1 Introduction

Since discovery of excellent magnets Nd–Fe–B in 1983, they have become the most widely used type of rare earth permanent magnets [1–3]. Rare earth permanent magnetic materials play important roles in economics and lives. However, it consumes mainly low-abundant rare earth elements to produce traditional rare earth permanent magnets, leading to the extreme unbalance in the utilization of resources. The consumption of Nd and Pr increases significantly due to the increasing production of Nd–Fe–B permanent magnets in the world. Moreover, the rare earth mineral is paragenic. Pr and Nd only account for 22 wt% of light rare earth metals, and La (≥25 wt%) and Ce (≥47 wt%) are the most abundant and the cheapest in the mineral. Developing high-abundant rare earth permanent magnets is necessary not only for reducing cost, but also for efficient utilization of the resources. Recently, the substitution of Pr–Nd by Ce [4–8] or misch metal (MM) alloys has attracted renewed attention [9–12]. However, the magnetic properties of magnet based on misch metal were found to be low [13–16]. The magnitude of coercivity does not meet the application of medium- and high-grade products.

In this work, it was taken advantage of double main phase alloy method to manufacture high-performance R–Fe–B permanent magnets with low production cost. The commonly associated rare earth permanent magnets were prepared using misch metal to save the Pr, Nd resources. This method avoids unnecessary separation process of rare earth elements, which utilizes the resources efficiently and protects the environment. Moreover, the method of doping PrNd nanoparticles was introduced in an attempt to improve the coercivity of sinter magnets containing MM [17–19]. Magnetic properties and magnetic domain of MM magnet with MM accounting for 30% of total rare earth elements were surveyed.

2 Experimental

The raw material MM is one of the typical misch metals from Bayan Obo mine in Baotou, China. It consists of 28.63 wt% La, 50.13 wt% Ce, 4.81 wt% Pr, 16.38 wt% Nd and other inevitable impurities. According to the dual-alloy method, two alloys with nominal compositions of (Pr,Nd)30FebalB1 and ((PrNd)0.3MM0.7)30FebalB1 were prepared by strip casting (SC) technique. The two varieties cast strips independently performed hydrogen decrepitation and jet-milling to obtain powders with an average particle size of 3–5 μm. And then the two powders were mixed together according to the nominal composition of ((PrNd)1−x MM x )30FebalB1 (wt%, x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7), respectively. The powders were compacted after aligning under a constant magnetic field of 1.8 T and compacting under a pressure of 30 MPa. Then the green compacts were sintered at 1010–1060 °C for 2 h under a vacuum of 4.5 × 10−3 Pa. Subsequently, the sintered magnets were annealed at 900 °C for 110 min and then post-annealed at 600 °C for 120 min. In order to improve coercivity, the powders with nominal composition of ((PrNd)0.7MM0.3)30FebalB1 and the PrNd (a composition of Pr25.5Nd74.5 (wt%)) nanoparticles were homogeneously mixed. The PrNd nanoparticles powders were prepared by the inert gas condensing method using direct current (DC) arc plasma metal nanometer powder preparation equipment (ZJ-NM-KY). The magnets were prepared by blended powder. The density of the magnets was measured by water immersion method at room temperature. The magnetic properties of the samples were measured using an NIM-2000 magnetic measuring device. The heat flow measurements were taken by a differential scanning calorimeter (DSC, STA449C-2). The phases of the magnets were characterized using X-ray diffractometer (XRD, PANalytical X’pert Powder) with Co Kα radiation. Microstructural investigations of the samples were carried out using scanning electron microscope (SEM, FEI Nova Nano 200) equipped with energy-dispersive spectroscope (EDS) and transmission electron microscope (TEM, JEM-2100). Domain structure was observed using a BH-786IP-PK high-field Kerr microscopy.

3 Results and discussion

The magnetic properties of ((PrNd)1−x MM x )30FebalB1 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7) magnets are shown in Fig. 1. It can be noticed from Fig. 1 that remanence (B r), maximum energy product ((BH)max), intrinsic coercivity (H cj) and density (ρ) gradually decrease with the increase of MM content accounting for the total rare earth amount, referred to as MM/R. It is due to the fact that the intrinsic magnetic properties of La2Fe14B and Ce2Fe14B are far inferior to those of Nd2Fe14B. For La2Fe14B and Ce2Fe14B, the saturation polarizations (4πM s) are 1.38 and 1.17 T, and the magnetocrystalline anisotropy fields (H A) are 2000 and 3600 kA·m−1 [2], respectively. For the substitution of MM for Pr and Nd, the R2Fe14B (R = La, Ce, Pr, Nd) phase forms, which decreases the magnetic polarization and anisotropy field in the magnet. It can be clearly seen that H cj decreases with the increase of MM/R. H cj declines rapidly to the value of 668.4 kA·m−1 for MM/R ≤ 20% and then slowly for MM/R ≥ 20%. For MM/R = 20%, the magnetic properties reach a preferable level of B r = 1.326 T, H cj = 666.0 kA·m−1 and (BH)max = 330.2 kJ·m−3. In addition, H cj declines significantly for MM/R = 30%.

Fig. 1
figure 1

Magnetic properties of ((PrNd)1−x MM x )30FebalB1 magnets at room temperature

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XRD patterns of ((PrNd)1−x MM x )30FebalB1 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7) magnets are shown in Fig. 2. It can be seen from Fig. 2 that the magnets consist mainly of R2Fe14B phase (R = La, Ce, Pr, Nd). A small amount of CeFe2 is observed (2θ ≈ 41° and 47°) in the magnet with low coercivity. CeFe2 is not only stable in equilibrium R–Fe binary alloy system, but also paramagnetic at room temperature [20]. In addition, the R-rich phases are also observed (2θ ≈ 35°) in the magnet. It can be noticed that the intensity of diffraction peak of R-rich phase becomes disorganized with the increase of MM/R. Therefore, the existence of impurity phase may harm the coercivity among the boundaries of the grains.

Fig. 2
figure 2

XRD patterns of ((PrNd)1−x MM x )30FebalB1 magnet samples after sintering and annealing

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Figure 3 displays DSC curves of ((PrNd)1−x MM x )30FebalB1 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7). Endothermic peaks at T C in Fig. 3 refer to the Curie temperatures of the magnets, at which the main phase changes from ferromagnetic to paramagnetic. It can be noticed from Fig. 3 that the Curie temperatures of the ((PrNd)1−x MM x )30FebalB1 magnets decrease with the increase of MM contents. The Curie temperatures of R2Fe14B phase are 304.0, 296.7, 288.6, 281.0, 280.1, 255.0 and 253.0 °C for x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7, respectively. Endothermic peaks at T 1 in Fig. 3 refer to the melting temperatures of the R2Fe14B phases, at which the main phase changes from solid to liquid phase. It can be noticed from Fig. 3 that the melting temperatures of R2Fe14B phase are 1180.7, 1169.8, 1167.7, 1158.3, 1157.9, 1155.6 and 1142.7 °C for x = 0, 0.1, 0.2, 0.3, 0.4, 0.5 and 0.7, respectively. Above results are due to that the Curie temperatures of Ce2Fe14B and La2Fe14B are lower than that of Nd2Fe14B [21]. In addition, some La and Ce enter into R2Fe14B phase, resulting in the decrease of the melting temperature of R2Fe14B phase [21, 22]. This indicates that adding MM can reduce the Curie temperatures and the melting temperatures of R2Fe14B phases. So the magnets can be sintered and annealed at lower temperature.

Fig. 3
figure 3

DSC curves of ((PrNd)1−x MM x )30FebalB1 magnets with different MM contents

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SEM images of magnets are shown in Fig. 4, and EDS analysis data are listed in Table 1 for different regions of the sample with nominal composition of ((PrNd)0.8MM0.2)30FebalB1 and ((PrNd)0.3MM0.7)30FebalB1. The gray areas correspond to R2Fe14B phase, and the white areas correspond to intergranular phase. It can be seen in Fig. 4 that LaCe-rich (marked as C, E) and LaCe-lean (marked as A, B, F, G, H) matrix grains exist in the magnets. Many triple-junction phases of R-rich phase with more La and Ce elements are observed in Fig. 4, as listed in Table 1. La and Ce contents in Regions 3 and 4 are more than those in Regions 1 and 5 in Fig. 4, respectively. Agglomeration occurs easily in LaCe-rich intergranular phase. The R-rich phase exists among the boundaries of the grains, which weakens the coupling between the hard phase grains [23–26]. The decrease of partial R-rich phase results in the abnormal triple-junction R-rich phases. Besides, since the coupling between the hard phase grains reinforces which is related to the abnormal triple-junction R-rich phases, the coercivity gradually decreases with the increase of MM contents. The agglomeration of the R-rich phase is observed in sintering R–Fe–B permanent magnets containing MM, which leads to such a low magnetic properties in the magnets.

Fig. 4
figure 4

SEM back-scattered images of ((PrNd)1-x MM x )30FebalB1 magnets with a x = 0.2 and b x = 0.7

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Table 1 Compositions of regions in Fig. 4a, b determined by EDS (wt%)
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SEM image of ((PrNd)1−x MM x )30FebalB1 powders and TEM image of PrNd nanoparticles are shown in Fig. 5. The size of ((PrNd)1-x MM x )30FebalB1 powders and PrNd nanoparticles are about 3–5 μm and 50–80 nm, respectively. In order to improve coercivity, PrNd nanoparticles were added to the ((PrNd)0.7MM0.3)30FebalB1 magnet by dual-alloy method. Table 2 shows the rare earth contents of the magnets with MM/R = 20% (S1) and MM/R = 30% with 5 wt% PrNd nanoparticles (S2). The total rare earth content (∑RE/Rm) increases from 30 wt% in S1 to 33.4 wt% in S2, which may lead to the enhancement of the coercivity. Nd(Pr)/Rm increases from 24 wt% in S1 to 24.8 wt% in S2 and MM/Rm increases from 6 wt% in S1 to 8.6 wt% in S2. Consequently, the production cost is reduced by the increase of MM content. Demagnetization curves of the magnets as a function of doping content of PrNd nanoparticles are shown in Fig. 6. The coercivity improves with the increase of PrNd nanoparticles content. The magnetic properties (B r = 1.332 T, H cj = 872.9 kA·m−1, (BH)max = 318.6 kJ·m−3) of the magnet used dual-alloy method with 5 wt% PrNd nanoparticles are better than those of the magnet without PrNd nanoparticles doping. The MM/R reduces to 25.7% for the magnet with 5 wt% PrNd nanoparticles, but the magnetic properties exceed those of the magnet with MM/R = 20%.

Fig. 5
figure 5

SEM image of ((PrNd)0.7MM0.3)30FebalB1 a and TEM image of PrNd nanoparticles b

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Table 2 Rare earth content of MM/R = 20% (S1) and MM/R = 30% with 5 wt% PrNd nanoparticles (S2) magnets (wt%)
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Fig. 6
figure 6

Demagnetization curves of sintered ((PrNd)0.7MM0.3)30FebalB1 magnets with different doping contents of PrNd nanoparticles

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To study the reason for coercivity change, the magnetic domain structures of the magnet with and without PrNd nanoparticles doping are shown in Fig. 7. The magnets were cut into a dimension of 3 mm × 8 mm × 5 mm with the easy axis along the long direction of the sample. The surfaces of the samples were polished for clearly revealing domain configuration. It can be seen from Fig. 7 that bright and dark domains are observed in same grain. Many continuous domains are found in the magnet without PrNd nanoparticles through the grain boundary in Fig. 7a. For the magnet with 5wt% PrNd nanoparticles doping, continuous domains decrease with the increase of PrNd nanoparticles content in the magnet in Fig. 7b. Magnets with homogenous distribution of R-rich phase normally exhibit excellent magnetic properties [27–30]. Improvement of the coercivity may be caused by the fact that the Re-rich phase becomes homogeneous in the magnet with the addition of PrNd nanopowders. It not only favors the isolation of adjacent hard magnetic grains, but also attaches to the surface of the Nd2Fe14B grains by taking the form of particle phase.

Fig. 7
figure 7

Domain patterns of ((PrNd)0.3MM0.7)30FebalB1 magnet a without and b with 5 wt% PrNd nanoparticles doping

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4 Conclusion

The magnetic properties of R–Fe–B prepared with different ratios of MM–Fe–B and Nd–Fe–B were systematically investigated. The results show that the magnetic properties decrease gradually with the increase of MM content, but they still have potential to manufacture permanent magnets. H cj gradually decreases with the increase of MM contents because of the agglomeration of R-rich phase. The coercivity enhances with PrNd nanoparticles content increasing, whereas both the remanence and the maximum energy product reach the maximum values with 5 wt% PrNd doping. The maximum energy product is 318.6 kJ·m−3 and coercivity is 872.9 kA·m−1 for MM/R = 25.7%. Domain investigation shows that the continuous domains are reduced by PrNd nanoparticle doping. Therefore, it is concluded that PrNd nanoparticle doping is an efficient method to improve coercivity in R–Fe–B magnets containing MM.