France



Introduction:&lt;br/&gt;Non-oriented silicon steel is widely used in the manufacture of motors and other equipments due to its excellent magnetic properties. In the production process, non-oriented electrical steel often needs to be punched to form the required shape. During the punching process the magnetic properties of silicon steel sheets will be deteriorated due to the mechanical stress [1-4]. As a result, the plastic deformation and residual stress emerge around the area of cutting edge, which have a significant impact on the magnetic properties and magnetic domain structure of electrical steel sheets.&lt;br/&gt;In this paper the effects of punching on the magnetic properties and magnetic domain structure of non-oriented electrical steel are studied. And the deterioration degree of the magnetic properties caused by different punching load are analyzed quantitatively.&lt;br/&gt;Ι Test samples and methods&lt;br/&gt;The test samples of cold-rolled non-oriented electrical steel B35A270 are punched on the precision punching machine by the 25-ton and 80-ton punching load, respectively. Then, two kinds of samples with a size of 49.5*49.5*0.35mm are obtained. The sample processed by 25 tons of punching load is denoted as sample A, and the one processed by 80 tons of punching load is denoted as sample B.&lt;br/&gt;The 2D magnetic properties measurement apparatus consists of two single C-shaped yokes, excitation windings and silicon steel sheet samples. An AC magnetic field with a frequency of 50 ~ 200Hz is applied along the rolling direction of the sample The samples supported by the sample holder are placed in the gap between the two yokes, and the yoke air gap matches the size of the sample. The sensing system is specially designed to measure the local magnetic properties in different position of the sample. In the sensing system, the micro-size &lt;i&gt;H &lt;/i&gt;coil is wound on a 0.05mm thick PCB substrate with a length of 3mm to measure the &lt;i&gt;H&lt;/i&gt; signal in the rolling direction of the sample. Two retractable &lt;i&gt;B&lt;/i&gt; probes are fixed on a 1.5mm thick PCB with a spacing of 3.5mm to measure the &lt;i&gt;B&lt;/i&gt; signal in the rolling direction of the sample.&lt;br/&gt;Taking the midline of the sample along the tangential direction as the measurement characteristic sampling line, the micro-size composite &lt;i&gt;B&lt;/i&gt;-&lt;i&gt;H&lt;/i&gt; sensor can be moved smoothly along the characteristic sampling line with a 2 mm distance step.&lt;br/&gt;ΙΙ Experimental results and analysis&lt;br/&gt;The core loss curves of the two samples measured at 50Hz sinusoidal signals are depicted in Fig.1 (a) and (b). The loss of the sample A increases by 0.9% ~ 15.25% in the range of 8 mm from the edge, especially in the range of 6 mm. The loss of the sample B increases by 6.8 % ~ 29.55%, while the deteriorated area expands to 14 mm.&lt;br/&gt;According to the Bertotti loss separation, the hysteresis, eddy current and excess loss coffficient are identified by the experimental results. It is concluded that the punching process cause an increase of the hysteresis loss. It can be attributed to plastic deformation and residual stress around the punching edge. The moving of magnetic domains can be hindered with the increase of the width of plastic deformation area. During the magnetization process, the movement resistance of the edge domain increases, so the hysteresis loss increases, which increases the iron loss of the sample.&lt;br/&gt;At the edge of the sample A, the magnatic coercivity increases by 5.5%, the magnetic remanence and permeability decreases by 11.7% and 29.2%. The magnatic coercivity increases by 24.3%, the remanence and the permeability decreases by 35.9% and 47.9% at the edge of the sample B.&lt;br/&gt;ΙΙΙ Observation of magnetic domain structure&lt;br/&gt;In order to investigate the mesoscopic properties of punching effect, the 8*8mm sample need to be prepared for Kerr microscope, which is cut by wire electrical discharge machining from the above sample A and B in the edge and center area of the samples.After sanding, mechanical polishing and other steps, shiny mirror-like samples are obtained. Using a Kerr microscope to observe the magnetic domains of the sample, as shown in Fig. 2(a), the center of the sample A is dominated by sheet-like magnetic domains, while there are a large number of labyrinth domains near the punching edge as shown in Fig. 2 (b). The size and shape of the domains also change. In addition, the magnetic domain in the center is parallel to the rolling direction, while the direction of the magnetic domain at the edge changes, which makes it difficult to magnetize. This leads to the increase of iron loss and the deterioration of magnetic properties.&lt;br/&gt;IV Conclusion&lt;br/&gt;Based on the local 2D magnetic measurement and magnetic domain observation, the magnetic properties of non-oriented electrical steel deteriorate by the punching processing. Furthermore, the deterioration degree of the magnetic properties become more significant by larger punching load.

**References:**

[1] Campos M F D , Sablik M J, and Landgraf F J G . Journal of Magnetism &amp; Magnetic Materials., Vol.320, p.e377-e380(2008)&lt;br/&gt;[2] René Siebert, Wetzig A, and Beyer E. IEEE Electric Drives Production Conference., (2013)&lt;br/&gt;[3] Nakata T. IEEE Translation Journal on Magnetics in Japan., (2008)&lt;br/&gt;[4] Senda K , Ishida M, and Nakasu Y. Journal of Magnetism and Magnetic Materials., Vol.304, p.e513-e515, (2006)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/2e1b0ece638ad406cd008f90715dcef9.jpg)
Fig.1 Core loss curves of sample A and B

![](https://assets.underline.io/uploads/markdown_image/1/image/b2c57ff57baf50342830843dfc45b66b.jpg)
Fig.2 Magnetic domain structure of sample A

INTERMAG 2021

 Effect of Punching on Magnetic Properties of Non-Oriented Electrical Steel

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Introduction: Non-oriented silicon steel is widely used in the manufacture of motors and other equipments due to its excellent magnetic properties. In the production process, non-oriented electrical steel often needs to be punched to form the required shape. During the punching process the magnetic properties of silicon steel sheets will be deteriorated due to the mechanical stress [1-4]. As a result, the plastic deformation and residual stress emerge around the area of cutting edge, which have a significant impact on the magnetic properties and magnetic domain structure of electrical steel sheets. In this paper the effects of punching on the magnetic properties and magnetic domain structure of non-oriented electrical steel are studied. And the deterioration degree of the magnetic properties caused by different punching load are analyzed quantitatively. Ι Test samples and methods The test samples of cold-rolled non-oriented electrical steel B35A270 are punched on the precision punching machine by the 25-ton and 80-ton punching load, respectively. Then, two kinds of samples with a size of 49.5*49.5*0.35mm are obtained. The sample processed by 25 tons of punching load is denoted as sample A, and the one processed by 80 tons of punching load is denoted as sample B. The 2D magnetic properties measurement apparatus consists of two single C-shaped yokes, excitation windings and silicon steel sheet samples. An AC magnetic field with a frequency of 50 ~ 200Hz is applied along the rolling direction of the sample The samples supported by the sample holder are placed in the gap between the two yokes, and the yoke air gap matches the size of the sample. The sensing system is specially designed to measure the local magnetic properties in different position of the sample. In the sensing system, the micro-size H coil is wound on a 0.05mm thick PCB substrate with a length of 3mm to measure the H signal in the rolling direction of the sample. Two retractable B probes are fixed on a 1.5mm thick PCB with a spacing of 3.5mm to measure the B signal in the rolling direction of the sample. Taking the midline of the sample along the tangential direction as the measurement characteristic sampling line, the micro-size composite B-H sensor can be moved smoothly along the characteristic sampling line with a 2 mm distance step. ΙΙ Experimental results and analysis The core loss curves of the two samples measured at 50Hz sinusoidal signals are depicted in Fig.1 (a) and (b). The loss of the sample A increases by 0.9% ~ 15.25% in the range of 8 mm from the edge, especially in the range of 6 mm. The loss of the sample B increases by 6.8 % ~ 29.55%, while the deteriorated area expands to 14 mm. According to the Bertotti loss separation, the hysteresis, eddy current and excess loss coffficient are identified by the experimental results. It is concluded that the punching process cause an increase of the hysteresis loss. It can be attributed to plastic deformation and residual stress around the punching edge. The moving of magnetic domains can be hindered with the increase of the width of plastic deformation area. During the magnetization process, the movement resistance of the edge domain increases, so the hysteresis loss increases, which increases the iron loss of the sample. At the edge of the sample A, the magnatic coercivity increases by 5.5%, the magnetic remanence and permeability decreases by 11.7% and 29.2%. The magnatic coercivity increases by 24.3%, the remanence and the permeability decreases by 35.9% and 47.9% at the edge of the sample B. ΙΙΙ Observation of magnetic domain structure In order to investigate the mesoscopic properties of punching effect, the 8*8mm sample need to be prepared for Kerr microscope, which is cut by wire electrical discharge machining from the above sample A and B in the edge and center area of the samples.After sanding, mechanical polishing and other steps, shiny mirror-like samples are obtained. Using a Kerr microscope to observe the magnetic domains of the sample, as shown in Fig. 2(a), the center of the sample A is dominated by sheet-like magnetic domains, while there are a large number of labyrinth domains near the punching edge as shown in Fig. 2 (b). The size and shape of the domains also change. In addition, the magnetic domain in the center is parallel to the rolling direction, while the direction of the magnetic domain at the edge changes, which makes it difficult to magnetize. This leads to the increase of iron loss and the deterioration of magnetic properties. IV Conclusion Based on the local 2D magnetic measurement and magnetic domain observation, the magnetic properties of non-oriented electrical steel deteriorate by the punching processing. Furthermore, the deterioration degree of the magnetic properties become more significant by larger punching load.

**References:**

[1] Campos M F D , Sablik M J, and Landgraf F J G . Journal of Magnetism & Magnetic Materials., Vol.320, p.e377-e380(2008) [2] René Siebert, Wetzig A, and Beyer E. IEEE Electric Drives Production Conference., (2013) [3] Nakata T. IEEE Translation Journal on Magnetics in Japan., (2008) [4] Senda K , Ishida M, and Nakasu Y. Journal of Magnetism and Magnetic Materials., Vol.304, p.e513-e515, (2006)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/2e1b0ece638ad406cd008f90715dcef9.jpg)
Fig.1 Core loss curves of sample A and B

![](https://assets.underline.io/uploads/markdown_image/1/image/b2c57ff57baf50342830843dfc45b66b.jpg)
Fig.2 Magnetic domain structure of sample A

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The influence of Ce3+ ions substitution for Fe3+ ions in Ni0.42Zn0.58CexFe2-xO4 (x = 0, 0.02, 0.04, 0.06, 0.08, 0.10) compounds on the selected magnetic properties and structure were investigated. The samples were synthesized by means of wet glycine-nitrate technology annealed at the temperature 1200°C for 6 hours. The wide spectrum of characterization methods were used for the study of relevant properties of synthetized materials. X-ray diffraction patterns were recorded by a diffractometer using conventional X-ray tube (CuKα radiation, λ = 1.5418 Å). EDAX were performed by the SEM JEOL 7600-F equipped with both EDX and WDX system. Temperature dependences of magnetic susceptibility was measured using bridge method. The hysteresis loops were recorded using computer controlled digital hysteresisgraph at the same amplitude of applied field Hmax = 1500 A/m, measured at the frequency f = 50 Hz and at room temperature. The determination of real µ' and imaginary µ'' parts of complex (relative) permeability have been performed by means of an impedance analysis procedure using a vector network analyser Keysight E5063A in the frequency range from 1 MHz to 3 GHz. Results and Discussion XRD results indicate a presence of multi-phase structure. The major is found to be spinel phase and secondary, increasing with the addition of Ce3+ ions, corresponds to CeO2. The final number of possible substitution of Ce3+ ions for Fe3+ ions in the B sites, and the difference in their ionic radii, is the cause of the formation of aggregates on the grain boundary as secondary phase [1]. The distribution of elements and their amounts in weight percentages for NiZnCe samples were studied by means of EDAX analysis. The real content of Ni2+, Zn2+, Ce3+, Fe3+ and O2- shown in Tab. I confirmed the substitution of Fe3+ ions by Ce3+ ions, since the increase of Ce3+ content decreases the amount of Fe3+. Temperature dependences of magnetic susceptibility were recorded within the range from room temperature up to 450°C at rise of 4°C/min. Ferrite powder specimens of amount approximately 50 mg were measured; susceptibility values were related to the same weight (1 mg) for all cases. These dependences are decreasing monotonously and they exhibit one strong drop to paramagnetic state. From that, we can conclude that no unwilled phase is present in the ferrite to a significant extent. Curie temperature TC was determined from the point of inflexion at strong decrease of temperature dependence. Even the paramagnetic substituent (Ce3+) ions occupy B-sites dominantly at expense of Fe3+ ions, it does not change the A-B exchange interactions, because TC remains almost constant [2]. The hysteresis loops exhibit quite complicated behavior. By adding larger amount of Ce3+ ions, the maximum flux density Bmax at given maximum applied field sometimes raises and sometimes decreases comparing to the non-substituted ferrite. Note that the value of Bmax at given maximum applied field Hmax = 1500 A/m approaches the largest value for x = 0.08 (only slightly larger than for x = 0.04). The differences can be explained by the fact that the saturation magnetization is an intrinsic parameter (thus, structurally independent) meanwhile the Bmax corresponds to structurally sensitive behavior of material when magnetizing and approaching the saturation. In addition, coercive field is nearly linearly increasing with increment of x. This behavior is similar to that observed in case of Cu-substituted ferrites with the similar chemical formula (except for Cu instead of Ce), [3]. This may be justified by the complexity of the spinel structure, presence of other crystalline phases, etc. In Fig. 1 the frequency dependences of real µ' and imaginary µ'' parts of complex (relative) permeability µ = µ' - jµ'' are presented for prepared spinel ferrites. The obtained frequency dispersion of complex permeability is caused by the domain wall and ferrimagnetic (or spin precession) resonance [4], [5]. The achieved results may be explained as follows. Magnetic properties of nickel-zinc ferrites as representatives of spinel-like ferrites are affected not only by the method of preparation and chemical structure (and composition), but also by properties such as porosity, crystallite size (or grain size), lattice impurities (or defects), concentration of cations, etc. Also the super exchange interactions may strongly change the resultant properties: in spinel ferrite system AA and BB interactions are poorer than AB interaction. In the prepared ferrites, the shrinkage of lattice is present because of the larger ionic radius of cations such as Ce3+ (1.143Å), Zn2+ (0.74Å) and Ni2+ (0.69Å) in comparison with Fe3+ (0.65Å). The magnetic moments of these cations are as follows (in Bohr magnetons): 2.5µB, 0µB, 2.3µB and 5µB. Values of magnetic moments relates the cation distribution of Ni0.42Zn0.58CexFe2-xO4 ferrite system (Ce3+ ions have preference on B sites): (Zn0.582+Fe1-x3+)A[Ni0.422+Cex3+Fe1+x3+]BO42-. Due to the values of ionic radius and magnetic moment, Ce3+ ions replace more Fe3+ ions in the crystal lattice and this leads to the raise of coercive field which in turns changes the value of permeability [5], [6]. Conclusion From the above experimental results, it can be seen that adding even a small amount of cerium element has evident influence on selected properties. Consequently, by adding defined amount of Ce3+ ions, the desired parameters can be easily controlled by a straightforward way and thus tailored for any particular application required in industrial practice.

**References:**

[1] R. S. Yadav, J. Havlica, J. Masilko, J. Magn. Magn. Mater., Vol. 399, pp. 109-117 (2016) [2] M. A. Ahmed, S. F. Mansour, S. I. El-Dek, Solid State Ionics, Vol. 181, pp. 1149-1155 (2010) [3] J. Sláma, A. Grusková, M. Ušáková, J. Magn. Magn. Mater., Vol. 321, pp. 3346-3351 (2009) [4] M. Šoka, R. Dosoudil, M. Ušáková, Acta Phys. Pol. A, Vol. 131, pp. 690-692 (2017) [5] R. Dosoudil, M. Ušáková, A. Grusková, IEEE Trans. Magn., Vol. 50, no. 4 (2014) [6] M. N. Akhtar, M. A. Khan, J. Magn. Magn. Mater., Vol. 460, pp. 268-277 (2018)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/eccf1c3546b0486f5775ad28c48f37bb.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/2358bea70a2a67441e24abb7319de58f.jpg)

 Magnetic and Structural Properties Analysis of Cerium Substituted Nickel Zinc Ferrites

Rare heart substituted iron garnets (REIG) are characterized by a high magnetic susceptibility that results in a large magnetization in the presence of small magnetic fields [1]. The magnetic microstructures are influenced by natural or artificial defects; in particular, domain walls (DWs) undergo to nucleation, segmentation and pinning in correspondence to them. In REIG with in-plane spontaneous magnetization, the Néel spike domain structure is one of the most common magnetic microstructure nucleated at microdefects and its charged DWs can produce not negligible stray fields above the REIG. These DWs are interesting both for fundamental studies and for innovative applications such as micro positioning by contactless actuators [2]. To make the most out of this, performing a quantitative analysis is crucial. The experimental characterization by means of quantitative magneto-optical imaging with an indicator film [3] showed that stray fields are detectable above the surface of (Lu, Bi) substituted iron garnets, characterized by a spontaneous in-plane magnetization around 150 kA/m, in correspondence of DWs delimiting a Néel spike that is generated by a micro-crack [4]. In order to understand how the externally detected stray fields are correlated to the pattern of the internal magnetization, here we propose 2D and 3D models based on FEM developed through the COMSOL Multiphysics© software. These models take into account the magnetic microstructure of the Néel spike, i.e., either single DWs that run parallel into the bulk of the domain and a couple of DWs at the Néel spike tip (2D) or DWs that merge in a single point (3D). This approach allows us to describe the magnetic microstructure by means of CAD defined patterns, whose complex structure can be implemented easily in the FEM simulation. It turns out that the results are in agreement both with experiments and with previous numerical calculations performed with a simplified model in MATLAB© [4]. The comparison with experiments allowed us also to determine the distance of measurement between the indicator film of MOI experiments and the REIG under investigation. Figure 1 shows linear profiles across the DWs, along with experimental data. The distance of measurement is about 7 μm and the magnetic induction at the tip of the Néel spike is twice the value found for the single charged DW at the domain boundary. The simplicity of the 2D model limits the computational cost required; on the other hand, the 3D simulation can be useful for reproducing in detail the stray field distribution generated by the complex structure of the Néel spike domain, as shown in Figure 2. In summary, we validated these models for reproducing the stray field distribution above charged DWs, opening the way for the design of artificial structures in order to obtain controlled magnetic patterns, thanks to the powerful 3D engine of the COMSOL Multiphysics© software.

**References:**

[1] L. E. Helseth, M. Kurniawan, R. W. Hansen, and T. H. Johansen, PHYSICAL REVIEW B77, 064433 (2008) [2] L. E. Helseth, R. W. Hansen, T. H. Johansen, and T. M. Fischer, Appl. Phys. Lett.85, 2556 (2004) [3] F. Laviano, D. Botta, A. Chiodoni, R. Gerbaldo, G. Ghigo, L. Gozzelino, S. Zannella and E. Mezzetti, 16, (2003) 71; Physica B 403 (2008) 293 [4] A. Napolitano, C. Ragusa, S. Guastella, S. Musumeci, P. Rivolo, F. Laviano, JMMM, Vol 504, 2020, 166556 [5] O. Khan, C. Ragusa, and F. Khan, IEEE transactions on magnetics, vol. 50, no. 11, November 2014

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/7f0da98d3c563c66b54b53692a817a42.jpg)
Figure 1. Cross sectional profiles of the stray field, simulated and experimental values, for a single DW (a) and the tip of the Nèel spike (b).

![](https://assets.underline.io/uploads/markdown_image/1/image/2928dd60fc86135c15716d29f81cc2da.jpg)
Figure 2. Magnetic field distribution of the stray field above a Néel spike: 3D plot of the magnetic induction intensity that clearly shows the stray field peak in correspondence to the tip of the Nèel spike. The distance of measurement was chosen on the basis of the previous comparison between experimental data and 2D simulations.

 Modeling stray-field distribution generated by domain-walls in rare-earth substituted iron garnets

Intensive research on soft magnetic nanoparticles has been attracted by biomedical applications such as MRI markers and cancer therapy by magnetic hyperthermia. These nanoparticles can also be used in additive manufacturing processes, such as inkjet printing, to elaborate magnetic coatings to meet emerging needs such as RFID antenna shielding or wireless power transfer. The nanoparticles are synthesized by thermal decomposition with organic solvents. Two different ways of heating are employed: the conventional one using a heating mantle and the other one using a microwave oven[1]. A good control of the particle size of Fe3O4 through the addition of ethylene glycol has been reported2. Other groups have used this technique to synthesize core-shell Fe3O4 nanoparticles. We have synthetized nanoparticles of mixed oxides ZnxFe3-xO4, where the magnetic and dielectric properties can be tuned by the Zn/Fe ratio[2]. The two processing methods were compared in order to determine the most appropriate way to control the size and morphology spread of the obtained nanoparticles. The influence of the microstructure on the magnetic properties of Zn0.4Fe2.6O4 and Zn0.2Fe2.8O4 (theoretical compositions) nanoparticles was studied by X-ray diffraction analysis and 57Fe Mössbauer spectrometry, under zero field and in-field. Experimental details: The conventional way in a heating mantle allows obtaining various morphologies: cubes, octahedra... However, it induces heterogeneity in terms of size and shape in a batch, due to insufficient monitoring of the temperature during the reaction. The microwave-assisted thermal decomposition enables a fast and more uniform heating, leading reproducibly to nanoparticles with homogeneous size and shape. The difference lies in the choice of solvents and surfactants: respectively benzyl ether and oleic acid for the conventional thermal decomposition, and octadecene for the solvent with a combination of oleic acid and oleylamine for the surfactant for the microwave reactor. These surfactants have two complementary functions in the synthesis. Oleic acid stimulates the growth of nanoparticles along a preferential axis while oleylamine, in addition to redox possibilities, slows down the nanoparticle growth[3]. Tuning their compositional ratio yields a good control of the nanoparticle morphology. Results: The morphology and the size were observed by scanning electron microscopy and transmission electron microscopy. Two different morphologies are highlighted. While 5 nm-nanoparticles (written SP(F)M) are quite spherical, the 35 nm-particles (FM) have a “raspberry” morphology (Fig. 1). These observations are consistent with those of Y-J. Liang’s1 on 50 nm particles. It should be assimilated to a binding between crystallites and the continuity of the crystal planes in a raspberry structure is noticed, as observed by S. Begin-Colin’s group[4] with structures of about 250 nm. The phase purity was probed by X-ray diffraction (XRD). The size of the coherent diffraction domains was calculated with two different methods: using Scherrer’s formula, and through a Rietveld refinement. The results are compared and correlated to TEM observations. The saturation magnetization was determined through the hysteresis cycles measured with a vibrating sample magnetometer (VSM) at room temperature. The superparamagnetism is observed for the 5 nm particles. The magnetization curve has been fitted with a Langevin function[5] which allows recovering the intrinsic saturation magnetization and an apparent size of 5 nm, close to that observed by TEM. The intrinsic saturation magnetization (85.6 emu/g) from the model is in good agreement with that determined at 2 K with a VSM/SQUID (85 emu/g). The saturation magnetization of the two types of nanoparticles is lower (85.6 emu/g for the 5 nm and 80.6 emu/g for the 35 nm) than that of the bulk ferrite Zn0.4Fe2.6O4 which is about 98.8 emu/g. Several contributions have to be taken into account to explain this discrepancy: - A surface spin-canting effect[6] could impact the magnetic moment of particles, especially for the smallest ones. The thickness of the surface spin-canting was estimated and the magnetic moments were thus corrected (Fig. 2). - A cationic disorder with a fraction of Zn2+ ions at the octahedral sites of the spinel structure, and not only at the tetrahedral (A) ones, can be deduced from a Rietveld analysis of X-ray patterns, and in-field Mössbauer spectra. Indeed, these techniques have made it possible to quantify the distribution of cations at the tetrahedral (A) and octahedral (B) sites of the spinel, and thus to calculate the magnetization according to a molecular field model[7] (Fig. 2). - Oxidation of the nanoparticles with conversion of Fe2+ into Fe3+ could lead to the reduction of the saturation magnetization[8]. The analysis of Mössbauer spectra enables to recover values close to the measured magnetic moments, which makes the cations distribution reliable. It therefore makes it possible to predict a magnetic moment close to the experimental one, keeping a spinel-type chemical formula with presence of vacancies to take into account the oxidation of the theoretical composition. The understanding of the effect of the microstructure on the magnetic moment opens up possibilities by tuning magnetic properties of iron oxide nanoparticles.

**References:**

[1] Y-J. Liang, F. Fan and M. Ma, Colloids and Surfaces A., Vol. 530, p.191-199 (2017) [2] A. Fujiwara, M. Tada and T. Nakagawa, J. Magn. Magn. Mater., Vol. 320, p.L67-L69 (2008) [3] S. Mourdikoudis and L. M. Liz-Marzán, Chem. Mater., Vol. 25, no. 9, p.1465–1476 (2013) [4] O. Gerber, B. Pichon, and C. Ulhaq, J. Phys. Chem. C., Vol. 119, p.24665-24673 (2015) [5] J.L. Dormann, D. Fiorani and E. Tronc, Adv. Chem. Phys., Vol. 98, p.283-494 (1997) [6] E. Tronc, P. Prené, J.P. Jolivet, J.L. Dormann and J-M. Greneche, Hyper. Inter., Vol. 112, p.97-100 (1997) [7] P. G. Bercoff and H. R. Bertorello, J. Magn. Magn. Mater., Vol. 169, p.314–322 (1997) [8] A. Demortiere, P. Panissod and B.P. Pichon, Nanoscale., Vol. 3, p.225-232 (2011)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/d0fab6a72fe682a60b495b03c6b78453.jpg)
Fig. 1. TEM picture of raspberry-like nanoparticles

![](https://assets.underline.io/uploads/markdown_image/1/image/3f6da719018112feef39b97921fb4966.jpg)
Fig. 2. Magnetic moments predicted from microstructural characterizations for SP(F)M and FM particles with different Zn contents (0.2 and 0.4).

 Quantitative retrieving of the magnetic moment of iron oxide nanoparticles through structural characterizations

A new intermetallic ternary compound La6Pd2.28Sb15 was synthesized from the reaction of lanthanum and palladium metals in a molten antimony flux. The compound crystallizes in the orthorhombic space group Immm with a= 4.3082(9) Å, b=15.399(3) Å, and c=19.689(4) Å. The crystal structure contains a three-dimensional framework of Sb squares and ribbons that extend along the a axis, including complex Sb–Sb bonding. Magnetization measurements indicate that the intrinsic magnetic behavior of La6Pd2+xSb15 (x = 0.28) is diamagnetic, which we associate with Landau diamagnetism slightly larger than Pauli paramagnetism in this compound. The resistivity (ρ) decreases when lowering the temperature, indicating metallic behavior, and at low temperatures r depends quadratically on T. Interestingly, both the Hall resistivity and magnetoresistance present a nonlinear dependence on the applied magnetic field, suggesting a multiband behavior. This is supported by density-functional electronic structure calculations which show a complex Fermi surface originated in the antimonide substructures and containing both electron and hole pockets as well as open sheets.

**References:**

M. I. Sturza, M. L. Amigó, J. I. Facio et al., Journal of Solid State Chemistry, Vol. 291, pp. 121578 (2020),

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/809fee0358e093416661fd556c5c2f76.jpg)
Fig 1. Ball-and-stick representations of three-dimensional structure of La6Pd2+xSb15 (x = 0.28) viewed down the crystallographic a-axis. The red dashed lines outline the La6 trigonal prisms.

![](https://assets.underline.io/uploads/markdown_image/1/image/0a2b239811e304dece46ec3eebd0d4db.jpg)
Fig 2. Temperature dependence of the susceptibility for zero field cooling (ZFC) and field cooling (FC) procedures. Inset: Field dependence of the magnetization for several temperatures. In all cases, the background of the glue and the sample holder was subtracted, as well as, the core diamagnetic contribution.

 La6Pd2+xSb15 (x = 0.28): A Rare-earth Palladium Intermetallic Compound with Extended Pnictogen Ribbons

Heusler alloys have been intensively investigated to develop functional materials as half-metals, shape memory alloys, and thermoelectric conversion materials. Heusler alloy Co2TiZ (Z = Si, Ge, Sn) was expected to be a half-metal with characteristics of Weyl fermions by first-principles band calculations [1]. The Weyl semimetals are of great interest materials, due to their unusual magnetoelectronic and magnetothermal properties. In the case of magnetic materials, the spatial and temporal correlations of the magnetization M(r,t) show a tendency to diverge at the critical temperature. The phenomenon in which magnetic properties show peculiar temperature and magnetic field dependence at the critical temperature is called a critical phenomenon. The critical exponent δ is defined at T = TC by H = DM(T,H)δ . (1) In later work, to find the critical exponent δ, the temperature at which the plot, taking logarithms on both sides of equation (1), is linear in log(M)-log(H) scale is defined as the Curie temperature TC [2]. Theoretically, the molecular field theory predicts δ = 3.0, the three-dimensional Heisenberg model δ = 4.80 [2], the three-dimensional Ising model δ = 4.82 [2], and the spin fluctuation theory by Takahashi δ = 5.0 [3]. For half-metals with characteristics of Weyl fermions, there are few reports of systematic analyses of the critical exponents. Therefore, we focus on the element substitution effect between Co2TiSn and Co2TiGa. Co2TiSn possesses characteristics of both half-metal and Weyl fermions, and Co2TiGa is the itinerant electron ferromagnet with the high spin polarization. We measured the magnetization in the mixed crystal compound Co2TiGa0.8Sn0.2 and investigated the critical exponent δ near TC. Polycrystalline Co2TiGa0.8Sn0.2 samples were synthesized by arc-melting high-purity constituent elements under high purity Ar atmosphere. To homogenize the synthesized sample, it was annealed at 1100°C for 3 days and then quenched in water. The magnetization of Co2TiGa0.8Sn0.2 was measured using a superconducting quantum interference device (SQUID) (Quantum Design MPMS-XL). From the magnetization measurements of Co2TiGa0.8Sn0.2, the spontaneous magnetization Ms was estimated to be Ms = 0.907 μB/f.u. The critical exponent δ was analyzed from the magnetic field dependence of the magnetization near TC using equation (1). The magnetization curves obtained at each temperature were plotted for log(M/Mmax) against log(H/Hmax) in Fig.1. The values of δ and its fitting error, obtained from the analysis of the magnetization curves in Fig.1, are shown in Fig.2. Since the temperature at which the fitting error is minimized agrees best with equation (1), that temperature can be considered as the Curie temperature TC. From Fig.2, TC = 105 K and δ = 4.45 were obtained. The δ value obtained in this study was different from the value 3 by the molecular field theory, but was close to the value 5 predicted by the spin fluctuation theory for a weakly ferromagnetic itinerant electron system. Therefore, we also reanalyzed the magnetization below 10 kOe, where the magnitude of the applied magnetic field is 1/5 of 50 kOe, because the strong magnetic field may suppress the spin fluctuations. From the analysis of the data below 10 kOe, TC = 104 K and δ = 4.82 were obtained. Thus, a change of δ due to the suppression of the critical fluctuation by the strong applied magnetic field seems to be likely at least in this case. The discrepancy between our results and the theoretical models seem to arise from a distribution of the Curie temperatures in the samples since the δ value changes even at 1 K.

**References:**

[1] G. Chang et al., Sci. Rep. 6, 38839 (2016). [2] M. Seeger et al., Phys. Rev. B 51, 12585 (1995) and References: cited therein. [3] K. Shimizu et al., J. Phys. Soc. Jpn. 59, 305 (1990).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/bc2f95d7cb6fc566aa6579d1bb42e012.jpg)
Fig.1. The log(M/Mmax)-log(H/Hmax) plots of Co2TiGa0.8Sn0.2 near TC.

![](https://assets.underline.io/uploads/markdown_image/1/image/829df8173ea5fbd2b9218fdb9ac04f21.jpg)
Fig.2. The critical exponent δ and its fitting error of Co2TiGa0.8Sn0.2.

 Critical behavior of the magnetization in Heusler alloy Co2TiGa0.8Sn0.2

Mn2YZ based inverse Heusler alloys with cubic or tetragonal structures and high Curie temperature (TC) display exceptional tunability of their properties either by controlling the degree and the type of atomic ordering [1] or by controlling the total valence electron concentration (VEC) of the alloy [2]. Here the Mn atoms occupy both the (0,0,0) or (½, ½, ½) and (¼, ¼, ¼) or (¾, ¾, ¾) sites in the Heusler structure with antiparallel spins of different magnitude resulting in a ferrimagnetic order. Several theoretical and experimental studies on Mn2YZ alloys have predicted interesting properties like half-metallicity, spin gapless semiconductor with a high magnetic ordering temperature, anomalous hall effect and even zero field exchange bias, spin valve behaviour [3,4] etc. Mn2NiSn is one such cubic ferrimagnetic Heusler alloy with a moderately high magnetic moment of ~2.5 mB [5]. Density functional calculations as well as neutron diffraction studies have shown that the magnetic moment of Mn2NiSn can be tuned by controlling antisite disorder [5,6]. There are also reports suggesting a strong relation between the magnetic moment and the VEC controlled essentially by the Z element in Mn2NiZ alloys [7]. In the present work we aim to understand the role of antisite disorder on the magnetic properties of Mn2NiSn. We have employed two routes to introduce the antisite disorder. Firstly, we have systematically replaced Ni by Mn at the X site in Ni2MnSn by preparing alloys of the type Ni2-xMn1+xSn (where x = 0, 0.25, 0.5, 0.75 and 1) Secondly, the two alloy compositions, Ni1.25Mn1.75Sn (x = 0.75) and Mn2NiSn (x = 1) were subjected to an additional thermal treatment. The alloys were prepared by arc melting constituent elements in argon atmosphere and subsequently annealed at 800 oC for 72 hours in evacuated quartz tubes and quenched in ice cold water (referred to as rapid quenched (RQ)). The arc melted Ni1.25Mn1.75Sn and Mn2NiSn were cut into two halves and were sealed in four different quartz tubes. While two of these tubes containing the respective alloys were quenched from 800oC in ice water, the temperature of the other two tubes with the same compositions as that of the RQ alloys was lowered from 800oC to 427oC after 72 hours of annealing and were annealed at this lower temperature for the next 28 hours and cooled in furnace (referred as FC). All the alloys were characterized by x-ray diffraction (XRD) and SEM-EDX for the determination of their structure and composition. SEM-EDX indicated the compositions to be within 1% of their starting values. Rietveld refinement of the XRD patterns for the RQ alloys showed them to be largely single phase with a cubic L21 Heusler structure (SG: Fm-3m). Small percentage (~ 1 – 2 %) of the MnO impurity phase was noticed in a few compositions. In the case of FC alloys, while Ni1.25Mn1.75Sn alloy was cubic, Mn2NiSn was biphasic with major L21 phase (~93%) and minor (~7%) Mn3Sn type hexagonal phase (SG: P63/mmc) (Fig 1) . Magnetization measurements as a function of temperature (3 K to 390 K) were carried out in 100 Oe applied magnetic field employing zero-field cooling and field cooling protocols. TC increased from 340 K in x = 0 to above the highest measured temperature for all alloys with x =0.5 and above . Isothermal magnetization at 5 K was measured in the field interval of ± 7 T to determine the saturation magnetic moment. All the compositions showed ferromagnetic like M(H) loop with very low coercivity. The saturation magnetization in Ni2MnSn gave an effective moment of 4.06 mB/f.u. good agreement with reported values [8]. In the RQ alloys, with increase in Mn content, the magnetic moment decreases steadily from 4.06 mB/f.u. to 1.93 mB/f.u. in Ni1.25Mn1.75Sn and then increased to 2.80 mB/f.u. in NiMn2Sn. The FC Ni1.25 Mn1.75Sn showed a significant increase in its magnetic moment to 3.31 mB/f.u. This increase is surprising considering its structure to be the same as that of its RQ counterpart. In the phase separated Mn2NiSn the magnetic moment showed a marginal increase to 2.83 mB/f.u. To understand the above magnetic properties, local structures of Ni, Mn and Sn in all the alloys were studies by employing extended x-ray absorption fine structure (EXAFS) spectroscopy. In all the alloys EXAFS data were analysed together using a model based on their respective crystal structure. Excellent fits were obtained for RQ alloys with x = 0, 0.25 and 0.5. In case of Ni1.25Mn1.75Sn, only Ni EXAFS could be well described with the crystal structure model. Both Mn and Sn EXAFS could not be fitted and relaxation of structural constraints indicated a much lower value for the Mn – Sn bond distance than the predicted 3.025 Å from its crystal structure. Introduction of additional Mn – Sn and Mn – Mn correlations derived from hexagonal Mn3Sn resulted in a good fit to the EXAFS data of x = 0.75 and 1 RQ and FC alloys. Mn3Sn like environment can exist in Mn2NiSn type alloys if, in addition to the A2 type disorder, there is also an antisite disorder between Mn and Sn atoms (Fig 2). EXAFS analysis indicated a defect concentration ~ 20 ± 7 % in RQ and FC Mn2NiSn and much less in Ni1.25Mn1.75Sn. Such a segregation of defect within the overall Heusler structure due to antisite disorder seems to be the reason for the increased magnetic moment in Mn2YZ alloys.

**References:**

[1] T. Graf, C. Felser, S. Parkin, Prog. Solid State Chem. 39, 1 (2011). [2] C. Felser, L. Wollmann, S. Chadov, Chem Int Ed. 46:668-99, (2007). [3] S. Ouardi, G. Fecher, C. Felser, Phys. Rev. Lett. 110, 100401 (2013). [4] A. K. Nayak, M. Nicklas, S. Chadov et al, Phys. Rev. Lett. 110, 127204 (2013). [5] R. B. Helmholdt and K. H. J. Buschow, J. Less-Common Met. 128, 167 (1987). [6] S. Paul, A. Kundu, B. Sanyal, J. Appl. Phys. 116, 133903 (2014). [7] H. Luo, G. Liu, F. Meng, J. of Magn. Magn. Mater. 321, 4063–4066 (2009). [8] M. Nazmunnahar, T. Ryba, J. J. delVal, J. of Magn. Magn. Mater. 386, 98–101 (2015).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/caad83c069f8c7210cc235e36301d783.jpg)
X-ray diffraction patterns for Ni2-xMn1+xSn series of RQ alloys where (x = 0, 0.25, 0.5, 0.75, 1) and Ni1.25Mn1.75Sn, NiMn2Sn FC alloys.

![](https://assets.underline.io/uploads/markdown_image/1/image/147e5a0f09a0160a24d74cfbbc2a0169.jpg)
Schematic diagram of the ab plane of Mn2NiSn showing the segregation of Mn3Sn- type Hexagonal defect phase.

 Antisite disorder and defect phase segregation and its role in magnetic properties of Mn2NiSn

The search for new functional magnetic materials has led to the exploration into multi-component alloys (MCAs), which contain four or more elements with percentages between 12.5 % and 33 %. These include the sub-class of high entropy alloys (HEAs) which often exhibit single phases, thought to be stabilised by the high entropy of mixing of these compounds amongst other competing thermodynamic contributions. The resulting alloys often have interesting functional properties, which are linked to their structure. This work has explored a number of different compositions (CoFeNi0.5Cr0.5Alx, CoFeNi0.5CrxAl CoFeNiCryCux) in both bulk form, to understand how the phases present in the alloys influence the magnetic properties. The work has also used semi-empirical models to predict the different phases within the compositions and machine learning has been used to predict the saturation magnetisation. For example, in bulk CoFeNi0.5Cr0.5Al it was found that Fe-Cr nanoparticles formed in a weakly magnetic CoNiAl matrix [1]. This formation of round nanoparticles increased the magnetisation of the CoFeNi0.5Cr0.5Al alloy by a factor 4 at room temperature. For all the CoFeNi0.5Cr0.5Alx samples, the coercive field was small and the saturation magnetisation comparable with NiFe. While for the bulk CoFeNi0.5CrxAl samples, nanoparticles formed in all the samples containing Cr, with the CoFeNi0.5Cr1.0Al sample having square nanoparticles. Thus the amount of Cr in CoFeNi0.5Al dominates the geometry of the FeCr nanoparticles that form. Plus encapsulating the Cr in Fe-Cr nanoparticles meant that the Curie Temperatures for all the alloys containing Cr and Al were above 300 K (Fig. 1a). While for bulk CoFeNiCryCux (x = 0.0 - 1.5 and y = 0.0 - 1.0) samples, it was found that the addition of Cu could be used to tune the Curie Temperature over a 100 K range, with less than a 30 % reduction in the saturation magnetisation at 10 K, while the addition of Cr changed the Curie Temperature by 700 K and reduced the saturation magnetisation by 60 % at 10 K. By tuning the Cu to Cr ratio it is possible to achieve a Curie temperature around room temperature. The magnetocaloric effect (MCE) was studied in these materials, as the compositions tend to involve cheaper, more widely available elements, plus careful tuning of the composition will allow for the Curie Temperature to be around 300 K. It was found that the refrigeration capacity for the CoFeNiCryCux alloys was ~ 70 J/kg and for the CoFeNi0.5Cr0.5Alx alloys was ~ 75 J/kg at 2 T (Fig. 1b) [2]. Although these are smaller than existing materials such as LaFeSi, these HEAs still have advantages due to the presence of second order transitions, and having a wider operational temperature range for the MCE. In conclusion, the MCAs studied have promising functional magnetic properties, with small coercive fields, and reasonable saturation magnetisations. The phases present can be tuned via the composition, allowing for two-magnetic phase systems to be fabricated, that improve the magnetic properties compared to a single solid solution. The research carried out is only the tip of the iceberg, but these alloys are showing encouraging results.

**References:**

[1] N. A. Morley, C. R. B. Lim, J. Xi, A. Quintana-Nedelcos, and Z. Leong, Scientific Reports, 10, 14506, (2020) [2] A. Quintana-Nedelcos, Z. Leong, and N.A. Morley, Materials Today Energy, (accepted 2021)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/279780d9001aefe47b6797a2a7923640.jpg)
Figure 1a. Curie Temperature for CoFeNi0.5Cr0.5Alx, CoFeNi0.5CrxAl and CoFeNiCryCux alloys. 1b. Refrigeration capacity at 2 T for CoFeNi0.5Cr0.5Alx, CoFeNi0.5CrxAl and CoFeNiCryCux alloys. Shaded area useful temperature range.

 High Entropy Alloys: The next big thing in functional magnetic alloys

Fe-Co alloys have been ideal materials for various applications where high saturation magnetization (Ms), high permeability (µ) and low coercivity (Hc) are needed, especially for the applications where the minimization of mass and volume is a design parameter, such as aerospace industries [1-6]. In addition, their high Curie temperature makes them competitive candidates for high temperature applications [4]. In this study, in order to fully understand the compositional effect of Fe-Co alloys on their magnetic behaviour, Fe25Co75, Fe50Co50 and Fe75Co25 alloys were produced by using arc melting technique and characterized with the help of X-ray diffraction (XRD) analysis and superconducting quantum interference magnetometer (SQUID). XRD analysis indicated that all the samples were single phase (BCC), proving that oxidation did not occur during the manufacturing process. Furthermore, from their hysteresis loops, their magnetic saturations, remnant magnetizations and coercivities were found in the ranges 253-280 emu/gr, 2.06-6.15 emu/gr and 8.67-40.24 Oe, [4] respectively (Fig. 1). It was observed that as the % Co (wt.) content increases, saturation and remnant magnetization decreases, thus the Fe-rich alloys have higher magnetic properties. Also, after arc melting, stress-relief annealing was employed at 1073 K for 1 hr, resulting in reduction of the magnetic saturation (220-272 emu/gr), remnant magnetization (1.08-3.8 emu/gr) and coercivity (4.73-13.28 Oe) [5] due to the microstructural refinement, as it can be seen clearly in Fig.1. This study confirms that arc melting helps to improve the magnetic properties of Fe-Co alloys, because their magnetic properties is higher than the magnetic properties of those fabricated by mechanical alloying (Ms: 150-225 emu/gr, Hc: 25-70 Oe) [4] and cold rolling (Ms: 217.8 emu/gr, Hc: 87.7 Oe) [3]. Curie temperatures of the as-cast samples (1097-1181 K), are slightly lower than the values in the literature [5], also decreased slightly to the value range of 1071-1107 K after the heat treatment. By looking at Fig. 2, for the high temperature magnetic performances of as cast and annealed specimens, it can be seen that the annealed samples have a higher saturation magnetisation as a function of temperature. Further study has been undertaken to design and develop a new Fe-based amorphous alloy, with similar properties to Fe-Co alloys. The alloy of Fe45Co15B28Al10B2 was arc-melted followed by stress-relief annealing at 573 K for 1 hr to achieve the desired magnetic properties and amorphous morphology. Using DCS and temperature dependence magnetisation measurements up to 1000 K, the temperature dependence structural and magnetic behaviour was studied. DSC analysis indicates that the sample has glassy phase with the melting point of 1200 K and glass transition temperature of 800 K. Moreover, its Curie temperature was found as 940 K. Thus, this composition has promising results as a new soft magnetic material for high temperature applications.

**References:**

[1] S. Bergheul, A. Haddad, H. Tafat and M. Azzaz, “Magnetic , Microwave and Absorbing Properties Fe-Co of Alloy Synthesized by Mechnanical Alloying Process,” in Application of Contemporary Non-Destructive Testing in Engineering, Portoroz, Slovenia, 2005. [2] R. Fingers and G. Kozlowski, “Microstructure and magnetic properties of Fe–Co alloys,” Journal of Applied Physics, vol. 81, no. 8, pp. 4110-4111, 1997. [3] S. Honjo, S. Kikuchi, A. Sakuma, S. Ishio and J. Balachandran, “Preparation of magnetic Fe-Co alloy using texture controlled technique,” Journal of Magnetism and Magnetic Materials, vol. 465, pp. 64-69, 2018. [4] F. Sanchez-De Jesus, A. Bolarin-Miro, C. Cortes Escobedo, G. Torres-Villasenor and P. Vera-Serna, “Structural Analysis and Magnetic Properties of FeCo Alloys Obtained by Mechanical Alloying,” Journal of Metallurgy, vol. 2016, p. 8347063, 2016. [5] R. Sundar and S. Deevi, “Soft magnetic FeCo alloys: alloy development, processing, and properties,” International Materials Reviews, vol. 50, no. 3, pp. 157-192, 2005. [6] Y. Zhang, R. Ma, S. Feng, L. Cheng, P. Davies and P. Yu, “Microstructures and magnetic properties of Fe-35%Co alloy fabricated bymetal injection molding,” Journal of Magnetism and Magnetic Materials, vol. 497, p. 165982, 2020.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b10699c413116d6f4d080b35248523dd.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/a2da2cf697ccf1f003fbfd7cb478a821.jpg)

 Heat treatment investigations of Fe-based alloys

Permanent Magnet Synchronous Machines (PMSM) have gradually become the mainstream solution for drive motors. One main trend for advanced PMSM optimization is focused on new material application. Conventional PMSM iron core is fabricated by Non Grain-Oriented (NGO) silicon steel sheet lamination, based on traditional technic, the high-strength NGO silicon steel uses solid solution strengthening and precipitation strengthening methods to enhance mechanical strength [1]. As mentioned in [2-3], due to the increase in mechanical strength of high-strength NGO steel, a narrower magnetic bridge design can be applied in the rotor design. Thereby reducing the leakage flux to increase the output torque and the overall efficiency. Even though the high-strength NGO had higher iron loss, counting the reduction in the rotor bridge width, the overall performance can be significantly enhanced. The maximum speed of advanced PMSM is mostly greater than 12000rpm, thus the rotor sustains significant tensile stress due to centrifugal force. It’s easier to form the striped 180° domain with higher permeability under tensile stress [4]. While the frequency influence and thermal condition has not been discussed in [4]. In this paper, we designed an equivalent sample to simulate the condition of an iron core, similarly in [5], with different excitation and manufacture parameter such as laminated number and thermal condition. Three types of high-strength NGO (SW35YS600\SW35YS900\B35AHS500) and a controlled group (conventional NGO B35A270) has been tested. Magnetic performance of equivalent sample is measured by ring sample method. Two narrow sides of the sample are subjected to large tensile stress concentration, the lamination number is 5, 10, and 20. At high excitation frequencies, the level of iron loss improvement caused by tensile stress is significant. Four materials have also been annealed respectively at 720°C for 10 hours under argon atmosphere. After annealing, the tensile stress has lower effect on the magnetic properties for B35AHS500 and B35A270, but is significant for SW35YS600 and SW35YS900. Therefore, it is speculated that a large proportion of heat treatment process has been adopted during SW35YS600 and SW35YS900. In addition, we also conduct metallographic experimental observations on the material of the sample and the tensile fracture, and analyze the reasons for this result at the grain level. And in the subsequent experiment process, we further explored the influence of different environmental temperatures and different tensile stresses on the magnetic properties of these four NGO materials.

**References:**

[1] GONG, Jian, and Hai-wen LUO. "Progress on the research of high-strength non-oriented silicon steel sheets in traction motors of hybrid/electrical vehicles." Journal of Materials Engineering 43.6 (2015): 102-112. [2] Tietz, Marco, et al. "Effects and advantages of high-strength non grain oriented (NGO) electrical steel for traction drives." 2013 3rd International Electric Drives Production Conference (EDPC). IEEE, 2013. [3] Tanaka, Ichiro, et al. "Application of high-strength nonoriented electrical steel to interior permanent magnet synchronous motor." IEEE transactions on magnetics 49.6 (2012): 2997-3001. [4] Senda, Kunihiro, et al. "Magnetic properties and domain structure of non-oriented electrical steel under stress." IJTFM 131.10 (2011): 884-890. [5] Gao, Lingyu, et al. "Application of grain-oriented electrical steel used in super-high speed electric machines." AIP Advances 10.1 (2020): 015127.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/ccde1a3a9953b57d9bd60bd246980599.jpg)
metallographic experimental observations

![](https://assets.underline.io/uploads/markdown_image/1/image/793eec53db2c5ec9c61420c606acc711.jpg)
sample of the tensile experiment and result of the magnetic test

 Magnetic Performance Improvement caused by Tensile Stress in Equivalent Iron Core fabricated by High-Strength Non-Oriented Electrical Steel

Grain-oriented (GO) electrical steel sheets, as a core material used in power transformers broadly, often holed at particular and suitable regions to make the stacking of the sheet easier. The magnetic properties around the hole will be deteriorated, and the local magnetic domain structure will be impacted [1]. These phenomena will significantly influence the performance of the transformer [2]. It is necessary to accurately measure and study the influence of hole on local vector magnetic properties to optimize the core structure. Some studies use the search coils wound by drilling through the sample to analyze the flux orientations [3]-[4]. Yet the distributions of local magnetic properties around the hole have not been clarified. To obtain a more comprehensive and accurate distribution of local vector magnetic properties, and avoid damaging the sample to wind the search coils, the needle probes and double H-coils, which were discussed in [5] and [6], are adopted and developed. In this paper, a novel local magnetic property testing system using an accurately-movable B-H sensor is designed and constructed to study the local vector magnetic properties around the hole in a transformer type core made of grain-oriented electric steel sheets. The influence of sizes and positions of the holes on the properties is compared and analyzed at 50-400 Hz. In addition, the changes of the magnetic domain are observed and analyzed. Fig. 1 shows the local testing system with an accurately-movable B-H sensor and sample with 300 mm × 300 mm studied in this paper. The distance between the needles of B sensor is 3.5 mm, and H-sensor coil is 2 mm × 2 mm assembled using wire with 0.01 mm diameter. The number of turns is 600. The sensor is mounted to a XYZ-precise-positioning device, hence, the magnetic flux density, B, magnetic field strength, H, and core loss, P, of each point in the local area can be measured. Eight holes with diameter of D and interval of 100 mm are punched in the transformer sheets. W is the distance of the center of hole to the inner edge. Different D and W are studied in the paper. Single layer grain-oriented steel sheet (namely, B30P120) is tested. To study the vector magnetic properties around the hole, the square area (47 × 47 mm2) centered at the hole is measured and studied. The H = H (x, y), B = B (x, y) and P = P (x, y) are measured when the sensor is moved 1 mm in interval along the x-, or y-directions and the total number of measured points are 2000. As a criterion of measurements, the peak flux density is measured and controlled with a search coil, which wound uniformly around the sample. Fig. 2 shows P distributions around the hole with D = 4 mm and W = 25 mm when B is controlled to be 50 Hz sinusoidal waveform with 1.5 T amplitude. The domain structures at particular areas A1 and A2 are also shown in Fig. 2. The results show that the core loss P around the hole is much larger than that around points far to the hole. It could be attributed to the non-uniform flux density distributions, in particular, increasing flux density close to the hole. It is consistent with the observation of domain structure around the hole. The domain structure is disordered close to the hole due to strain introduced by hole punching while it is uniform at the point far to the hole. In addition, the distribution of the flux density B, field strength H, and core loss P around the hole with different diameters (D = 4, 6, 8, 10 mm) and different positions (W = 20, 25, 30 mm) are measured and analyzed in the flux density ranges of 0.1 – 1.5 T at 50 – 400 Hz. The experiment results show that the local B and H distributions are altered by the hole. The domain pattern is also restructured due to strain introduced by hole punching. As a result, the local losses are increased. In addition, the losses around the hole increase to 150 – 300 % compared with the losses at the far point depending on D value of 4 – 10 mm. However, the location of hole has minor influence on the losses, in particular, when the hole is located at the symmetrical position (W = 25 mm). In conclusion, small size of the hole and symmetrical position of the hole during stacking of the core can minimize the deterioration of local magnetic properties and improve the performance of the stacked electrical steel sheets ultimately.

**References:**

[1] T.Gunes, R. Schäfer and N.Derebasi, IEEE Transactions on Magnetics., Vol. 54, p.1-8 (2018) [2] G. teNyenhuis, S. Girgis and F. Mechler, IEEE Transactions on Power Delivery., Vol. 16, p.648-653 (2001) [3] T.Gunes, N.Derebasi and C.Erdonmez, IEEE Transactions on Magnetics., Vol. 51, p.1-4 (2015) [4] N Derebasi. C Erdonmez, Journal of Superconductivity and Novel Magnetism., Vol. 30, p.3309-3313 (2017) [5] S. Imamori, S. Aihara and H. Shimoji, Journal of Magnetism and Magnetic Materials., Vol. 500 (2020) [6] S Aihara, H Shimoji and T TodakaM, IEEE Transactions on Magnetics., Vol. 48, p.4499-4502 (2012)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/dd07283974bbc51d0a8991dc1207b2ad.jpg)
Fig. 1.Novel local measurement system and measurement sample.

![](https://assets.underline.io/uploads/markdown_image/1/image/cf2651f0e011db64b56bc65193d42edc.jpg)
Fig. 2. Distributions of the core loss around the hole and magnetic domain structure at area A1, A2.

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