and Pr₂Fe₁₄B (7.6 and 9.0 T, respectively). However, the introduction of nonmagnetic Ga atoms decreases the saturated magnetization monotonically. As a result, the saturated magnetization is only about 1.0 T, which is much lower than those of Nd₂Fe₁₄B and Pr₂Fe₁₄B. In order to obtain high-energy product, it is necessary to improve the saturated magnetization and remanence, while keeping the coercivity sufficiently high. Exchange coupled nanocomposite magnets meet this requirement as magnetically hard phase Sm₂Fe₁₅Ga₂C can provide large coercivity while magnetically soft phase α-Fe can provide high magnetization.

Since Sm is easy to evaporate during alloying, it may be difficult to control the grain size and its distribution precisely by directly melt spinning, especially for the Sm₂Fe₁₅Ga₂C-based nanocomposite with low Sm concentration. Cheng et al. first fabricated the Sm₂Fe₁₅Ga₂C/α-Fe nanocomposites with low Sm contents directly by melt spinning without a subsequent heat treatment in 1998.^[241] The alloy with nominal composition Sm₈Fe₇₂Ga₈C₁₂ was melt-spun at a wheel speed of 15–35 m/s. It was found that the wheel speed affects the microstructure sensitively. The ribbons quenched at 15 and 17.5 m/s are composed of a phase with the Th₂Zn₁₇-type structure and α-Fe, while those quenched at 25 m/s and above show nearly complete amorphous state. The average grain size of Sm₂Fe₁₅Ga₂C and α-Fe are found to decrease with increasing the wheel speed from 140 and 50–60 nm for 15 m/s to 60–70 and 20–30 nm for 18.5 m/s, respectively. A maximum energy product of 5.9 MGOe was obtained for ribbons spun at 17.5 m/s with remanence of 0.71 T and coercivity of 0.55 T. The authors further found that Zr additive in the above nanocomposites could lead to the microstructure refinement and significantly improve the magnetic properties.^[242] Figure 32 presents the TEM micrographs and the grain size distributions of the Zr-free and Zr-containing samples. It shows that the ribbons contain two different phases, i.e., the major phase Sm₂Fe₁₅Ga₂C with larger grains and the minor phase α-Fe with smaller grains. The α-Fe grains are located as isolated particles at the grain boundaries of the major phase. The grain size of the Zr-free sample has very broad distributions from 50 to 140 nm for 2:17 phase and from 20 to 50 nm for α-Fe. The Zr additive appears to be very successful in retarding the grain growths during crystallization. The grain size of the Zr-containing sample becomes more homogeneous and distributes in the range of 40–70 nm for the 2:17 phase and 13–30 nm for α-Fe. The improved microstructure therefore leads to a smoother hysteresis loop and higher remanence in the Zr-containing ribbon, as shown in Fig. 33. As a result, magnetic properties with remanence of 0.72 T, coercivity of 0.77 T, and energy product of 9.0 MGOe were achieved through Zr substitution in the Sm₂Fe₁₅Ga₂C/α-Fe nanocomposites.

Fig. 32.

	Figure Option ViewDownloadNew Window
	Fig. 32. (a) TEM images and (b) grain size distributions for the Zr-free (Sm8Fe72Ga8C12) and Zr-containing (Sm7.9Fe71.3Ga7.9Zr1C11.9) ribbons.[242]

Fig. 33.

	Figure Option ViewDownloadNew Window
	Fig. 33. Room-temperature hysteresis loops of the Zr-free (Sm8Fe72Ga8C12) and Zr-containing (Sm7.9Fe71.3Ga7.9Zr1C11.9) ribbons.[242]

Zhang et al. studied the magnetic properties of Sm₂Fe₁₅Si₂C/α-Fe nanocomposites with Cu additive, which are prepared by melt spinning at wheel speed of 20 m/s.^[243] The average grain size of 2:17 carbides and α-Fe in the ribbons is about 50 and 30 nm, respectively. The ribbons exhibit a single hard magnetic phase behavior and remanence enhancement due to inter-grain exchange coupling between magnetically hard and soft phases. The remanence increases with the increase of α-Fe phase content in ribbons, while the coercivity decreases rapidly, which agrees with the correlation between remanence and coercivity in other nanocomposite magnets. The highest maximum energy product around 10.7 MGOe was obtained for Sm₈Fe₇₆Cu₄Si₈C₄ ribbons with a coercivity of 0.5 T.

The effect of pressure on the microstructure of Sm₂(Fe,Si)₁₇C_x/α-Fe nanocomposite magnets has been studied by Zhang et al. through pressing the amorphous Sm₈Fe₈₅Si₂C₅ alloy under pressures up to 6 GPa at 650 °C.^[244] The grain size of α-Fe and Sm₂(Fe,Si)₁₇C_x decreases from 30.6 to 6.4 and from 28.5 to 5.8 nm, respectively. The volume fraction of the 2:17-type carbide increases with increasing pressure. A significant increase in both remanence and coercivity, comparing to those prepared under normal pressure, is observed in the nanocomposites prepared under high pressure. As a result, an energy product is significantly improved in the Sm₂(Fe,Si)₁₇C_x/α-Fe nanocomposites.

Table 11.

Table 11.

Table 11. μ0Mr, μ0Hc, and (BH)max of Sm2Fe15M2C/α-Fe (M = Ga, Si) nanocomposite magnets with different compositions and different element substitutions . Composition μ0Mr/T μ0Hc/T (BH)max/MGOe Reference Sm8Fe72Ga8C12 0.71 0.55 5.9 [241] Sm7.9Fe71.3Ga7.9Zr1C11.9 0.72 0.77 9.0 [242] Sm8Fe76Cu4Si8C4 0.78 0.50 10.7 [243] Sm8Fe85Si2C5 (4 GPa) 0.88 0.62 N/A [244]

Table 11. μ0Mr, μ0Hc, and (BH)max of Sm2Fe15M2C/α-Fe (M = Ga, Si) nanocomposite magnets with different compositions and different element substitutions .

Doll et al. fabricated the precipitation-hardened ribbons through the melt spinning process and appropriate post heat treatment for the first time.^[249] Although the alloys are prepared by melt spinning, however, the annealing procedure is very similar to that of conventional sintered magnets, i.e., T_h = 1160 °C, t_h = 30 min, T_a = 840 °C, t_a = 19 h, and cooling rate of 0.75 °C/min. TEM analysis shows that a rhombic structure parallel to the c axis and a hexagonal network in the basal plane are formed in the ribbons after annealing, which is very similar to the microstructure of the sintered magnets. The characteristic cell size is about 50–70 nm, which depends on the composition. Figure 35 shows the temperature dependence of the coercivity for both magnets in comparison with three other common permanent magnetic materials, namely Nd₂Fe₁₄B, Sr ferrite and AlNiCo, respectively. For the Cu/Fe-rich sample, the coercivity decreases with increasing temperature, i.e., a normal behavior of coercivity with negative temperature coefficient. In contrast, the coercivity of the Cu/Fe-poor magnet changes only slightly from room temperature to 500 °C and even shows a positive temperature coefficient in the temperature range between 227 °C and 427 °C. The coercivity reached a maximum of 0.4 T at 427 °C. To understand the different coercivity behavior, EDX was scanned from one 2:17 cell over a 1:5 cell boundary into another 2:17 cell for the Cu/Fe-poor magnet. It is clear that Cu have much higher concentration in the grain boundary 1:5 phase than the cell 2:17 phase. Although the mechanism is still a matter of debate, it is agreed that the abnormal coercivity behavior is related to the Cu-profile along the grain boundary phase.

Fig. 35.

	Figure Option ViewDownloadNew Window
	Fig. 35. (color online) (a) Temperature dependence of the coercivity of Cu/Fe-rich and Cu/Fe-poor in comparison with other common permanent magnetic materials; (b) EDX profiles of Sm, Cu, Fe, and Zr (contents standardized to the Co content) in the vicinity of a 2:17/1:5 phase boundary in Cu/Fe-poor magnet.[249]

Yan et al. further simplified the process for the melt-spun Sm(CoFeCuZr)_z ribbons and found that the homogenizing process (solid solution treatment) actually is not necessary.^[250] They compared the magnetic hardening processes of the Sm(Co_0.74Fe_0.1Cu_0.12Zr_0.04)_7.5 ribbons with and without solid solution treatment (homogenizing at 1160 °C for 4 hours). The coercivity of the ribbons with solid solution treatment continues to increase with longer isothermal aging time and a value of 2.4 T is obtained after aging for 12 h at 850 °C, as shown in Fig. 36(a). This behavior is very similar to that found in conventional precipitation-hardened Sm(CoFeCuZr)_z magnets. However, quite different behavior is observed in the ribbons without solid solution treatment. A very high coercivity of 2.8 T is developed by simply slow cooling from 850 to 400 °C, without isothermal aging, which is much higher than that of homogenized ribbons subjected to the same aging process (1.64 T). A short aging time about 3 h could bring the coercivity to its maximum of about 3.0 T. This behavior could be attributed to the already formed 1:7 phase in the melt-spun ribbon which is the final product with supersaturated Cu and Zr in the phase of the homogenizing process for the conventional Sm(CoFeCuZr)_z alloys. The 1:7 phase is believed to play important roles in developing the cellular/lamellar microstructure and high coercivity. TEM images shown in Fig. 36(b) confirmed that the microstructure, consisting of cellular and lamellar phases, was formed in the ribbon without solid solution treatment. Therefore, the time consuming and intricate aging regimen, which is normally required to develop high coercivity for bulk Sm(CoFeCuZr)_z permanent magnets, is significantly simplified in melt-spun materials, i.e., the precipitation hardening is realized in melt-spun ribbons by simply slow cooling from 850 to 400 °C.

Fig. 36.

	Figure Option ViewDownloadNew Window
	Fig. 36. (a) Evolution of the coercivity dependent on the isothermal aging time at 850 °C for the Sm(Co0.74Fe0.1u0.12Zr0.04)7.5 ribbons with and without solid solution treatment; (b) typical TEM micrographs of samples without solution treatment.[250]

It is very interesting that Yan et al. observed the disappearance of the large coercivity in the above melt-spun Sm(CoFeCuZr)_z magnet by an additional short time aging.^[251] Figure 37(a) shows the coercivity dependence of the ribbons on the aging temperature of the additional aging process while the aging time is fixed at 5 min. One can see that higher aging temperatures lead to a more significant loss in the coercivity and the ribbons subjected to aging at 850 °C for only 5 minutes showed a very low coercivity of 0.16 T, which is far lower than the coercivity of 3.0 T of the original ribbon without aging. It should be noted that no change in the coercivity was observed after further increasing the aging time. The dramatic change of coercivity should be attributed to the Cu distribution across the 1:5 grain boundary phase. Figure 37(b) shows the nanoprobe EDX microanalysis profiles of Cu obtained by a line scan across the interface between the 2:17 matrix and the 1:5 in the ribbons with and without additional aging process. A very sharp peak of Cu content (24 at.%) was detected in the middle of the 1:5 phase indicating that the Cu distribution in the cell boundary is highly inhomogeneous, which is similar to that observed in Ref. [249]. After the additional short-time aging at 850 °C, the Cu distribution in the 1:5 boundary phase becomes more homogenous and the large gradient of Cu content is no longer detectable. Therefore, the high coercivity in the melt-spun precipitation hardened magnet is proposed to be originated from the large gradient of domain wall energy within the 1:5 grain boundary phase and not from the Cu diffusion from 2:17 phase to 1:5 phase.

Fig. 37.

	Figure Option ViewDownloadNew Window
	Fig. 37. (a) Dependence of the coercivity of Sm(Co0.74Fe0.1Cu0.12Zr0.04)7.5 ribbons on aging temperature (additional isothermal aging for 5 min); (b) nanoprobe EDX profiles of Cu across the interface between the 2:17 matrix and the 1:5 grain boundary of the ribbons before (sample A) and after (sample B) additional aging at 850 °C for 5 min.[251]

Rong et al. found that the TbCu₇-type 1:7 phase is formed by direct melt spinning in the Gd(Co_0.88−xCu_xFe_0.09Zr_0.037 alloy as well.^[252] It is confirmed that solid solution treatment and the isothermal aging are not necessary for precipitation hardening in the melt-spun materials, i.e., high coercivity can be obtained by simply slow cooling the ribbons from 850 to 400 °C. Figure 38 shows the microstructure of the as-spun and precipitation-hardened Gd(Co_0.88−xCu_xFe_0.09Zr_0.037 ribbons. TEM studies reveal that the as-spun ribbons are quite heterogeneous consisting of large crystalline areas, which mainly have a TbCu₇-type structure, as shown in Fig. 38(a). The grain boundaries remain in the precipitation-hardened ribbons, as shown in Fig. 38(b). EDX analysis shows that the 1:7 phase boundary is composed of a Gd-rich phase and its composition is almost unchanged before and after the heat treatment. The mean size of the 1:7 regions is about . The microstructure of one 1:7 region, as presented in Figs. 38(c) and 38(d), shows clearly that the cellular structure, which consists of a Fe-rich 2:17 cell phase, a Cu-rich 1:5 cell boundary phase, and a Zr-rich platelet phase, is obtained by the simple annealing treatment. Interestingly, the cell size is only about 10–20 nm which is much smaller than that of typical sintered magnets. The cell boundaries are about 2–5 nm in width and are coherent to the cellular grains. However, no cellular microstructure is formed in the 1:7 phase boundaries. The temperature dependences of remanence and coercivity of precipitation-hardened Gd(Co_0.88−xCu_xFe_0.09Zr_0.037 ribbons without isothermal aging are given in Fig. 39. All ribbons exhibit a positive temperature coefficient of remanence from RT to about 400 °C. This is a consequence of the antiparallel coupling between Gd and Co magnetic moments. Since the rare-earth sublattice disorders rapidly increase with temperature, the ferromagnetic coupling of Gd–Co magnets results in an increasing magnetization with increasing temperature. Abnormal coercivity (positive temperature coefficient) is only observed in ribbons with x = 0.10. Further addition of Cu causes the monotonic temperature dependence of coercivity.

Fig. 38.

	Figure Option ViewDownloadNew Window
	Fig. 38. TEM micrographs of the Gd(Co0.88−xCuxFe0.09Zr0.03)7 ribbons. (a) As-spun; (b) precipitation-hardened without isothermal aging; (c) and (d) one 1:7 region of (b) which is observed perpendicular and parallel to the c axis, respectively.[252]

Fig. 39.

	Figure Option ViewDownloadNew Window
	Fig. 39. (color online) Temperature dependences of remanence and coercivity of the precipitation-hardened Gd(Co0.88−xCuxFe0.09Zr0.03)7 ribbons without isothermal aging.[252]

Besides the simplified precipitation hardening process, texture can be developed in melt-spun Sm(CoFeCuZr)_z ribbons, without the conventional sintering process. Therefore, a relatively high energy product can be obtained in such ribbons. Rong et al.^[253,254] and Yan et al.^[139] have studied the effects of composition on the crystallographic texture of the precipitation-hardened (x = 0–0.25, y = 0–0.30, w = 0–0.10, and –7.5) ribbons, which were melt-spun at low wheel speed of 5 m/s. The results are summarized below.

Effect of z value (or Sm content). The studies show that the z value or Sm content has little influence on the crystallographic texture of the melt-spun ribbon, as reported by both Rong et al.^[253] and Yan et al.^[139]

Effects of Cu. Figure 40 shows the typical XRD patterns of the melt-spun and precipitation-hardened Sm(Co_0.87−xCu_xFe_0.1Zr_0.037 ribbons with different Cu contents. The significantly enhanced intensity of the (200) and (110) peaks for the Cu-free ribbon suggests that they are crystallographically anisotropic with easy magnetization direction parallel to the longitudinal direction of the ribbons. However, a small quantity of Cu substitution (x = 0.05) leads to a significant increase in the intensity of (111) reflection, which is close to that expected for isotropic materials, indicating that no crystallographic texture is present in the ribbons. Further addition of Cu leads to stronger intensity of (111) peak and disappearance of (200) and (110) peaks, indicating that the degree of crystallographic texture increases with increasing Cu content. After precipitation hardening through isothermal aging at 850 °C for 3 hours, the 1:7 phase is segregated into 1:5 and 2:17 phases. The broadened peaks of the precipitation-hardened samples indicate that phase transformation leads to the formation of a fine microstructure. The crystallographic texture is preserved in ribbons after precipitation hardening.

Fig. 40.

	Figure Option ViewDownloadNew Window
	Fig. 40. XRD patterns of melt-spun and precipitation-hardened Sm(Co0.87−xCuxFe0.1Zr0.03)7 ribbons with different Cu contents.[253]

Effects of Fe. Figure 41(a) shows the XRD patterns of the melt-spun Sm(Co_0.87−yFe_yCu_0.1Zr_0.037 ribbons with different Fe contents. The as-spun ribbons consist mainly of a single TbCu₇-type phase. The intensity of the diffraction peaks of the Fe-free ribbons is close to those of isotropic materials, indicating no obvious texture in the ribbons. The addition of Fe leads to significantly enhanced intensity of (111) peak in the ribbons with y = 0.2, which indicates the increase of texture degree with increasing Fe content. However, further addition of Fe leads to the disappearance of crystallographic texture in the ribbons.

Fig. 41.

	Figure Option ViewDownloadNew Window
	Fig. 41. XRD patterns of the melt-spun ribbons (a) Sm(Co0.87−yFeyCu0.1Zr0.03)7 with different Fe contents; (b) Sm(Co0.8−wFe0.1Cu0.1Zrw)7 with different Zr contents.[253,254]

Effects of Zr. Figure 41(b) shows the XRD patterns of the melt-spun Sm(Co_0.8−wFe_0.1Cu_0.1Zr_w)₇ ribbons with different Zr contents. Again, all as-spun ribbons consist of a single TbCu₇-type phase. The significantly enhanced intensity of the (200) and (110) peaks for the Zr-free ribbons suggests that they are crystallographic anisotropic with easy magnetization direction parallel to the longitudinal direction of the ribbons. However, no obviously preferential orientation is observed in the samples with , which agrees with the results reported in Ref. [139]. It is interesting to note that partially preferential orientation starts to form with Zr content w = 0.04. Further Zr substitution leads to a more pronounced increase in the intensity of (110) peak and the perfect crystallographic texture is obtained in the ribbons with w = 0.08. Due to the appearance of 2:7 phase, the crystallographic texture disappears in the ribbons with . It is also found that Zr substitution results in a significant modification to the microstructure of the ribbons surface. Figure 42 shows the SEM micrographs of the free surface of the as-spun ribbons with different Zr content. The small grains of approximately in diameter are obtained in the isotropic ribbon with w=0.01. Whereas, the columnar dendrite grains, with their long axes mainly parallel to the longitudinal axes of the ribbons, start to form in the ribbon with w=0.04. The average grain size is about . With further Zr substitution, it can be found clearly that the columnar dendrite grains arrange more formally and compactly than those of low Zr content ribbons. It is suggested that the Zr addition favors the formation of preferential orientation, which agrees closely with the XRD analyses given in Fig. 41(b).

Fig. 42.

	Figure Option ViewDownloadNew Window
	Fig. 42. (color online) SEM images of the free surface of melt-spun Sm(Co0.8−wFe0.1Cu0.1Zrw)7 ribbons: (a) w = 0.01, (b) w = 0.04, and (c) w = 0.08.[254]

The desired magnetic properties of Sm(CoFeCuZr)_z ribbons are controlled not only by the composition and texture, but also by the heat-treatment process, i.e., the isothermal temperature and time, the cooling rate, and the final quench temperature. The systematically work to optimize the magnetic properties of Sm(CoFeCuZr)_z melt-spun ribbons was reported in Refs. [253] and [255].

Figure 43 summarizes the compositional dependence of μ₀M_r, μ₀H_c, and (BH)_max of precipitation-hardened ribbons. It should be noted that the remanence is determined by both the composition and texture. As a general rule, remanence increases with increasing the content of Fe/Co content and the preferred crystallographic texture, while decreases with increase of Zr/Cu content, as shown in Fig. 43(a). Coercivity, on the other hand, mainly relies on the composition and its related effects on the microstructure, but not the texture. As one can see in Fig. 43(b), all elements play an important role in the development of coercivity, which has been reported in both melt-spun ribbons and conventional sintered Sm(CoFeCuZr)_z magnets. The coercivity is less than if any of the elements is missed. Among the elements, Cu has the biggest impact on coercivity of the precipitation-hardened magnets. The ribbons without Cu show coercivity about zero, while a high Cu content about x = 0.15–0.20 can lead to a coercivity about 3.0 T at room temperature. With a controlled composition and related texture, the maximum energy product around 11 MGOe at room temperature is achieved in the ribbons by the above-mentioned simple process.

Fig. 43.

	Figure Option ViewDownloadNew Window
	Fig. 43. (color online) The compositional dependence of (a) μ0Mr, (b) μ0Hc, and (c) (BH)max of precipitation-hardened ribbons.[253]

Besides the composition, the control of precipitation hardening process is another critical factor for the coercivity development in Sm(CoFeCuZr)_z ribbons. Figure 44 shows the room-temperature coercivity of Sm(Co_0.87−xFe_0.1Cu_xZr_0.037 ribbons as a function of isothermal aging temperature (T_a), aging time (t_a), cooling rate (v_c), and quenching temperature (T_q). It should be noted that the results in Fig. 44 give one fact at a time, i.e., all other parameters are fixed when one factor is studied. The coercivity dependences on these parameters are summarized below. (i)

Coercivity is low when T_a is lower than 750 °C, while it increases quickly with varying T_a up to 900 °C. Further increase of T_a to 950 °C and above leads to a drop of coercivity.

Fig. 44.

	Figure Option ViewDownloadNew Window
	Fig. 44. (color online) The effects of isothermal aging temperature (Ta), aging time (ta), cooling rate (vc), and quench temperature (Tq) on the room-temperature coercivity of Sm(Co0.72Fe0.1Cu0.15Zr0.03)7 ribbon.[255]

The coercivity increases fast with a decrease in the quenching temperature T_q. Similar to the effects of v_c during the slow cooling, the decrease in T_q leads to an increase of Cu concentration in the 1:5 boundary phase and thus an increase in coercivity.

As mentioned above, the composition affects temperature dependence of coercivity H_c(T). In fact, all parameters of the precipitation hardening process affect the temperature dependence of coercivity as well. Figure 45 shows the H_c(T) curves of the Sm(Co_0.77Fe_0.1Cu_0.1Zr_0.037 ribbons quenched at different temperature T_q, as an example. The anomalous temperature dependence of coercivity can be obtained in the ribbons with . In addition, the tendency of the abnormal temperature dependence of coercivity gradually disappears with decreasing T_q. Besides the quenching temperature, other annealing parameters such as T_a, t_a, and v_c significantly affect the H_c(T) curves as well. The details can be found in Ref. [255]. All relations between H_c(T) curves and the annealing parameters may be explained by the change of microstructure and microchemistry, which leads to a variation of the distribution and amount of Cu in the 1:5 cell boundaries.

Fig. 45.

	Figure Option ViewDownloadNew Window
	Fig. 45. (color online) Hc(T) curves of the Sm(Co0.77Fe0.1Cu0.1Zr0.03)7 ribbons with different quenching temperatures.[255]