Formation of the High-Coercive State in Rare-Earth Magnets

Abstract

The effect additives of TbAl₃ hydride have on the magnetic properties and structure of sintered permanent magnets based on Pr–Fe–Co–Cu–B alloy with increased temperature stability is studied. Such magnetic properties as B_r = 1.07 T, ВН_max = 216 kJ m⁻³, _jH_c = 2000 kA m⁻¹, H_k = 1680 kA m⁻¹, H_k/_jH_c = 0.84, and α = −0.030% °С⁻¹ (in the temperature range of 20–100°C) are achieved.

INTRODUCTION

There are currently two main groups of rare-earth sintered hard-magnetic materials (alloys of Nd–Fe–B and Sm–Co–Fe–Cu–Zr) in which such mechanisms of magnetization as the nucleation delay of reverse domains and the delay in the displacement of domain boundaries [1] are observed, respectively. The latter group of hard-magnetic materials has better temperature stability, due to its higher Curie temperatures, but requires more complex manufacturing technology [2]. The structural state of the Nd₂Fe₁₄B phase plays an important role in the first group of hard-magnetic materials [3–5]. The structural state of the main phase can be altered substantially (e.g., to improve its stability, to create an elastic-stressed state inside it, or to create a nanoheterogeneous distribution of alloying elements), by doping the base alloy in particular, along with features of the technological processes (e.g., mechanical doping, hydride dispersion, and using additives of hydrides of rare-earth metals). This allows us to greatly enhance the technological and operational resistance of permanent magnets (PMs). PMs are doped with such elements as Dy, Tb, Pr, and Co to improve their temperature stability on the basis of Nd₂Fe₁₄В intermetallic compound.

The aim of this work was to optimize the composition of Nd–Fe–B alloy and the process for reducing the reversible temperature coefficient of magnetic induction (α) to values ​​that correspond to sintered magnets based on Sm–Co alloys (0.03% °С⁻¹).

EXPERIMENTAL

Our initial alloys with the chemical composition shown in Table 1 were smelted in a vacuum induction furnace from pure charge materials in a medium of especially pure argon. Their chemical composition was monitored via atomic emission spectroscopy. Magnets with the chemical composition of the C-alloy and the ratio of this alloy with the addition of TbAl₃ were chosen such that the resulting chemical composition of the magnet corresponded to that of magnet B. The base alloys (1–3) and the additive alloy of TbAl₃ were subjected to hydride dispersion in a dry hydrogen channel at 400°C for 1 h, and were then finely ground inside a vibration mill for 50 min in isopropyl alcohol to a mean particle size of 3 microns. Joint fine grinding with the addition of TbAl₃ (1.2 wt %) was done for C magnets. The magnets were pressed in a transverse magnetic field and sintered for 2 h at T = 1100°C, after which they were thermally treated in vacuum for 2 h at T = 900°C and then slowly cooled (1–2°С min⁻¹) to 500°C and kept at this temperature for 1 h. Magnetic measurements were made using a МН-50 hysteresigraph in a closed magnetic circuit. The reversible temperature coefficient of magnetic induction (α) was measured in the temperature range of 20–100°C on magnet samples using a vibration magnetometer, and in a magnetic circuit using a teslameter and a microwebermeter. The Curie temperature (Т_С) was determined by measuring the temperature dependences of the initial magnetic permeability and magnetization. The microstructure of the magnets was investigated via optical and scanning microscopy (SEM) and local X-ray diffraction analysis (LXRDA). X-ray analysis of powders prepared from sintered magnets was done on a DRON-3M diffractometer using copper CuK_α radiation and a graphite monochromator.

Table 1.   Chemical composition of the alloys (wt %)

RESULTS AND DISCUSSION

The data from magnetic measurements of PMs at room temperature using the МН-50 hysteresisgraph in a closed magnetic circuit are presented in Table 2. Coefficient α was measured in the temperature range of 20–100°C. It was 0.060, 0.030, and 0.030% °С⁻¹ for magnets of types A, B, and C, respectively. Analysis of these measurements shows that such parameters as residual induction B_r, maximum energy product ВН_max, and coefficient α (in absolute value) were greater in magnitude for magnet A than for B- and C-type magnets. However, the other parameters (coercive magnetization force _jH_c and demagnetizing field intensity, at which the residual magnetization was 90% of H_k and H_k/_jH_c) are much lower. For B- and C-type magnets, such parameters as B_r, ВН_max, and α almost coincide. Parameters _jH_c, H_k, and H_k/_jH_c were greater in magnitude for C magnets than those of B magnets.

Table 2.   Magnetic properties of the alloys at room temperature

The results from measuring the magnetic properties of C-type magnets at elevated temperatures in the range of 20–100°C are presented in Table 3. These data allowed us to determine the temperature coefficient of the coercive magnetization force. It was 0.55% °С⁻¹ in the temperature range of 20–100°С. Curie temperature Т_С for the investigated magnets was found to be ~490, ~560, and ~570°C for magnets of the A, B, and C types, respectively.

Table 3.   Magnetic properties of magnet C at elevated temperatures

SEM and LXRDA studies of microstructure showed that the chemical composition of the main phase in A-type magnets is expressed by the formula (Nd_0.8Tb_0.2)₂(Fe_0.8Co_0.2)₁₄B (at %). In addition to the main phase, ones such as (Nd,Tb)₃(Fe,Co), (Nd,Tb)_rich, (Nd)_1.1(Fe,Co)₄B₄, (Nd,Tb)(Fe,Co)₂ were observed. The composition of the main phase for B- and C-type magnets was (Pr_0.7Tb_0.3)₂(Fe_0.72Co_0.28)₁₄B. In the latter case, there was a gradient distribution of terbium and aluminum in the grain (which was maximal at the grain boundaries). There was no boundary high-permeability Laves phase in B- and C-type magnets.

The lower structurally sensitive parameters (_jH_c, H_k, and H_k/_jH_c) for A-type magnets, relative to those of types B and C, can be explained by there being a lower terbium content in the main magnetic phase, and by the presence of high-permeability phase (Nd,Tb)(Fe,Co)₂. The data from magnetic measurements (i.e., the high values of such parameters as B_r and ВН_max for A-type magnets) also testify to this phenomenon. This follows from the magnetic moment of terbium atoms being directed antiparallel to the magnetic moments of cobalt and iron atoms in a phase lattice of the (Nd,Pr)₂Fe₁₄B type, while it has lower saturation magnetization and double the values ​​of the anisotropy field (~21 T). For comparison, the anisotropy fields for the Nd₂Fe₁₄B and Pr₂Fe₁₄B compounds were ~7 and ~9 T, respectively [1].

The more structurally sensitive parameters (_jH_c, H_k, and H_k/_jH_c) for C-type magnets compared to those of the B type are explained by the high content of terbium and aluminum in the boundary region of the main magnetic phase. A similar effect was observed earlier on sintered cobalt-free magnets of the (Nd,Pr,Tb)₂Fe₁₄B type, i.e., an increase in _jH_c when there was a gradient of terbium [3, 5] and/or aluminum [6].

CONCLUSIONS

It was shown that praseodymium stabilizes the structure of intermetallic compounds of the (Pr,Tb)₂(Fe,Co)₁₄B type with high contents of cobalt (up to 20 wt %). Adding TbAl₃ hydride during the manufacture of sintered magnets enhances the coercive magnetization force and improves the temperature stability of magnetic induction.