Introduction

In recent years, the oil crisis and energy problems have emerged one after another. Therefore, rare earth elements play a key role in many technologies. Among them, the development of permanent magnets has received increasing attention, and neodymium (Nd) has stood out in the manufacturing of permanent magnets and is prepared as so-called NdFeB permanent magnets with the strongest magnetic properties in today’s permanent magnets.1 With the advancement of science and technology, the market for NdFeB permanent magnets has gradually expanded. Not only was it maturely used in the voice coil motor (VCM) in hard disk drives in the early days,2 but in recent years, it has gradually expanded to the development of electric vehicles (EV),3 wind power turbines,4 and other equipment. However, the bottleneck of NdFeB permanent magnets is their low Curie temperature (Tc) of about 320°C or 608 °F.5 One significant limitation of NdFeB permanent magnets is their relatively low Curie temperature (Tc) of approximately 583 K (309.85°C)5 compared to other permanent magnets. This characteristic poses a challenge in applications requiring high operating temperatures. Factors such as magnetic stability, thermal stability, safety considerations (due to potential hazards when operating near Tc), and material integrity (as high temperatures can accelerate corrosion and oxidation) must be carefully evaluated. Consequently, NdFeB magnets may not be ideal for applications demanding higher operational temperatures. When the operating temperature exceeds the Curie temperature (Tc), the magnetic properties of NdFeB permanent magnets decrease significantly, directly affecting equipment stability.6 By increasing the Curie temperature (Tc) and coercivity (iHc) of a permanent magnet, better thermal stability can be achieved.7 To date, many reports have suggested the NdFeB permanent magnet with elemental substitutions as a viable candidate by substituting rare earth ions and other cations such as Dy, Ga, Pr, Co, and Tb to significantly improve their coercivity and thermal stability.8,9,10,11,12,13,14,15 Kim et al. used various sintered NdFeB magnets containing 28–31% Nd and 0–7% Dy to enhance their coercivity, which can be applied to drive motors in hybrid vehicles with operating temperatures as high as 200–220°C.16 Therefore, how to improve the coercivity of the magnets so that the NdFeB permanent magnets could still maintain good thermal stability and excellent magnetic properties under high temperature environment will be the focus of this study.

Due to the scarcity and rising costs of heavy rare earths, recent strategies to enhance the coercivity (iHc) of NdFeB permanent magnets aim to avoid resource depletion and cost increases from excessive use. Commonly adopted methods include "grain boundary diffusion" and "magnetic powder refinement." Grain boundary diffusion involves coating the surface of the sintered magnet with heavy rare earth elements such as dysprosium (Dy) and terbium (Tb). These elements are then diffused into the magnet along the grain boundaries at high temperatures through heat treatment. Although the grain boundary diffusion method can effectively reduce the amount of heavy rare earth elements and increase the iHc of the magnet, it causes impurities to remain inside the magnet, resulting in a decrease in Br and (BH)max17,18 and additional cost. The conventional sintered NdFeB magnet manufacturing process, as shown in Fig. 1a, includes powder milling, hydrogen decrepitation (HD), jet milling (JM), additive mixing, magnetic field forming, cold isostatic pressing (CIP), vacuum sintering, and heat treatment, etc. Through the hydrogen decrepitation (HD) and jet milling (JM) process, the size of the magnetic powder is more refined. In this way, a more concentrated particle size distribution of the magnetic powder can be obtained,19 thereby achieving the effect of enhancing the iHc of the permanent magnet. A novel low-oxygen powder metallurgy process was also proposed in 2021 and succeeded in producing anisotropic Sm2Fe17N3 sintered magnets from finely pulverized powder without coercivity degradation.20 In addition, the influence of grain size on the coercivity of NdFeB magnets was investigated by Bittner et al. in 2017.21 They found that if the average grain size can be reduced to 1 µm, sintered magnets with a coercivity > 2000 kA/m (about 25.13 kOe) can be attained without using any heavy rare earth elements.21

Fig. 1
figure 1

(a) The conventional sintered NdFeB magnet manufacturing process; (b) comparison of the conventional and modified manufacturing process of sintered NdFeB magnet.

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To further enhance the iHc and (BH)max of NdFeB permanent magnets without using heavy rare earths for grain boundary diffusion, the conventional NdFeB permanent magnet manufacturing process was modified in this study. In the new process, the dehydrogenation in the conventional process, which was originally performed in hydrogen decrepitation, will be moved to the sintering stage, followed by jet milling (JM) and subsequently the secondary fine pulverization, with the objective of reducing the ultrafine powder. The effect of different sintering process conditions on the formability of the magnets was also discussed. In this study, the green body will be sintered in vacuum and argon atmospheres, respectively, to investigate the effect of sintering at different heating rates on sintered NdFeB permanent magnets.

Experimental Procedure

Commercial strip cast alloys, with a refined composition of (NdPr)29.9Tb0.6Co1.0Cu0.1Al0.2B0.96Febal (total rare earths, TRE: 30.5 wt.%; Nd = 22.34 wt.%; Pr = 7.56 wt.%; Tb = 0.6 wt.%; Dy < 0.01 wt.%; Fe: Bal.; B = 0.96 wt.%; Co = 1 wt.%; Cu = 0.1 wt.%; Al = 0.2 wt.%; Ga < 0.01 wt.%; Nb < 0.01 wt.%; Ti < 0.01 wt.%; Zr < 0.01 wt.%; Mn = 0.02 wt.%; Si = 0.03 wt.%; Mg < 0.01 wt.%; Ni < 0.01 wt.%; Y < 0.01 wt.%; La < 0.01 wt.%; Ce < 0.01 wt.%; Sm < 0.01 wt.%; O = 100 ppm; C = 100 ppm; N < 0.01 ppm; H = 0.02 ppm), served as the initial material for the hydrogen decrepitation (HD) process. Details of this process are described elsewhere.22 The density of the green body is meticulously regulated to approximately 3.8 g/cm3 following magnetic field forming, with a field strength of 1.8 Tesla, to optimize its magnetic properties.23

The hydrogen decrepitation (HD) and jet milling (JM) process in the previously described conventional NdFeB permanent magnet manufacturing process was modified in this study, and their comparison is depicted in Fig. 1b. In the conventional process, the so-called hydrogen decrepitation (HD), the NdFeB magnetic powder is packed into a stainless steel container and vacuumized to 10−2 Pa first. Then, pure hydrogen is filled into the container, so that the hydrogen pressure inside the container reaches 105 Pa (about one atmosphere); after another 20–30 min, the temperature in the container will gradually increase.24 The heating of the stainless steel container is due to the exothermic effect after the hydrogenation reaction of the NdFeB magnetic powder. When the temperature rises to between 150°C and 300°C, the hydrogen is absorbed by the Nd (neodymium)-rich phase. As the temperature continues to rise to about 650°C, the hydrogen will be further strongly absorbed by the Nd-rich phase and the main Nd2Fe14B phase will crack.25 The reason is that the NdFeB magnetic powder forms hydride after absorbing hydrogen, which makes the inside of the magnetic powder crack. This reaction is called hydrogen decrepitation (HD). As shown in Fig. 2a, when Nd2Fe14B reacts with hydrogen to form hydrides Nd2Fe14BHx and NdHy, its volume will expand, resulting in local volume expansion and internal stress of hydrides during hydrogenation. Since NdFeB alloy compounds are brittle materials, hydrogen decrepitation (HD) occurs when the internal stress exceeds the fracture strength.26

Fig. 2
figure 2

Schematic diagram of (a) hydrogen decrepitation (HD) process; (b) the jet mill (JM) designed in this study.

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In the modified process, only hydrogen absorption is performed in this hydrogen decrepitation (HD) step without dehydrogenation. The atoms in NdFeB permanent magnets arrange themselves in a regular pattern, forming an Nd2Fe14B tetragonal crystal structure with numerous interstitial spaces, such as tetrahedral sites. Hydrogen atoms, being the smallest, can easily penetrate and migrate between these spaces. It is therefore conjectured that retaining the dehydrogenation step until the sintering process allows hydrogen to carry away nitrogen and carbon impurities from the magnet during subsequent evacuation processes. Furthermore, it is postulated that the effect of hydrogen embrittlement can fragment lubricants into smaller molecules in hydrogen-rich, high-temperature environments,27 thereby facilitating their removal during subsequent evacuation processes and consequently reducing carbon impurity content. As a result, carbon and nitrogen elements would be released together with hydrogen during the sintering process, which can reduce the impurities remaining in the magnetic powder, thereby increasing the coercivity iHc. After the hydrogen decrepitation (HD), the jet milling (JM) is carried out to control the average particle size of the magnetic powder within the range of 4–5 μm.28 Subsequently, after the jet mill (JM, Hosokawa-Alpine AFG100, Japan), a secondary fine pulverization process is executed to make the magnetic powder surface smoother, less ultrafine powder < 0.1 μm, and more concentrated particle size distribution. Due to the reduction in particle size and more concentrated particle size distribution, the intrinsic coercivity iHc could be further increased. On the other hand, the shape of the magnetic powder tends to be smooth, which leads to a higher degree of alignment in the magnetic field forming step and further increases the maximum energy product (BH)max.

The schematic diagram of the jet mill (JM) designed in this study is illustrated in Fig. 2b. In the process, initially, the speed of the classifying wheel above the jet mill (JM) is set to 9500 rpm, and the pressure of the air flow for pulverization is 0.6 MPa. Under the above process condition, the powders are driven by nitrogen injected into the grinding chamber from the nozzle and collide with each other and cause the effect of pulverization. When the powder is small enough, it can pass through the classifying wheel. Next, the magnetic powder is collected and sent to the jet mill (JM) again, and the rotating speed of the classifying wheel is increased to 22,000 rpm. The pressure of the air flow for pulverization is controlled at 0.45 MPa. Therefore, most of the magnetic powders are left in the grinding chamber, and the powder collide with each other driven by a lower air flow pressure. In this way, the pressure of the air flow for pulverization is moderated, and the collision between the magnetic powders can be slowed down. Therefore, the sharp protrusions on the surface of the magnetic powders can be modified, and the appearance of the magnetic powders becomes smoother. After 10 min, the rotating speed of the classifying wheel is reduced to 3000 rpm to discharge the magnetic powder in the grinding chamber. In addition, it should be noted that the secondary pulverization method used in the modified process of this study makes the magnetic powder stay in the grinding chamber of the jet mill for a longer time, resulting in a smaller average particle size of the magnetic powder. When the powder is small enough, it can pass through the classifying wheel, so there will be less ultrafine powder < 0.1 μm in the magnetic powder obtained by the secondary pulverization method. From what has been discussed above, so-called “secondary fine pulverization” is introduced as a new method to modify the conventional manufacturing process of sintered NdFeB magnet, given enhanced magnetic performance. To further improve the formability and magnetic properties of NdFeB permanent magnets, the green bodies were sintered in vacuum and argon atmospheres, respectively, and the different sintering conditions of magnets were systematically discussed as follows: (1) rapid heating in vacuum environment; (2) slow heating in vacuum environment; (3) rapid heating in argon atmosphere, and (4) slow heating in argon atmosphere.

In this study, the particle size and distribution of the magnetic powder were measured by the HELOS-BR RODOS dry laser particle size analyzer (Sympatec, Germany). Morphologies of the magnetic powder were observed by a field emission scanning electron microscopy (FE-SEM, Zeiss, Oberkochen, Germany). After sintering, the magnets were first magnetized to saturation state by charging and demagnetizing equipment (SJ-3560MD, Taiwan), and then the magnetic properties of the magnets were measured at room temperature of 25°C using the BH tracer (NIM-2000, Beijing Metrology Institute, China). To characterize the BH curve of the magnet, the BH Tracer necessitates samples with dimensions of 10 mm × 10 mm × 8 mm, with the C-axis aligned parallel to the 8 mm surface. The sample is initially saturated in a magnetic field of 4.5 Tesla, after which it is positioned within the measurement coil. The poles of the BH Tracer are precisely aligned in close contact with the sample. The BH Tracer first performs forward magnetization up to 1.2 Tesla followed by reverse demagnetization to evaluate the BH curve of the sample. The content of impurities, such as oxygen, carbon, and nitrogen, in the finished NdFeB magnet, was determined using a carbon analyzer (EMIA-Pro, HORIBA Scientific, Japan) and O/N/H Analyzer (EMGA-Expert, HORIBA Scientific, Japan), respectively.

Results and Discussion

Comparison of Powder Morphology and Particle Size Distribution Between Conventional and Modified Processes

In this study, the secondary fine pulverization method is introduced in the modified manufacturing process of sintered NdFeB magnet. The SEM photographs in Fig. 3a and b show the morphologies of the NdFeB magnetic powders obtained after pulverization in the conventional process and the modified process proposed in this study, respectively. The difference between the two is that the former is the powder obtained only by jet milling (JM), while the latter is the magnetic powder obtained by jet milling (JM) first and then the secondary fine pulverization. The results in Fig. 3 show that the powder obtained by the conventional process has sharper edges and more ultrafine powder < 0.1 μm, and the morphology of the NdFeB magnetic powder treated by the secondary fine pulverization method in the modified process is much smoother than that of the powder in the conventional process without the secondary fine pulverization, and the ultrafine powder is significantly reduced. The reason is that during the second fine pulverization step, the pressure of the airflow for pulverization is reduced to 0.45 MPa, and the collision between the magnetic powders can be alleviated. Thus, the sharp protrusions on the surface of the magnetic powders can be modified, and the magnetic powder morphology tends to be more rounded, which would help to enhance the intrinsic coercivity iHc of NdFeB permanent magnets. Interestingly, the secondary pulverization method used in the improvement process of this study allows the magnetic powder to stay in the grinding chamber of the jet mill for a longer time, resulting in a smaller average particle size of the magnetic powder. When the powder is small enough, it can pass through the classifying wheel. In this way, ultrafine powder < 0.1 μm in the magnetic powder obtained by the secondary pulverization method is further reduced. When the proportion of ultrafine powder in the magnetic powder is small, the oxygen content in the magnetic powder will be reduced,29 which can further enhance the magnetic properties of the magnet.30

Fig. 3
figure 3

SEM photograph of the NdFeB magnetic powders obtained after pulverization in (a) conventional process and (b) modified process proposed in this study.

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The particle size distributions of the NdFeB magnetic powders of the conventional process and the modified process are shown in Fig. 4, obtained by measuring three samples respectively. As illustrated in Fig. 4, the average particle size of the powder obtained by the conventional process is D50 = 3.74 μm, and the concentration degree of particle size distribution is D90/D10 = 4.94.21 In contrast, the average particle size of the powder obtained by the modified process is D50 = 3.51 μm, and the concentration degree of particle size distribution is D90/D10 = 4.66. Figure 4 shows that the particle size distribution of the magnetic powder in the modified process is more concentrated (i.e., smaller D90/D10),21 and the average particle size of the powder is relatively finer than that in the conventional process, which is in good agreement with the SEM result of Fig. 3.

Fig. 4
figure 4

Particle size distribution of the NdFeB magnetic powders of the conventional process and modified process.

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In conventional NdFeB permanent magnet processing, the magnetic powder exhibits a broad particle size distribution, with a D90/D10 ratio of 4.94. However, after undergoing secondary fine pulverization in this study, the particle size distribution of the magnetic powder narrows to a D90/D10 ratio of 4.66. Compared to traditional methods, the secondary fine pulverization process proposed in this study utilizes the jet mill classifier wheel to first separate and remove excessively fine powders, while retaining powders of normal size for lower-energy milling inside the jet mill’s grinding chamber. Under these grinding conditions, the impact energy between powders during the jet mill process is insufficient to further refine the powder but is effective in smoothing the powder's surface morphology. Consequently, while the particle size D50 after secondary fine pulverization is 3.51 µm—slightly lower than the D50 of 3.74 µm in the traditional process—the reduction in excessively fine powders contributes to an improvement in the final magnet’s properties from a magnetic perspective.

Influence of Sintering Temperature and Heating Rate on NdFeB Permanent Magnets

Conventional Sintering Process for NdFeB Permanent Magnets

The sintering process of NdFeB magnets is carried out in a vacuum sintering furnace to avoid the increase of oxygen content and the deterioration of magnetic properties. Figure 5 illustrates the sintering profiles for the conventional process currently in use and the modified processes introduced in this study, with all stages conducted under vacuum conditions except for the final cooling stage to 35°C. The sintering profile of the conventional process, as depicted in Fig. 5, is detailed as follows: Initially, the temperature is held at 20°C for 5 min, then gradually increased to 400°C > 120 min and maintained at this temperature for 90 min. The temperature is subsequently raised from 400°C to 580°C over 120 min and held at 580°C for 90 min. Following this, the temperature is ramped up to 800°C in 210 min and held at 800°C for 90 min. The process continues with an increase from 800°C to 900°C in 90 min, with a subsequent hold at 900°C for 90 min. The temperature is then further elevated to 1000°C over 120 min and maintained at this peak for 420 min. Finally, the temperature is reduced from 1000°C to 200°C in 30 min. Subsequently, argon gas is introduced into the furnace, and the temperature is gradually lowered from 200°C to 35°C within 60 min in an argon atmosphere to complete the sintering process. This conventional method remains widely employed in the industry for the preparation and mass production of commercial NdFeB permanent magnets.22

Fig. 5
figure 5

Comparison of the sintering profiles for the conventional and modified processes, conducted under vacuum conditions except for the final cooling to 35°C.

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In the following, the formability and magnetic properties of the NdFeB magnets will be systematically investigated through the modified process of the secondary fine pulverization combined with sintering method proposed in this study.

Modified Sintering Process with Rapid Heating in Vacuum

To improve the magnetic properties of the NdFeB magnet, the subsequent experiments used the powder after the secondary pulverization treatment for magnetic field formation and sintering process. The sintering profile for the modified process with rapid heating in vacuum, also shown in Fig. 5, is described as follows: The temperature was initially held at 20°C for 5 min, then ramped from 20°C to 400°C over 240 min and maintained at 400°C for 90 min. It was then increased from 400°C to 580°C over 120 min and held at 580°C for 150 min. Subsequently, the temperature rose from 580°C to 800°C in 210 min and was held at 800°C for 90 min. The process continued with a ramp from 800°C to 900°C in 60 min, followed by a 90-min hold at 900°C. The temperature was then increased from 900°C to 1000°C over 60 min and maintained at 1000°C for 240 min. Next, the temperature decreased from 1000°C to 600°C in 30 min, then increased from 600°C to 1000°C in 120 min and was held at 1000°C for 240 min. Finally, the temperature was reduced from 1000°C to 200°C in 30 min. Argon gas was then introduced into the furnace, and the temperature was cooled from 200°C to 35°C within 60 min in an argon atmosphere to complete the sintering process. Figure 6 depicts the appearance of the magnet after sintering with the modified processes introduced in this study. All stages were performed under vacuum conditions, except for the final cooling stage to 35°C. As observed in Fig. 6a, the magnets sintered using the modified rapid heating process in a vacuum exhibit obvious cracks. From a powder metallurgy perspective, it is inferred that the suboptimal forming performance results from the reduced particle size of the powder subjected to secondary fine milling in this study (as indicated by Figs. 3 and 4), leading to insufficient forming strength of the magnets. It is important to note that, according to powder metallurgy principles, finer powders present greater challenges in the forming process. Given the differences in the NdFeB permanent magnet manufacturing process, the following analysis offers a more detailed explanation of why the new process (i.e., moving the dehydrogenation step to the sintering stage) introduced in this study could result in magnet cracking after sintering if the sintering process is not further modified. In the conventional processing of NdFeB permanent magnets, the strip-cast alloy undergoes hydrogen decrepitation (HD) in two distinct stages: (1) Hydrogen Absorption: At room temperature and under a hydrogen overpressure environment, a significant amount of hydrogen infiltrates the alloy's grain boundary phases and main phase lattice, leading to alloy fragmentation. (2) Dehydrogenation: The temperature is subsequently elevated to 580°C, and hydrogen is extracted from the grain boundary phases and main phase lattice within a vacuum environment. Consequently, in the conventional NdFeB permanent magnet sintering process, the initial hydrogen release is attributed to the residual hydrogen remaining after the dehydrogenation process. In this study, only the hydrogen absorption stage was conducted during hydrogen decrepitation (HD), omitting the dehydrogenation stage (i.e., the dehydrogenation step moved to the sintering stage). As a result, the considerable amount of hydrogen absorbed by the alloy remains within the green body prior to sintering. During sintering, the hydrogen that would typically be released during the dehydrogenation stage is instead rapidly and significantly expelled during the initial hydrogen release of the sintering process. Due to the exceptionally low strength of the green body before sintering, this rapid and substantial hydrogen release can lead to visible cracking of the green body. However, since NdFeB undergoes liquid-phase sintering, the liquid phase formed at high temperatures can assist in densification by filling and repairing micro-cracks. Therefore, in the metallographic examination of sintered NdFeB permanent magnets, these micro-cracks may be eliminated and thus not detected. Although the implementation of a secondary fine milling process in this experiment enhanced the magnetic properties of the NdFeB magnets, the forming characteristics remained suboptimal, suggesting that additional process parameters may require further optimization.

Fig. 6
figure 6

Photos of magnets after sintering with the modified processes: (a) rapid heating and (b) slow heating, both under vacuum except for final cooling to 35°C.

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Modified Sintering Process with Slow Heating in Vacuum

To overcome the above experimental results of magnet crack, the sintering heating rate of the magnets was further modified in this study to improve the formability of the magnets. The sintering profile for the modified process with slow heating in vacuum, also illustrated in Fig. 5, is outlined as follows: Initially, the temperature was held at 20°C for 5 min, then gradually increased to 100°C over 20 min and maintained at 100°C for 60 min. The temperature was then raised from 100°C to 200°C over 20 min and held at 200°C for 60 min. Next, the temperature increased from 200°C to 300°C in 20 min and was held at 300°C for 60 min. This was followed by a ramp from 300°C to 400°C in 20 min, with a 60-min hold at 400°C. The temperature then rose from 400°C to 500°C over 20 min and was maintained at 500°C for 60 min. Subsequently, the temperature increased from 500°C to 580°C in 20 min and was held at 580°C for 120 min. The process continued with a ramp from 580°C to 800°C in 240 min and a 90-min hold at 800°C. The temperature was then raised from 800°C to 900°C over 60 min and held at 900°C for 90 min. Following this, the temperature increased from 900°C to 1000°C in 120 min and was maintained at 1000°C for 480 min. Finally, the temperature rapidly decreased from 1000°C to 100°C in 30 min. Argon gas was then introduced into the furnace, and the temperature was gradually reduced from 100°C to 35°C within 60 min in an argon atmosphere to complete the sintering process. Figure 5 reveals that the modified process with slow heating in vacuum effectively reduces the hydrogen release rate by extending the hold times during the heating stages, thereby lowering the overall heating rate. This strategy helps to mitigate the simultaneous release of large quantities of hydrogen gas, which can otherwise lead to green body cracking. Consequently, this approach preserves the integrity of the green body, preventing any cracking after the sintering process. Figure 6b illustrates the appearance of the magnet after sintering with the modified process described above, which is completely free of cracks. By adjusting the sintering heating rate in this study's modified process, the limitations in forming strength associated with secondary fine pulverization treatment are effectively addressed. Additionally, this adjustment helps to mitigate the issue of green body cracking that can occur after sintering.

Modified Sintering Process with Rapid Heating in Ar

After overcoming the issue of the magnet crack caused by the secondary fine pulverization treatment of the modified process, the original sintering profile (i.e., the sintering profile in Section "Modified Sintering Process with Rapid Heating in Vacuum") was used and argon gas was introduced to replace the original vacuum environment during the heating stage of sintering to further improve the magnetic properties of the magnet. To prevent the hydrogen from being released before reducing the carbon and nitrogen elements in the magnet during vacuum sintering, Ar is introduced into the sintering process to prolong the residence time of hydrogen in the magnet. Figure 7 compares the sintering profiles of the conventional and modified processes, with all stages conducted under vacuum except for the pre−600°C warming stage and the final cooling to 35°C, during which argon was introduced. The sintering profile for the modified process with rapid heating in Ar as depicted in Fig. 7 is detailed as follows: Prior to reaching 580°C, the initial vacuum is terminated, and argon gas is introduced at 0.9 atm. This substitution of the vacuum environment extends the residence time of hydrogen within the magnet, thereby reducing internal impurities. Figure 8 depicts the appearance of the magnet after sintering with the modified processes introduced in this study, with all stages conducted under vacuum except for the pre−600°C warming stage and the final cooling to 35°C, during which argon was introduced. As observed in Fig. 8a, the magnets sintered using the modified rapid heating process in Ar still exhibit noticeable cracking. It is speculated that since the heating rate of dehydrogenation step in sintering process is still too fast currently, the rapid release of hydrogen causes the green body to crack. Therefore, in the following experiments, the dehydrogenation step in the sintering process will be further divided into more stages. In this way, the stop point in the heating stage will be increased to reduce the heating rate and further improve the crack of magnet.

Fig. 7
figure 7

Comparison of the sintering profiles of the conventional and modified processes, with all stages conducted under vacuum except for the pre−600°C warming stage and the final cooling to 35°C, during which argon was introduced.

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Fig. 8
figure 8

Photos of magnets after sintering with the modified processes: (a) rapid heating and (b) slow heating, with all stages conducted under vacuum except for the pre−600°C warming stage and the final cooling to 35°C, during which argon was introduced.

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Modified Sintering Process with Slow Heating in Ar

To improve the experimental results of magnet crack in Section "Modified Sintering Process with Rapid Heating in Ar," the sintering heating rate of the magnets was further modified in this study to improve the formability of the magnet sintered in Ar. The sintering profile for the modified process with slow heating in Ar, as depicted in Fig. 7, is also outlined as follows: Prior to raising the sintering temperature to 580°C, argon is introduced at 0.9 atm to facilitate dehydrogenation under partial pressure. This approach not only maximizes the effectiveness of hydrogen as a reducing agent but also incorporates temperature hold points to mitigate the hydrogen release rate, thereby preventing excessive hydrogen release that could compromise the integrity of the sintered green body and avoid fracturing. Figure 8b illustrates the appearance of the magnet after sintering with the modified process described above and demonstrates that the slow heating in an argon atmosphere effectively addresses the issue of insufficient forming strength caused by secondary fine pulverization, resulting in sintered magnets with a flawless appearance.

Comparison of Magnetic Properties of Sintered Magnets Prepared by Conventional Process and Modified Process with Slow Heating in Ar

In this study, the magnetic properties of 5 mm × 5 mm × 5 mm samples prepared using both the conventional and modified processes with slow heating in vacuum and Ar were measured simultaneously. Figure 9 presents a comparison of the demagnetization curves for magnets produced by these methods. The magnetic properties of the sample prepared using the conventional process were found to be Br = 14.50 kG, iHc = 12.44 kOe, (BH)max = 51.19 MGOe, BHH = 63.63, as illustrated in Fig. 9. Here, BHH = iHc + (BH)max is defined as the technical capability index for NdFeB permanent magnets.31 An elevated BHH value signifies enhanced overall magnetic properties of NdFeB permanent magnets.31 The sintering density of the magnet prepared by the conventional process is 7.53 g/cm3. For comparison, Fig. 9 also presents the demagnetization curves for magnets prepared using the modified slow heating process in vacuum and Ar. The magnetic properties of the sample processed with slow heating in vacuum are as follows: Br = 14.51 kG, iHc = 14.35 kOe, (BH)max = 51.42 MGOe, BHH = 65.77. The sintering density of this magnet is 7.54 g/cm3. Additionally, the magnetic properties of the sample processed with slow heating in Ar are: Br = 14.60 kG, iHc = 16.04 kOe, (BH)max = 52.46 MGOe, BHH = 68.50, with a sintering density of 7.58 g/cm3. These results highlight the enhanced magnetic properties achieved through the modified process with slow heating in Ar. In the modified process, both coercivity (iHc) and BHH exhibit significantly higher values compared to those of magnets produced by the conventional process. This demonstrates that the secondary fine pulverization and sintering techniques harnessed in the modified process of this study substantially enhance the magnetic properties of NdFeB permanent magnets.

Fig. 9
figure 9

Comparison of the demagnetization curves for magnets prepared using the conventional process versus that prepared with the modified slow heating process in vacuum and Ar.

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Table I details the process parameters for both the conventional and modified slow heating processes in Ar. It demonstrates that not only can modifying the hydrogen decrepitation timing and introducing secondary fine pulverization enhance the coercivity of the magnets, but these adjustments also improve thermal stability and formability. Moreover, Table II shows that the new process significantly reduces carbon and nitrogen content, a result of the hydrogen effect during dehydrogenation in the sintering process. The reduction in oxygen content is likely due to the decreased surface area of the magnetic powder exposed to air following the secondary fine pulverization, which minimizes the presence of ultra-fine magnetic powder.32

Table I Process parameters of the conventional and modified slow heating process in Ar
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Table II Comparison of carbon, nitrogen, and oxygen content of NdFeB permanent magnets prepared by conventional and modified slow heating process in Ar
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Figure 10 presents the SEM BED (backscattered electron detector) images of NdFeB permanent magnet samples prepared using the conventional process and the optimal process conditions (i.e., modified slow heating processes in Ar) in this study, respectively. From Fig. 10a, it is evident that sintered NdFeB permanent magnets produced using the conventional process exhibit noticeable fine grains (highlighted by the yellow dashed circle), a broad grain size distribution, and less smooth grain boundaries. However, Fig. 10b shows that sintered NdFeB permanent magnets prepared using the modified process with slow heating in Ar do not exhibit fine grains, have a narrower grain size distribution, and possess smoother grain boundaries. This improvement enhances the overall magnetic properties and is consistent with the results presented in Fig. 9.

Fig. 10
figure 10

SEM images comparing the microstructure of sintered NdFeB magnets prepared using (a) the conventional process and (b) the optimal process conditions (i.e., modified slow heating processes in Ar) in this study, respectively.

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Conclusion

In this study, the enhanced performance of sintered NdFeB permanent magnets without using heavy rare earths for grain boundary diffusion was realized by modifying the conventional process of NdFeB permanent magnets, such as pulverization and sintering. Hydrogen easily combines with the carbon and nitrogen in the NdFeB powder. Therefore, carbon and nitrogen elements would be released together with hydrogen during the sintering process, thereby reducing impurities remaining in the powder. Therefore, the dehydrogenation in the conventional process, which was originally carried out in the hydrogen decrepitation, was moved to the sintering stage in the new process, which contributes to the improvement of the magnetic properties of magnets. Furthermore, the magnetic powder will be treated by the so-called secondary fine pulverization method after the jet milling (JM) process. Through the secondary fine pulverization, the degree of collision between the powders is alleviated, and the appearance of the powders tends to be smoother. In addition, since the powder stays in the grinding chamber of the jet mill for a longer time during the secondary fine pulverization process, the ultrafine powder in the magnetic powder is further significantly reduced, and the oxygen content is also decreased. To prolong the residence time of hydrogen in the magnet, argon is introduced in the sintering of the modified process to further maximize the effect of hydrogen to reduce internal impurities, which further improve iHc. Besides, the heating rate of dehydrogenation step in sintering process was further controlled to avoid the instantaneous release of a large amount of hydrogen due to the too fast heating rate, which would cause the green body cracking during the sintering process. To sum up, the NdFeB magnet has a complete appearance without cracks and achieves the best magnetic properties with iHc = 16.04 kOe, (BH)max = 52.46 MGOe, and BHH = 68.50 are obtained by the new process modified in this study. However, the best magnetic properties of the magnets prepared by the conventional process only reach iHc = 12.50 kOe and (BH)max = 51.28 MGOe and BHH = 63.78. The iHc, (BH)max and BHH of NdFeB permanent magnets obtained by the modified process in this study are 22.07%, 2.25%, and 6.89% higher than those obtained by the conventional process, respectively. The magnetic properties of the magnets prepared by the modified process are upgraded from the original N52 to N50H grade of Hitachi Metals NEOMAX specification.33,34 According to the research results, the iHc in this study has been greatly improved, making NdFeB magnets more beneficial for applications in equipment operating in high temperature environments. We expect this study’s approach and process condition, possessing enhanced magnetic properties, will find applications far beyond the limitations of conventional NdFeB magnetic materials, and we envision that it can be widely used in new types of end applications, such as offshore wind turbines and hybrid electric vehicles.