Hard magnetic properties of spacer-layer-tuned NdFeB/Ta/Fe nanocomposite films
Introduction
Nd–Fe–B-based permanent magnets have attracted a lot of attention recently, owing to their broadening applications in the traction motors of hybrid and pure electric vehicles and wind generators. A large maximum energy product (BH)max can be obtained in anisotropic Nd–Fe–B sintered magnet because of its high remanence and a coercivity higher than μ0Mr/2 along with good squareness. The typical value of the (BH)max of the commercial Nd–Fe–B magnet is ∼400 kJ m−3 with a coercivity of 1.2 T. A higher (BH)max of 475 kJ m−3 was reported for a laboratory-scale Nd–Fe–B sintered magnet with highly oriented grains and minimized non-ferromagnetic Nd-rich phases, but its coercivity was only 0.8 T, barely larger than μ0Mr/2, which is too low for most applications [1].
The increasing cost and scarcity of rare-earth (RE) elements has stimulated the development of permanent magnets using fewer RE elements. In 1988, Coehoorn et al. [2] developed isotropic Nd2Fe14B/Fe3B nanocomposite magnets with moderate coercivity and enhanced remanence by melt-spinning. The RE content of the alloy was only 4.5%, much less compared with ∼12.5–14% for commercial sintered magnets. However, the isotropic magnets cannot exhibit a higher (BH)max than that of anisotropic sintered magnets because of low remanence and poor squareness. For exchange coupling of soft and hard magnetic phases, there is an upper limit for the size of the soft-magnetic (SM) phase. In one dimension, the grain size of a SM phase must be less than the twice the domain wall width δh of a hard-magnetic (HM) phase [3]. Here, δh is estimated by δh = π(Ah/Kh)1/2. Taking Nd2Fe14B as the HM phase, Ah ≈ 10−11 J m−1, Kh = 4.5 × 106 J m−3 [4], 2δh is ∼8.4 nm. Based on a three-dimensional calculation, in exchange coupled Sm2Fe17N3/Fe67Co33 anisotropic nanocomposite magnets, a large (BH)max of 1 MJ m−3 has been optimistically predicted [5].
Many experiments on anisotropic nanocomposite magnets [6], [7] and their exchange coupling mechanism [8], [9] have been carried out. Compared with the bulk nanocomposite magnet, controlling nanostructure is easier in thin films. So many investigations on anisotropic nanocomposite magnet films were done as proof-of-principle experiments, e.g. Nd–Fe–B/Fe [10], [11] and SmCo/Fe [12], [13] multilayer films. However, the reported values of (BH)max were disappointing, owing to small coercivity and degraded squareness caused by poor anisotropic texture. The reason for the small coercivity was thought to be the initiation of magnetic reversal from exchange-coupled SM layers. So, attention was paid to how to improve the coercivity in NdFeB-based nanocomposite magnets. A coercivity of 0.8 T was reported for Mo-inserted NdFeB/Fe multilayer films [14], and the coercivity mechanism was studied as a result of the modified interfaces [15], [16]. The importance of the interface between SM/HM layers for coercivity was further demonstrated by the insertion of a thin Ta layer between Nd2Fe14B and FeCo layers. By controlling the interlayer, a high coercivity of 1.38 T and (BH)max of 486 kJ m−3 were achieved in a Nd–Fe–B/Ta/FeCo nanocomposite multilayer film [17]. The superior performance in the Ta-inserted Nd2Fe14B/Ta/FeCo nanocomposite enabled a feasible approach for a high-coercivity SM/HM nanocomposite. The direct contact between SM and HM layers will cause strong exchange coupling, which causes a dramatic reduction in the coercivity of a hard phase. However, the coercivity of Nd2Fe14B/FeCo films could be kept comparable with that of single-layer (SL) Nd–Fe–B film by introducing a 2-nm-thick Ta spacer layer as a result of the weakened SM/HM exchange coupling [17]. These results indicate that the Ta spacer layer plays a critical role between SM/HM layers, possibly leading to a different coupling mechanism. In the present work, simplified Nd2Fe14B/Ta/Fe trilayer model films with various Ta and Fe thicknesses were grown, to gain a better understanding of the coupling mechanism and demagnetization process in spacer-layer-tuned SM/HM multilayer systems.
Section snippets
Experiments
An alloy target with a composition of Nd13Fe77B10 was used to grow Nd–Fe–B hard magnetic layers. The base pressure was better than 10−6 Pa, and the Ar pressure was kept at 1.3 Pa. The film structure was Ti(20 nm)/NdFeB(100 nm)/Nd(10 nm)/Ta/Fe/Ti(20 nm). The Ti(20 nm) underlayer and cover layer were sputtered at room temperature to suppress oxidation. These two layers are omitted in the expression of the thin film structure for simplicity. The NdFeB(100 nm) layer was deposited at 600 °C to get a textured
Results
Fig. 1 shows the X-ray diffraction (XRD) patterns of NdFeB(100 nm)Nd(10 nm) SL film compared with those of NdFeB(100 nm)Nd(10 nm)/Ta(x nm)/Fe(10 nm) films (x = 2 nm, 10 nm, 100 nm, 250 nm). Strong (0 0 4) and (0 0 6) peaks are observed in the XRD of the SL film. A peak ∼2θ = 31° is from the Nd-rich phase, which is of double hexagonal close-packed structure. In the thin films with a Ta(2 nm) spacer layer and a Ta(10 nm) spacer layer, besides the (0 0 2 l) peaks, a minor peak corresponding to (1 1 4) of Nd2Fe14B phase
Discussion
Kinks are considered to be a sign that a SM phase cannot be fully coupled with a HM phase, and some free moments appear in the SM phase. As shown in Fig. 2, Fig. 4, along the IP direction, for a fixed Ta(2 nm) spacer layer, a kink appears when Fe thickness is increased from 5 nm (see Fig. 4a) to 10 nm (see Fig 2c). In contrast, for a fixed Fe(10 nm) layer, along the IP direction, the kink appears when Ta thickness is increased from 0 nm (see Fig. 2b) to 2 nm (Fig. 2c), indicating decoupling between
Conclusions
The anisotropic Ti(20 nm)/NdFeB(100 nm)Nd(10 nm)/Ta(x nm)/Fe(y nm)/Ti(20 nm) thin films were prepared with varying thickness of the Ta spacer layer and the Fe layer. The demagnetization process and recoil behaviors were studied experimentally. The effects of exchange coupling and magnetostatic coupling in NdFeB/Ta/Fe trilayer films on their HM properties were discussed based on the experimental observations as well as micromagnetic simulations. It was found that exchange coupling was dominant when
Acknowledgement
This research was partly supported by JST, CREST.
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2020, Journal of Rare EarthsCitation Excerpt :The high Br = 12.7 kG originates from the high-content α-Fe(Co) phase with a small grain size combined with the (00l) texture for the Nd2Fe14B and SmCo7 nanocrystals and the magnetic coupling of the obtained HPTC magnets.16,17,31,32 The magnetic coupling in the HPTC magnets may come from the combination of exchange coupling and magnetostatic coupling, where the strong exchange coupling takes place between the small nanocrystals of SmCo7 and α-Fe(Co) while the magnetostatic coupling takes place among the layers.17,31,32 The large Hci −6.1 kOe is mainly because that the achieved magnet possesses the Nd2Fe14B nanograins with the Nd-rich phase,33 the high magnetocrystalline anisotropic SmCo7 phase,34 and introduced layered interfaces, all of these provide strong pinning sites for domain-wall movements.