France



Crystallization from an amorphous precursor has been historically exploited to synthesize metastable soft magnetic nanocomposite systems comprised of transition metal rich ferromagnetic nanocrystals embedded within an intergranular amorphous phase [1,2]. Subsequent thermal processing steps are designed to produce a highly refined microstructure with grain sizes in the nm-range to enable effective exchange averaging and highly reversible magnetization processes as described by the random anisotropy model [3,4]. Successful realization of superior soft magnetic properties results from ubiquitous nucleation and self-limiting growth of nanocrystals due to expulsion of glass formers to the intergranular amorphous matrix, thereby requiring significant amounts of metalloids and early transition metals which reduce the effective saturation magnetization.&lt;br/&gt;Recent work has discovered the balance between nucleation and growth can alternatively be optimized through highly controlled thermal processing steps involving extremely rapid heating and subsequent cooling rates on the order of 10&lt;sup&gt;2&lt;/sup&gt;-10&lt;sup&gt;3&lt;/sup&gt;K/s referred to as “flash annealing”, such that a high density of nucleation sites can be achieved with limited nanocrystal growth even for alloy chemistries which do not produce nm-scale microstructures for conventional annealing [5]. High heating and cooling rates generated during laser radiation have offered great advantages for rapid thermal processing and surface treatment of ferromagnetic amorphous ribbons [6–8].&lt;br/&gt;In this work, we report on pulsed-laser processing of a series of Fe-based and Co-based amorphous ribbons with emphasis being placed on phase transformations within the laser-irradiated regions as well as impacts on surrounding regions due to the detailed thermal effects associated with the laser pulse.&lt;br/&gt;Amorphous ribbons with nominal compositions Fe&lt;sub&gt;73.5&lt;/sub&gt;Si&lt;sub&gt;13.5&lt;/sub&gt;Nb&lt;sub&gt;3&lt;/sub&gt;B&lt;sub&gt;9&lt;/sub&gt;Cu&lt;sub&gt;1&lt;/sub&gt; and Co&lt;sub&gt;77.2&lt;/sub&gt;Fe&lt;sub&gt;1.4&lt;/sub&gt;Mn&lt;sub&gt;1.4&lt;/sub&gt;Nb&lt;sub&gt;4&lt;/sub&gt;B&lt;sub&gt;14&lt;/sub&gt;Si&lt;sub&gt;2&lt;/sub&gt; were used for laser experiments. Laser annealing has been carried out using a Q-smart pulsed Nd-YAG (wavelength of 1064 nm) and laser beam diameter around 0.6 mm. The laser power was varied from 50-150 mJ and the time delay associated with each laser power has been calculated using the FWHM of the laser output pulse. Microstructure and crystalline phases after laser annealing were characterized using both XRD and SEM. Temperature profiles achieved during laser annealing have been estimated using FEA heat transient analysis.&lt;br/&gt;Fig. 1 shows XRD patterns of as-cast and laser annealed Fe- and Co-based ribbons at different laser power. As can be seen, as-cast and samples annealed at 60 mJ are fully amorphous, whereas samples treated at higher laser power (i.e., at 100mJ and above) present a mixture of boride and α-Fe phases with an average grain size of the latter being on the order of 8-10 nm estimated from the FWHM of the corresponding crystalline peak. In conventional annealing methods, the formation of boride phases are typically observed only after reaching secondary crystallization temperatures in these alloy systems. Because the borides are already observed in laser annealed samples with significant volume fraction for the lowest pulse power levels at which crystalline peaks can be detected, a need exists to revisit the crystallization mechanism in light of the detailed thermal histories experienced across and surrounding the irradiated region and the Gaussian distribution of the laser beam across the irradiated surface. Laser energy density decreases exponentially in the radial direction and has a maximum at the center of the laser spot leading to greater temperature values at the center and a temperature gradient from the center towards the edges. In addition, laser beam intensity also produces a stress distribution which can be associated with its Gaussian profile resulting in an inhomogeneous stress across the beam diameter, tensile at the center and compressive at the edges. Such temporally and spatially dependent temperature and stress profiles can alter free volume, nucleation rate, and diffusion kinetics leading to rapid crystallization events and alternative pathways towards the crystalline phase.&lt;br/&gt;Temperature as a function of time of a simulated laser-irradiated spot at the top surface and below the top surface of an amorphous ribbon is shown in Fig. 2. The temporal temperature profile suggests temperatures above 2000K at the top surface which is beyond the melting point and most likely results in a local melting at the top surface layer with crystallization possible upon cooling. Temperatures estimated approximately 2 microns below the surface (within the skin penetration depth of Fe at the optical wavelength of Nd-YAG) are on the order of 1000K and likely not to reach melting temperatures, consistent with the surface deformation observed as illustrated in the SEM images of Fig. 2c. These temperatures significantly decrease within the distance from the center towards the edge of the laser spots suggesting diffusion of heat from the centers outwards and therefore different crystallization kinetics are expected.

**References:**

[1] M. E. McHenry, M. A. Willard, and D. E. Laughlin, &lt;i&gt;Progress in Materials Science&lt;/i&gt;, vol. 44, no. 4, pp. 291–433, Oct. 1999&lt;br/&gt;[2] A. Talaat, D. W. Greve, M. V. Suraj, and P. R. Ohodnicki, &lt;i&gt;Journal of Alloys and Compounds&lt;/i&gt;, vol. 854, p. 156480, Feb. 2021&lt;br/&gt;[3] G. Herzer, &lt;i&gt;Acta Materialia&lt;/i&gt;, vol. 61, no. 3, pp. 718–734, Feb. 2013, doi: 10.1016/j.actamat.2012.10.040.&lt;br/&gt;[4] Y. Yoshizawa, S. Oguma, and K. Yamauchi, &lt;i&gt;Journal of Applied Physics&lt;/i&gt;, vol. 64, no. 10, pp. 6044–6046, Nov. 1988&lt;br/&gt;[5] K. Suzuki &lt;i&gt;et al &lt;/i&gt;&lt;i&gt;AIP Advances&lt;/i&gt;, vol. 9, no. 3, p. 035311, Mar. 2019, doi: 10.1063/1.5079778.&lt;br/&gt;[6] S. Katakam &lt;i&gt;et al.,&lt;/i&gt; &lt;i&gt;Scripta Materialia&lt;/i&gt;, vol. 66, no. 8, pp. 538–541, Apr. 2012&lt;br/&gt;[7] C. Smith &lt;i&gt;et al.&lt;/i&gt;, &lt;i&gt;Materials Letters&lt;/i&gt;, vol. 122, pp. 155–158, May 2014&lt;br/&gt;[8] L. Lanotte and V. Iannotti, &lt;i&gt;Journal of Applied Physics&lt;/i&gt;, vol. 78, no. 5, pp. 3531–3533, Sep. 1995

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b76b25d19ac56909b73662b71acd2ab3.jpg)
Fig. 1 XRD patterns of as-cast and laser-annealed Fe-based (a), and Co-based ribbons (b)

![](https://assets.underline.io/uploads/markdown_image/1/image/428c178d2a796bb0b353dfa63449a262.jpg)
Fig. 2. FEA simulations of temperature as a function of time at a laser power of 100mJ for (a) top surface and (b) below top surface of an amorphous ribbon. Red (0.6 mm), green (0.4 mm), and blue (0.2 mm) lines represent the distance from the center towards the edge of the laser spot. (c) SEM image of the ribbon surface layer showing evidence for local surface melting at the center of the laser irradiated region.

INTERMAG 2021

 Laser Processing of Soft Magnetic Metal Amorphous Nanocomposites

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Crystallization from an amorphous precursor has been historically exploited to synthesize metastable soft magnetic nanocomposite systems comprised of transition metal rich ferromagnetic nanocrystals embedded within an intergranular amorphous phase [1,2]. Subsequent thermal processing steps are designed to produce a highly refined microstructure with grain sizes in the nm-range to enable effective exchange averaging and highly reversible magnetization processes as described by the random anisotropy model [3,4]. Successful realization of superior soft magnetic properties results from ubiquitous nucleation and self-limiting growth of nanocrystals due to expulsion of glass formers to the intergranular amorphous matrix, thereby requiring significant amounts of metalloids and early transition metals which reduce the effective saturation magnetization. Recent work has discovered the balance between nucleation and growth can alternatively be optimized through highly controlled thermal processing steps involving extremely rapid heating and subsequent cooling rates on the order of 102-103K/s referred to as “flash annealing”, such that a high density of nucleation sites can be achieved with limited nanocrystal growth even for alloy chemistries which do not produce nm-scale microstructures for conventional annealing [5]. High heating and cooling rates generated during laser radiation have offered great advantages for rapid thermal processing and surface treatment of ferromagnetic amorphous ribbons [6–8]. In this work, we report on pulsed-laser processing of a series of Fe-based and Co-based amorphous ribbons with emphasis being placed on phase transformations within the laser-irradiated regions as well as impacts on surrounding regions due to the detailed thermal effects associated with the laser pulse. Amorphous ribbons with nominal compositions Fe73.5Si13.5Nb3B9Cu1 and Co77.2Fe1.4Mn1.4Nb4B14Si2 were used for laser experiments. Laser annealing has been carried out using a Q-smart pulsed Nd-YAG (wavelength of 1064 nm) and laser beam diameter around 0.6 mm. The laser power was varied from 50-150 mJ and the time delay associated with each laser power has been calculated using the FWHM of the laser output pulse. Microstructure and crystalline phases after laser annealing were characterized using both XRD and SEM. Temperature profiles achieved during laser annealing have been estimated using FEA heat transient analysis. Fig. 1 shows XRD patterns of as-cast and laser annealed Fe- and Co-based ribbons at different laser power. As can be seen, as-cast and samples annealed at 60 mJ are fully amorphous, whereas samples treated at higher laser power (i.e., at 100mJ and above) present a mixture of boride and α-Fe phases with an average grain size of the latter being on the order of 8-10 nm estimated from the FWHM of the corresponding crystalline peak. In conventional annealing methods, the formation of boride phases are typically observed only after reaching secondary crystallization temperatures in these alloy systems. Because the borides are already observed in laser annealed samples with significant volume fraction for the lowest pulse power levels at which crystalline peaks can be detected, a need exists to revisit the crystallization mechanism in light of the detailed thermal histories experienced across and surrounding the irradiated region and the Gaussian distribution of the laser beam across the irradiated surface. Laser energy density decreases exponentially in the radial direction and has a maximum at the center of the laser spot leading to greater temperature values at the center and a temperature gradient from the center towards the edges. In addition, laser beam intensity also produces a stress distribution which can be associated with its Gaussian profile resulting in an inhomogeneous stress across the beam diameter, tensile at the center and compressive at the edges. Such temporally and spatially dependent temperature and stress profiles can alter free volume, nucleation rate, and diffusion kinetics leading to rapid crystallization events and alternative pathways towards the crystalline phase. Temperature as a function of time of a simulated laser-irradiated spot at the top surface and below the top surface of an amorphous ribbon is shown in Fig. 2. The temporal temperature profile suggests temperatures above 2000K at the top surface which is beyond the melting point and most likely results in a local melting at the top surface layer with crystallization possible upon cooling. Temperatures estimated approximately 2 microns below the surface (within the skin penetration depth of Fe at the optical wavelength of Nd-YAG) are on the order of 1000K and likely not to reach melting temperatures, consistent with the surface deformation observed as illustrated in the SEM images of Fig. 2c. These temperatures significantly decrease within the distance from the center towards the edge of the laser spots suggesting diffusion of heat from the centers outwards and therefore different crystallization kinetics are expected.

**References:**

[1] M. E. McHenry, M. A. Willard, and D. E. Laughlin, Progress in Materials Science, vol. 44, no. 4, pp. 291–433, Oct. 1999 [2] A. Talaat, D. W. Greve, M. V. Suraj, and P. R. Ohodnicki, Journal of Alloys and Compounds, vol. 854, p. 156480, Feb. 2021 [3] G. Herzer, Acta Materialia, vol. 61, no. 3, pp. 718–734, Feb. 2013, doi: 10.1016/j.actamat.2012.10.040. [4] Y. Yoshizawa, S. Oguma, and K. Yamauchi, Journal of Applied Physics, vol. 64, no. 10, pp. 6044–6046, Nov. 1988 [5] K. Suzuki et al AIP Advances, vol. 9, no. 3, p. 035311, Mar. 2019, doi: 10.1063/1.5079778. [6] S. Katakam et al., Scripta Materialia, vol. 66, no. 8, pp. 538–541, Apr. 2012 [7] C. Smith et al., Materials Letters, vol. 122, pp. 155–158, May 2014 [8] L. Lanotte and V. Iannotti, Journal of Applied Physics, vol. 78, no. 5, pp. 3531–3533, Sep. 1995

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b76b25d19ac56909b73662b71acd2ab3.jpg)
Fig. 1 XRD patterns of as-cast and laser-annealed Fe-based (a), and Co-based ribbons (b)

![](https://assets.underline.io/uploads/markdown_image/1/image/428c178d2a796bb0b353dfa63449a262.jpg)
Fig. 2. FEA simulations of temperature as a function of time at a laser power of 100mJ for (a) top surface and (b) below top surface of an amorphous ribbon. Red (0.6 mm), green (0.4 mm), and blue (0.2 mm) lines represent the distance from the center towards the edge of the laser spot. (c) SEM image of the ribbon surface layer showing evidence for local surface melting at the center of the laser irradiated region.

technical paper

Magnetic flux leakage (MFL) testing is a widely used method for detection of defects such as the metal loss due to the corrosion in steel bars located in reinforced concrete. The magnetic flux leak caused by the defect is then measured by a magnetic sensor, which gives information to the user of the presence of the defect [1]. For small defects, it is difficult for the sensor to determine the presence and location of the defects. Therefore, we propose a gradiometer measurement setup [2] using highly sensitive magnetic tunnel junction (MTJ) sensors, which is based on the tunnel magnetoresistance (TMR) effect, to increase the magnitude of the measured signal. In this work, 2 MTJ sensor modules were used to measure a gap of 0.1 mm between two steel bars (diameter = 20 mm, length = 500 mm). The length of the overall specimen is 1000 mm, where the gap is at the 500 mm mark (touching point of the two steel bars). The MTJ sensor modules consist of each sensor connected in a full-bridge circuit [3]. The signal is then passed through an amplifier circuit and low pass filter to further reduce noise in the measured signal. A robot is then used to move the sensor along the steel bars at a speed of 50 mm/s, measuring the magnetic flux signal. The lift-off value for the overall measurement is set at 50 mm. Before each measurement, the steel bars are magnetized using a solenoid (approximately 1.52 mT). A planar gradiometer setup was used, whereby the MTJ modules are placed in the same x-axis. The effect of the baseline (b), the distance between the first and second MTJ sensor module, was investigated (b = 50, 100, 150, 200 mm). From the results, the overall magnitude of the gradiometer sensor signal showed an increase, almost twice as large as the single sensor module signal (Figure 1). Furthermore, the gradiometer sensor signal produced a single peak response (gap is right above the peak), in contrast to the single sensor signal which produced a sine-like wave response (gap is between the peak and trough) (Figure 2). This is helpful in terms that the location of the gap can be determined in real time. For the estimated gap location, a loss in accuracy was observed in all the signals, with the most accurate location determined for b = 100 mm. We demonstrated that a planar gradiometer setup consisting of 2 MTJ sensors could improve the overall signal value of the measurement setup resulting in an easier way to determine the location of a 0.1 mm gap at 50 mm lift-off. The results obtained from this research are expected to contribute to the further application of MTJ sensors not only in MFL testing but in other fields of magnetic-based non-destructive testing. This work was partly supported by the Center of Innovative Integrated Electronic Systems (CIES) and the Center for Spintronics Research Network (CSRN). The sensors were provided by Mr. Kosuke Fujiwara from Spin Sensing Factory (SSF).

**References:**

[1] Hirose. M, Maeda. T, Matsuda. K, Yokota. M, Hattori. A, and Miyagawa. T, J. Stru. Eng., 58A, p.867-878 (2012). [2] W. Dehui, S. Lingxin, W. Xiaohong, and L. Zhitian, J Nondestruct Eval., 36, p.24 (2017). [3] J. Cao, and P.P Freitas, J. Appl. Phys., 107, 09E712 (2010).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/3227474d3e22de78e908e4d2ba9a7a6b.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/74835cefbcae5c802089d4aa85ecad44.jpg)

 Detection of Gap in Steel Bars by Magnetic Tunnel Junction Sensor Gradiometer based Magnetic Flux Leakage Testing Method

Non-destructive testing(NDT) is a method that detects discontinuous boundaries or defects that exist inside and outside of pipes without damaging or destroying pipe[1]. The magnetic flux leakage(MFL) method, one of the most representative NDT methods, uses permanent magnets to saturate the ferromagnetic pipe by applying a large magnetic field to the pipe and measure the leakage magnetic flux signal around the defect. Because the MFL signal contains information about the size and shape of the defect, an algorithm can be created to estimate the shape of the defect through analysis[2]-[3]. In general, many studies on non-destructive testing, including the MFL method, mainly focus on detecting defects or cracks in pipeline. However, not only defects but also metal objects outside the pipe affect the condition of the pipe. When external metal contact with the pipe wall, corrosion is more prone to the metal contact area on the pipe. In addition, pipes used in industry are coated with thick plastic to block contact with oxygen or external metal, However, the external metal damages the plastic coating. If a hole is made in the plastic coating of the pipe due to external metal, it is more prone to corrosion by contacting the pipe with oxygen as well as the external metal. Therefore, it is necessary to study not only defects but also detection of metal existing outside the pipe. Analysis of external metals requires research not only to prevent corrosion but also to prevent theft of energy sources in pipes such as oil and gas. People who steal energy sources eventually drill holes in pipes through which the energy source flows, and use metal pipes to steal the energy source. Metal pipe for theft can eventually be found through external metal analysis, so analysis of external metal also helps prevent theft of energy sources. Because the MFL method uses a large magnetic field, it is easier to detect and analyze external metals than other NDT methods. Fig.1 is here This paper analyzes the MFL sensor signals according to the shape and distance of external metal around underground pipeline, and presents an algorithm that estimates the shape and the distance of the external metal to the pipe using the MFL sensor signal. Fig. 2 shows the magnetic flux and the magnetic flux density distribution when there is a defect or external metal, respectively. Fig.2 is here When a defect occurs in the pipe, the cross-sectional area through which the magnetic flux passes decreases, and the magnetic flux flowing through the pipe leaks out of the pipe. In other words, when a defect occurs in the pipe, the magnitude of the signal detected by MFL sensors increases. Conversely, when external metal contacts the pipe, the cross-sectional area through which the magnetic flux passes increases, and the amount of magnetic flux leaking out of the pipe decreases. That is, when external metal contacts the pipe, the magnitude of the signal detected by MFL sensors decreases. In addition, like defects, when the shape or distance of the external metal changes, the size or shape of the signal detected by the MFL sensors also changes. As a result, the shape and distance of the external metal can be estimated by applying the signal change tendency and the defect shape estimation algorithm[2]-[3].

**References:**

[1] M. Rostami Kandroodi, B. Nadjar Araabi, and M. Nili Ahmadabadi, in IEEE Transactions on Magnetics, Vol. 53, No. 3, pp.1-10(2017) [2] H. M. Kim, C. G. Heo, and G. S. Park, in IEEE Transactions on Magnetics, Vol. 54, No. 11, pp.1-5(2018) [3] H. M. Kim, and G. S. Park, in IEEE Transactions on Magnetics, Vol. 50, No. 2, pp.109-112(2014)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/2b20fbd5d4eb98b21e434495c17d02dc.jpg)
Fig.1. Photographs of external metal and pipes

![](https://assets.underline.io/uploads/markdown_image/1/image/3c45be0487083f6756301fe5aa5bacd7.jpg)
Fig.2. Distribution of magnetic flux density around (a) the defect. (b) the external metal. Diagram of magnetic flux line around (c) the defect. (d) the external metal.

 A Study on the Estimation of the Shapes and Distance of External Metal around Underground Pipeline Using Magnetic Flux Leakage Sensors

Temperature measurements are vital in laboratory and industry settings. For that reason, many different measurement methods exist. Common temperature sensors are based on thermistors, resistance temperature detectors, thermocouples, semiconductors and fiber optics. Optical sensor technologies possess advantages for electrical measurements and analysis in high voltage applications like monitoring in smart grids, ranging from several Hz to some kHz in frequency. In general, in such systems current measurements are based on sensing the magnetic field generated by the electric current rather than the current itself. By this, possible electric hazards are eliminated, and a natural galvanic separation is built into the sensing system. Optical magnetic field or direct current measurement systems have been realized through fiber-optic current sensors, based on an optical fiber around a current conductor1, bulk magnetic garnet materials2, or single or stacked garnet thin films3. We demonstrate magnetooptical measurement schemes based on the Faraday effect for relating magnetic and micromagnetic features to magnetic field, electrical current and temperature4. For that, we use magnetooptical active only micrometer thin iron garnet film. The garnet film has a large Verdet constant (Faraday effect) together with a low absorption coefficient. The magnetooptical Faraday effect sensitivity of the magnetooptical indicator film is 1.8°/mT at room temperature. The polarimetric magnetooptical read-out is realized in a non-contact optical arrangement. For the measurements, the magnetic state of the garnet sensor is modulated through a low noise current amplifier and the magnetooptical response signal is acquired through a lock-in amplifier. The thin film sensor element has the ability for high sensitivity magnetic field and naturally electrical current sensing capabilities over a many magnitudes of field and current. High-resolution magnetic field sensitivities down to the nanotesla regime are demonstrated as shown in Fig. 1. For the micrometer thick garnet films we demonstrate a magnetic field detection limit of better than 20 nT/Hz0.5 and an electrical current detection limit of 3 mA/Hz0.5. These values are achieved over several orders of magnitude of magnetic field frequencies into the kHz regime. The magnetooptical response at small magnetic field amplitudes originates from small magnetic angle rotations inside the magnetic domain walls. Different schemes for the measurements of temperature are evaluated. In all cased the range of application is defined by the Néel temperature of the sensing garnet material5. Exemplary results of temperature sensing based on variations in the magnetic domain wall susceptibility with temperature, from analyzing the resulting magnetooptical signal amplitude in a magnetic modulation field, are shown in Fig. 2. Temperatures are derived with high accuracy of 0.1°C across a wide temperature range. Overall, we propose and demonstrate a multimode magnetooptical temperature sensor relying on combined magnetooptical Faraday effect amplitude variations with the related physical properties. The sensing range spans many orders of signal amplitude and a wide frequency range with equal sensitivity and a limit of detection. Temperatures are derived with high accuracy across a wide temperature range. The measurement schemes can be extended to the use of other magnetic sensing materials. These include ferromagnetic and ferromagnetic films with lower or higher Néel temperatures or Curie temperatures. The magnetooptical sensor allows for measuring magnetic field, direct electrical current, and temperature without electrical connection and will therefore be especially useful in applications requiring resistance to electrical noise. This work was funded by the German Research Foundation (DFG) through grant no. MC 9/20-1.

**References:**

 1. R. Silva, H. Martins, I.Nascimento et al., Applied Sciences, Vol. 2, p. 602-628(2012) 2. B.J.H. Stadler, and T. Mizumoto, IEEE Photonics Journal, Vol. 6, p.1-15(2014) 3. A. Garzarella, M.A. Shinn, and D.H. Wu, Appl. Phys. Lett., Vol. 108, p.241101 (2016) 4. F. Klingbeil,S. D. Stölting and J. McCord ,Appl. Phys. Lett.,(under review) 5. M. Kustov, R. Grechishkin, M. Gusev et al, Adv. Mater., Vol. 27, p.5017 (2015)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/ffc2beb949bb793dba35dd1aff13a879.jpg)
Fig. 1 Exemplary limit of detection plot for a base-to-peak magnetic field Hac amplitude with a frequency of fH = 1110 Hz.

![](https://assets.underline.io/uploads/markdown_image/1/image/ae6308b2c378f5878afb9133bee9a314.jpg)
Fig. 2 Magnetooptically measured temperature derivation with varying temperature. The inset shows the corresponding change of the magnetooptical signal amplitude versus a thermocouple temperature used for calibration.

 Multimodal Magnetooptic Sensing of Magnetic Field, Electric Current and Temperature

Due to their special features such as low background noise, high sensitivity, and the possibility of nanoscale fabrication, magnetoresistive (MR) sensors present a great interest for the scientific community, being used especially in biomedical applications [1]. Among MR sensors, those based on tunnel magnetoresistance (TMR) effect exhibit the highest magnetoresistance value, higher sensitivity, and lower power consumption [2, 3]. However, shortcomings as weak linearity in the variation of electric resistance with the applied magnetic field, high coercive field, or quite complicated fabrication process, make the TMR sensors less investigated for biosensing applications. The purpose of this work is to develop a model of TMR based sensor, with small size and high sensitivity, capable of real-time detection of magnetic particles, with possible applications in the detection of magnetic particles from tissues in hyperthermia. The operation of the sensor will be done without magnetic flux concentrators or permanent magnets, therefore no additional steps in the technological process are required. In order to obtain a TMR structure with high values of magnetoresistance and sensitivity, the thicknesses of the component layers were optimized reaching the final configuration: Ta (5 nm)/Ru (20nm)/Ta (5 nm)/CoFe (3 nm)/IrMn (20 nm)/CoFe (2.5 nm)/Ru (0.85 nm)/CoFeB (3 nm)/MgO (1.8 nm)/CoFeB (3 nm)/Ta (10 nm). The thin layers were deposited on Si/SiO2 support (18 x 18 mm2), using the ATC 2200 / AJA International deposition system, which allows both magnetron sputter deposition and electron beam evaporation. Excepting the MgO barrier layer, which was deposited by electron beam evaporation, all other thin layers were obtained by magnetron sputter deposition at a working pressure of 3 mTorr. To induce an in-plane easy axis in the ferromagnetic layers, a magnetic field (100 Oe) parallel to the plane of the thin films, was applied during the deposition. As a detection element, we proposed a structure with a low aspect ratio in the shape of a cylindrical pillar with 4 µm diameter. The total size of each device (sensor and contacts) is 400x250 μm2, as depicted in Fig. 1(a). These small dimensions allow the fabrication of 100 sensors onto single support of 18x18 mm2, leading to a simplification of the technological process. In the first step of the sensor microfabrication, rectangular structures of 40x80 μm2 (Fig. 1(b)) were patterned by e-beam lithography, then ion beam milling and removing of e-resist. The detection element (pillar) and contact surfaces, were defined by electron beam lithography and ion beam milling. In order to passivate the lateral sides of the pillar, 60 nm of insulating layer (SiO2) was sputter-deposited, then the resist and SiO2 from the top of the pillar and contact surfaces were removed by lift-off technique. For measurements, the CPP (“current perpendicular to plane”) configuration was used by passing the electric current perpendicular to the plane of the thin layers. For this step, the electrical contacts with coplanar waveguide geometry (ground-signal-ground), were defined by laser beam lithography followed by Ta (10 nm)/Cu (200 nm) deposition and then lift-off. The coplanar waveguide geometry consists of a middle pad which is the signal strip and two lateral pads that are used as ground lines. The middle pad is connected to the top of the pillar and the two outer pads are connected to the bottom part of the pillar. After microfabrication, TMR sensors were annealed at a temperature of 320 °C in an external magnetic field of 5 kOe, in the direction of anisotropy induced during the deposition. For the as-fabricated TMR sensor, we obtained a magnetoresistance value of 37.5 % and 0.26 %/Oe for the sensitivity. In order to test the TMR sensor, detection experiments were performed using FeCrNbB magnetic particles with a diameter of 50-500 nm obtained by high-energy ball milling from melt-spun ribbons (MSRs) precursors [4]. The magnetic nanoparticles were dispersed in cyanoacrylate adhesives in order to obtain solutions with concentrations of 20 % and 80 %. Then, a glass support was partially covered with these solutions and used to demonstrate the detection capability of the MR sensor. During the measurements, a magnetic field of 20 Oe was applied perpendicular to the direction of the sensor, in order to generate a dipole magnetic field of magnetic particles. The output voltage variation was measured for different concentrations of magnetic particles. As shown in Fig. 2, an output voltage variation of 0.044 mV was obtained for a magnetic particles concentration of 20 %. Increasing the particles concentration to 80%, the variation of the sensor output voltage also increased to 0.161 mV. We demonstrated a dependent relationship between the sensor output voltage and the concentration of magnetic particles. It should also be noted that the values of electrical resistances remain constant over time, indicating the stability of the sensor. Acknowledgment This work was supported by the Romanian Ministry of Research and Innovation (MCI) under the NUCLEU programme, Project PN 19 28 01 01.

**References:**

[1] D. Su, K. Wu, R. Saha et al., Micromachines, Vol. 11, p. 1 (2020) [2] D. D. Djayaprawira, K. Tsunekawa, M. Nagai et al., Applied Physics Letters, Vol. 86, p. 092502-1 (2005) [3] S. Cardoso Freitas, D. C. Leitao, L. Gameiro et al., Microsystem Technologies, Vol. 20, p. 793 (2014) [4] H. Chiriac, N. Lupu, M. Lostun et al., Journal of Applied Physics, Vol. 115, p.17B520-1 (2014)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/8bc6653b5ea2fe219e2b8925ce30340c.jpg)
Fig. 1. (a) Optical microscope image of a device with G-S-G type electrical contacts and (b) SEM image of a rectangular structure.

![](https://assets.underline.io/uploads/markdown_image/1/image/1b121658e4c6ea5578969a74a9676632.jpg)
Fig. 2. Real-time signal of the TMR sensor for different magnetic particles concentrations.

 Tunnel Magnetoresistance Sensors for Real-Time Detection of Magnetic Nanoparticles

The aerospace industry has become increasingly reliant on carbon fibre composite (CFC) as a structural component in aircraft, due to his strength to weight ratio. One of the disadvantages of CFC are their vulnerability to impact damage and external environment changes during flight, which causes barely visible damage to the structure. One solution to this is structural health monitoring (SHM), which monitors the composite structure to provide early detection for damage. Previous work has found that magnetostriction ribbons used as sensors have good resolution to detect damage but are time consuming to attach to the composite surface [1]. To solve this issue, research has focused on additive layer manufacturing to produce 3D printed sensors, which can be directly attached to the surface of the composite. To evaluate the magnetic performance of the 3D printed sensors, this project focused on the design and development of 3D printed magnetostrictive materials using computer simulations in parallel with experimental research to measure the magnetic and mechanical performance. The design of the magnetostrictive sensors was simulated using COMSOL multiphysics, to determine how design parameters such as thickness, track length and gap changed the magnetic response with and without applied force. A desktop bound metal deposition printer at the Royce Translational Centre, was then used to print the metal with a polymer binder (polypropylene) to create the chosen 3D structures (Fig 1). The magnetic material chosen for this work was stainless steel 17/4ph. The initial work assessed the current performance of the printer at each stage. The designed structure goes through two stages of printing and heat treatment processes. The development of the feedstock to the heat treated structure could affect the final performance of the magnetostrictive material. For example, after printing the printed part (green body) goes through a debinder and a sintering stage which ultimately changes the microstructure and allows grain growth, which could affect the saturation magnetisation magnitude and direction. Using a MPMS3-SQUID magnetometer, in-plane magnetisation hysteresis loops were measured of different track lengths taken from structures at the different processing stages (green body and sintering). The results show that there is an increase in saturation magnetisation from green body to sintered sample. The sintered sample showed isotropic magnetisation within the track while the green body showed an anisotropic magnetisation. This may be due to the homogenous grain growth. To evaluate the grid design performance as a strain sensor, a weight was applied to the sintered structure and the change in magnetic flux density was measured. The COMSOL simulations were used to validate the results by using the in-built magnetostriction module. Both simulations and experimental works show comparable results (Fig 2). Demonstrating the larger the track gap the larger the response to an applied force. Future work will explore a range of printable magnetostrictive materials such as nickel and magnetite. The work will explore different designs based on COMSOL simulations including mechanical and magnetic performance. Acknowledgements: We wish to acknowledge the Henry Royce Institute for Advanced Materials, funded through EPSRC grants EP/R00661X/1, EP/S019367/1, EP/P02470X/1 and EP/P025285/1, for the financial support and the Desktop Metal Studio+ system and MPMS3-SQUID access at The University of Sheffield.

**References:**

[1] Z. Leong, W. Holmes, J. Clarke, A. Padki, S. Hayes, and N. A. Morley, IEEE Trans. Magn., 55, 7, 1–6, 2019.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/623f74f02e43c68a9449a77427f35558.jpg)
Fig 1. Additive manufacture sintered samples with track gap of 1.5mm, 3mm and 5mm (left to right)

![](https://assets.underline.io/uploads/markdown_image/1/image/d026f6ebdbd80cd984b1658195b9c6b0.jpg)
Fig 2. Comparison of COMSOL model and experimental data

 Modelling magnetostrictive materials for structural health monitoring of carbon fibre composite

Introduction Rogowski coils (RC) are used for the measurement of large alternating current, particularly as a short circuit detector in electric substations. Another application is in electricity meters. A regular, air-core RC cannot be saturated and therefore its dynamic range is limited only by noise voltage and supply voltage of the accompanying integrator or ADC [1]. An RC used for precise current measurement should have a consistent cross-sectional area and turn-to-turn spacing, limiting the winding arrangement to a single layer. The sensitivity of such coils is poor, causing a low signal-to-noise ratio [2]. Our proposed design adds a low-permeability ferromagnetic core to improve the sensitivity and thus reduce the required number of turns. Ferromagnetic core has been already tested for different types of RC, e.g. self-integrating RCs which are used for measurement of short current pulses. In this case increasing the core permeability decreases the low frequency limit of self-integrating operation [3]. Contrary to that, our target application is precise current sensing at the mains frequency and its harmonics. Such a sensor should be able to work with high accuracy (typical requirement is 0.5% or 0.2%) in the nominal range, as well as above this range with reduced accuracy with partly saturated core as a fault current detector [4]. Because the coil does not require a load resistor on its winding, there is no concern about overheating during overload conditions.However, there are some open questions, how will the ferromagnetic core affect the frequency characteristics of the RC, its ability to suppress external currents and resistance against the position deviations of the measured (primary) conductor. Several different materials, either in the form of a finished toroidal core or a 3-D printer filament filled with magnetic particles were tested, both theoretically by a FEM simulation and practically by measurement. FEM Simulation We have made a series of 3D FEM simulations to evaluate the effect of the core permeability on rejection of external currents and position of the measured conductor. In both cases the perfectly homogenous coil would be absolutely immune against both these influences. Practical coil always has uneven winding density, in particular between the first and last turn [5]. Also the cross-sectional area of individual turns varies slightly. We simulated the role of these imperfections in crosstalk from an external conductor as a function of its position and core permeability. In general, the error is diminishing with increasing permeability due to demagnetization effect. However, for practical permeability of 20 the error reduction is only very small (from 0.23 % to 0.15 %). The permeability effect on the resistance against the position deviations of the measured (primary) conductor is even lower and for permeability of 20 it is negligible. The effect of nonuniformity of the core permeability along its circumference was found to be insignificant as well. Samples used for measurement The first batch of test specimens (labeled A, B, C) are toroidal powder cores with an outer diameter of 79 mm and rated relative permeability of 26 (Kool Mju and Sendust materials) or 14 (Sendust). The winding is electrostatically shielded by a copper adhesive tape, care was taken to avoid the creation of a short-circuited turn. We also tested 3D printed cores from two types of plastic filament filled with magnetic particles and for the full paper we plan to test resin-based cores of various sizes. Measurement procedure Two important sensor characteristics were measured: linearity and frequency response. Test current was provided by a Techron 7548 power amplifier with an additional step-down transformer. Reference current measurement was done by a LEM Ultrastab current sensor (IT60S for frequency response, IT1000S for linearity). Two different measurement approaches were tested: either integrating the induced voltage with a Lakeshore 480 fluxmeter in AC mode, or measuring the voltage directly with a HF2LI lock-in amplifier. Both methods give similar results in amplitude, but the phase can only be measured by a lock-in amplifier. A primary conductor with a loop length of 10 meters was passed once through the measured coil to minimize the influence of the return current. Fig. 1 shows the frequency dependence: there is no observable sensitivity drop due to frequency dependence of the permeability. The sensitivity increase at high frequencies is due to the resonance with parasitic coil capacitance and its amplitude could be reduced by a parallel damping resistor [6] at the expense of heat being generated in the associated circuitry. Fig. 2 shows the measured linearity error at 93 Hz, which is in a good agreement with the catalogue values of permeability vs. flux density in the core. A nominal flux line length of 0.196 m was used to convert flux density to primary current. Conclusion We have demonstrated that adding a ferromagnetic core to a RC of the classical dI/dt type is beneficial for its sensitivity, which would allow the design of very precise RCs for lower currents. The RC linearity is decreased, but remains acceptable for most applications. The amplitude error caused by self-resonance can be reduced by standard methods such as a damping resistor, or by using a different material with high core loss. Acknowledgement This study was supported by the Grant Agency of the Czech Republic within the Nanofluxgate project (GACR GA20-27150S).

**References:**

[1] S. Tumanski, Meas. Sci. Technol. 18, p. R31, 2007 [2] J. D. Ramboz, IEEE Trans. Instrum. Meas. 45, p. 515, 1996. [3] J. D. Zhu, Q. G. Zhang, J. B. Jia et al., Plasma Sci. Tech. 8, p. 457l, 2006. [4] L. A. Kojovic, M. T. Bishop, and D. Sharma, IEEE Trans. Ind. Appl. 49, pp. 8, 2013. [5] L. Ferkovic, D. Ilic, and R. Malaric, IEEE Trans. Instr. Meas. 58, p. 122, 2009. [6] M. H. Samimi, A. Mahari, M. A. Farahnakian et al., IEEE Sens. J. 15, p. 651 (2015).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b316cd355ace0c3db4ef22e26ec8fa2a.jpg)
Fig. 1 Frequency response

![](https://assets.underline.io/uploads/markdown_image/1/image/fd59bddadd4847af0a850bbd099da4d6.jpg)
Fig. 2 Linearity

 Rogowski Coil with Ferromagnetic Core

A growing interest in electric vehicles challenges soft magnetic materials to improve efficiency and power density. Metal amorphous nanocomposites (MANCs) have lower coercivities and eddy current losses, allowing for greater efficiency and higher switching frequency. The later allows for high motor speeds and improved power density. However, commercially available and laboratory demonstrated MANCs have shortcomings. Fe-based MANCs, such as FINEMET have excellent magnetic properties, but mechanical properties limit their application in high speed electric motors (HSMs). Fe-Co based MANCs, such as HITPERM, have high saturation inductions and good mechanical properties, but elevated magnetostrictive losses. Co-based MANCs and more recent Fe-Ni based MANCs have low losses and good mechanical properties, but lower saturation inductions (1-1.2 T), that limit power density in HSMs. Recently, much work has explored improving saturation induction in MANCs by increasing the content of magnetic elements, which comes at the cost of glass-forming ability (GFA). These efforts where generally limited to trial and error testing. Additionally, all work to date has focused on Fe-based alloys. In this work, a method of using Thermocalc simulation to locate near-eutectic compositions has been applied to Fe-Ni based alloys. Minima in liquidus temperature and solidification range were found for a ternary composition range and used to identify compositions that retain good GFA as the percentage of magnetic elements is increased. The (Fe70 Ni30)x(B-Si-Nb)100-x alloy system for x=82% and x=85% was explored by Thermocalc simulations. This is an increase in magnetic element content compared to previously developed x=80% alloys. 3 compositions in the x=82% system, and 1 alloy in the x=85% system where identified and successfully cast as amorphous ribbon. The amorphous nature of the ribbon was confirmed by a bend test and XRD. Magnetic testing was performed by PPMS to measure saturation induction and Curie temperature of the amorphous material. The Curie temperatures increased to 407-438 °C for the x=82% alloys, and 462 °C for the x=85% alloy. Saturation induction increased to 1.28-1.36 T for the x=82% alloy and 1.48 T for the x=85% alloy, which is a significant improvement over the ~370 °C Curie temperature and 1.2 T saturation of previous Fe-Ni alloy. Crystallization behavior was studied by XRD for the x=85% alloy, showing that crystallization process follows a 2-step process of Amorphous → BCC+FCC+Amorphous → BCC+FCC+Fe3B+Fe23B6. Post annealing magnetic properties were studied by strip testing, and saturation induction as high as 1.48 T was seen after optimal annealing. TEM was used to study structure of as cast and annealed material. 2 alloys with x=82% compositions were identified with an optimal crystalline size of 10-20 nm, while the x=85% alloy had this structure in the as cast state. This opens the possibility of using the alloy in the as cast state, without requiring annealing. In conclusion, alloys with good GFA and improved magnetic properties were identified by Thermocalc simulation.

**References:**

 N. Aronhime, V. DeGeorge, V. Keylin, JOM. 69 (2017) 2164–2170 N. Aronhime, E. Zoghlin, V. Keylin, Scr. Mater. 142 (2018) 133–137 J.M. Silveyra, A. Leary, V. DeGeorge, J. Appl. Phys. 115 (2014) 17A319 J.M. Silveyra, P. Xu, V. Keylin, J. Electron. Mater. 45 (2015) 219–225 A. Urata, H. Matsumoto, S. Sato, J. Appl. Phys. 105 (2009) 2007–2010 C. Wang, A. He, A. Wang, J. Pang, Intermetallics. 84 (2017) 142–147 A. Wang, C. Zhao, A. He, J. Alloys Compd. 656 (2016) 729–734 Y. Krimer, N. Aronhime, P. Ohodnicki, J. Alloys Compd. 814 (2020) 1-9

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/1be95150355009fb148c6d8dd5009eca.jpg)

![](https://assets.underline.io/uploads/markdown_image/1/image/46a90665e7662989fbfe624a29693a8d.png)

 Development of High Saturation Induction Fe-Ni Based Metal Amorphous Nanocomposite by Optimization of Glass Forming Ability

The D03 Fe3Sn phase shows higher magnetization and smaller magnetocrystalline anisotropy compared to Fe3Si based on the comprehensive first principles calculations. For this reason, it is important to develop nanocrystalline soft alloys which have D03 Fe3Sn phase as primary crystallite. A new alloy design method, motivated by the work of Villars, was used to choose the alloy composition. Element Co was found to stabilize the D03 Fe3Sn phase. In this work, a new alloy system (Fe1-x-yCoxSny)87Nb3B9Cu1 was studied to achieve D03 Fe3Sn phase with the help of the non-equilibrium processing method. Differential scanning calorimetry was used to determine the primary crystallization temperatures of the amorphous ribbons. The amorphous ribbons were annealed in tube furnace at 723K, 773K, and 823K for 3600 seconds after encapsulating them in a fused quartz ampoule. X-Ray Diffraction (XRD) has been used to analyze structural properties of as-spun and heat-treated ribbons of all alloy compositions. The single broad peak of Fe60Co20Sn7Nb3B9Cu1 as-spun ribbons’ XRD diffractograms indicates an amorphous phase due to their lack of long-range order. The rest as-spun ribbons’ XRD diffractograms shows partial crystalline during melt-spinning process due to the low glass-forming ability of these series of alloys. XRD diffractogram of the heat-treated ribbons are comprised of a B2 phase, which has a BCC structure. The crystallite sizes for each heat-treated sample were estimated by Scherrer broadening to have values between 7 and 10 nm. The magnetic properties of as-spun and heat-treated ribbons were studied by vibrating sample magnetometry. Saturation magnetization and coercivity were determined from room temperature magnetic hysteresis loops. The highest magnetization value achieved was 161.7 Am2/kg for the Fe75Co4Sn8Nb3B9Cu1 ribbon annealed at 723K.

**References:**

P. Villars, “Three-dimensional structural stability diagrams for 648 binary AB3 and 389 binary A3B5 intermetallic compounds: III,” J. Less-Common Met., vol. 102, no. 2, pp. 199–211, 1984 P. Villars, “A three-dimensional structural stability diagram for 998 binary AB intermetallic compounds,” J. Less-Common Met., vol. 92, pp. 215–238, 1983 Alex Zunger, “Systematization of the stable crystal structure of all AB-type binary compounds: A pseudopotential orbital-radii approach,” Phys. Rev. B, vol. 22, no. 12, pp. 5839–5872, 1980 S.S. Batsanov, “Dielectric Methods of Studying the Chemical Bond and the Concept of Electronegativity,” vol. 51, no. 7, 1982.

**Images:**

 Development of an (Fe, Sn)-based Nanocrystalline Soft Magnetic Alloy

INTRODUCTION Fe-based amorphous ribbons play an important role in high efficiency energy conversion devices and have been widely used in the field of electric power, power electronics and renewable energy as transformers [1], inductors [2] and reactors [3]. Small improvements in core losses can have large energy savings [4]. The demands of high efficiency and energy saving pose a challenge to further reduce the core losses of amorphous ribbons. To date, numerous studies have been investigated on materials development [5], subsequent annealing and structure design of magnetic cores [6] to further reduce the core losses of Fe-based amorphous cores and components. However, few reports are found on the influence of process parameters on core losses of Fe-based amorphous ribbon produced by planar flow casting (PFC) process. The objectives of this work are to study the influence of the thermal conductivity of cooling wheel substrate on cooling rate of Fe-based amorphous ribbon, to investigate the effect of cooling rate on core losses of Fe-Si-B amorphous ribbons. MATERIALS AND METHODS A batch of copper beryllium (Cu-2Be) cylindrical rings with different thermal conductivity was manufactured. A cooling wheel system is composed of a stainless steel wheel core and the cylindrical ring. Amorphous alloy ribbons with nominal composition of Fe80Si9B11 (at. %) were produced by using the Cu-2Be as the substrate of cooling wheel, the ribbon width is 50±0.5 mm and the ribbon thickness is 25±0.5 mm. The core losses were measured using IWATSU SY-8232 B-H analyzer. All the measured samples are toroidal cores with inner diameter of 30 mm, outer diameter of 35 mm and height of 50 mm subjected to heat treatment at 365 °C for 60 minutes with applied longitudinal field of 400 A/m. A 2D numerical model was established to calculate heat and mass transfer during PFC process by using ANSYS FLUENT software, the parameters used in the numerical calculation were the same as that of practical casting of Fe-Si-B alloy ribbons. RESULTS AND DISCUSSIONS The PFC process was simulated and the thermal conductivity of 175.3 W/m×K and 206.5 W/m×K for the Cu-2Be rings as a substrate of cooling wheel were used. The main process parameters employed for the simulation are as the follows: the applied pressure at the bottom of nozzle is 45 kPa, the cooling wheel outer diameter is 300 mm, the thickness of the Cu-2Be ring is 10 mm, the linear velocity of cooling wheel surface is 20 m/s, the nozzle slit width is 0.3 mm, the nozzle wheel gap distance is 0.15 mm, the temperature of molten alloy 1420 oC, the initial temperature of Cu-2Be substrate and circumferential air is 32 oC, the glass formation temperature of Fe-Si-B alloy is assumed to be 750 oC. Fig. 1 shows the simulation results of upstream meniscus, downstream meniscus and growth of liquid/solid interface at the 130th revolution for substrate thermal conductivity of 175.3 W/m×K and 206.5 W/m×K, respectively. In Fig. 1, upstream meniscus and down meniscus are plotted by the line of volume fraction of melt=1 and the liquid/solid interface is plotted by the isothermal line of glass formation temperature Tg=750 oC. The growth of liquid/solid interface is from P3 to P4 for this case. The solidification time is t=L/U and the average cooling rate is h=(Tg-Ts)/ t, where L is the puddle length from P3 to P4, U is the linear velocity of cooling wheel surface and Ts is the surface temperature of cooling wheel beneath the puddle (not shown). Based on the data extracted from the liquid/solid interface of the puddle and the surface temperature of cooling wheel beneath the puddle, the average cooling rates are 1.26e-6 K/s and 1.44e-6 K/s for k=175.3 W/m×K and 206.5W/m×K, respectively. The cooling rate increases by 14.3% due to the increase of substrate thermal conductivity. To investigate the cooling rate on core losses of Fe-Si-B alloy during PFC process, Cu-2Be rings with thermal conductivity of 175.3 W/m×K and 206.5 W/m×K were shrink fitted onto the wheel core as the substrate of cooling wheel. Fig. 2 shows the core losses at 50 Hz as a function of induction for Fe-Si-B ribbons produced by cooling wheel with different thermal conductivity. The increase of thermal conductivity of cooling wheel from 175.3 W/m×K to 206.5 W/m×K results in the decrease of core losses from 0.171 W/kg to 0.133 W/kg, with the core losses reduction up to 22.1%. CONCLUSIONS The influence of thermal conductivity of cooling wheel substrate on cooling rate during PFC process was simulated by numerical method and the results show that the increase of thermal conductivity of cooling wheel substrate from 175.3 W/m×K to 206.5 W/m×K leads to the increase of cooling rate from 1.26e-6 K/s to 1.44e-6 K/s in PFC process. The effect of cooling rate on structure and soft magnetic properties of Fe-Si-B alloy ribbons was investigated. The results show that the increase of thermal conductivity of cooling wheel substrate from 175.3 W/m×K to 206.5 W/m×K during PFC process results in the decrease of core losses from 0.171 W/kg to 0.133 W/kg, with the core losses reduction up to 22.1%.

**References:**

[1] D. Azuma, R. Hasegawa, IEEE Trans. Magn., 44, 4104-4106 (2008) [2] R. Hasegawa, J. Non-Cryst. Solids, 387, 405-412 (2001) [3] N. Kurita, K. Onda, K. Nakanoue, IEEE Trans. Power Electr., 29, 3657-3668 (2014) [4] O. Gutfleisch, M. A. Willard, E. Brück, Adv. Mater., 23, 821-842 (2011) [5] C. Dong, A. Inoue, X. H. Wang, J. Non-Cryst. Solids, 500, 173-180 (2018) [6] D. Azuma, N. Ito, M. Ohta, J. Magn. Magn. Mater., 501, 166373 (2020)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/eb7f8d32dca65c6294b9b09eda7cc5f5.jpg)
Fig. 1 The simulation results of upstream meniscus, downstream meniscus and growth of liquid/solid interface in quasi-static process for substrate thermal conductivity of 175.3 W/m×K and 206.5 W/m×K.

![](https://assets.underline.io/uploads/markdown_image/1/image/544d845cdb3c2ee849ef1324426ec190.jpg)
Fig. 2 Core losses at 50 Hz as a function of induction for Fe-Si-B ribbons produced by cooling wheel with different thermal conductivity.

 Reducing the Core Losses of Fe-Si-B Amorphous Ribbons by High Cooling Rate Planar Flow Casting

Fe-Co-Ni magnetic thin films with continuously varying compositions have been fabricated by the magnetron co-sputtering method. The films were characterized by high-throughput techniques. The results reveal the correlation between a range of compositions and the electrical properties of the films fabricated at various process parameters. Introduction: Soft magnetic materials are essential elements of electro-magnetic energy transformation technologies and they are widely used in various distribution, conversion, generation devices, such as transformers, motors, converters, generators, actuators, sensors etc. However, it usually takes long time for a new material before its acceptance for commercial application. Therefore, it is very crucial to speed up the discovery of new materials and incorporate them into various systems. High-throughput methods are efficient research methods [1-3] to explore this immense search space to identify new or optimized materials. Fe-Co based alloys received great attention due to high saturation magnetization. And what makes them even more fascinating is that their magnetic and electric properties can be tuned by controlling the alloy composition and microstructure. Among these compositions, Fe-Co-Ni films revealed significant induced magnetic anisotropy, low coercivity and high anisotropy of magnetoresistivity, what makes them a good candidate for technical applications. The resistivity is an important parameter to determine power loss in many soft magnetic applications. In this work we studied the effect of process parameters on the structure and resistivity of compositionally graded Fe-Co-Ni films to find process conditions at which optimum properties are obtained. Experimental details: Fe-Co-Ni films with continuously varying alloy compositions were fabricated by the magnetron sputtering method. Two experimental procedures were designed: a) Fe-Co-Ni films were deposited on SiO2/Si substrates by co-sputtering of permalloy and Co target materials using DC sources at different substrate temperatures (Ts): in particular, ambient temperature, 300 °C and 500 °C; b) the film was deposited at ambient temperature and annealed for 2 hours in vacuum at 500 °C. The base pressure in the chamber was controlled at 4 x 10 - 4 Pa. Depositions were performed in argon gas environment while the pressure was kept approximately at 0.3 Pa. For compositional, structural and electrical properties high-throughput Electron Probe Micro-Analysis (EPMA), X-ray Diffraction, four-point probe measurements were performed to determine chemical compositions, structure and resistivities of the films. Results and discussion: Fe-Co-Ni films with continuously varying alloy compositions were fabricated by magnetron sputtering method. The results of high-throughput Electron Probe Micro-Analysis measurements across Fe-Co-Ni films (Fig.1) revealed that atomic composition of Ni was highest, while Co was lowest in the film fabricated at Ts = 300 °C. The ratios of Co and Ni atomic compositions in deposited graded films were calculated as well. We found that Fe-Co-Ni film fabricated at Ts = 300 °C shows the lowest ratio as compared to other films. XRD measurements and analysis of multiple selected areas on Fe-Co-Ni film fabricated at ambient substrate temperature exhibited amorphous structure. As the substrate temperature increased diffraction peaks appeared, showing a clear change in structure. The diffraction patterns showed that only one FCC phase appeared. The peak intensity increased with temperature and the grain sizes became larger. The increase of peak intensity and average grain size indicated that increasing the substrate temperature from room temperature to 500 °C enhances the crystallinity of Fe-Co-Ni films. As for the annealed samples the intensity of the peak was low, and peak shifted towards lower angles indicating crystal lattice expansion. In addition, oxide peaks have been observed. Fig. 2 shows the measured resistivities as a function of composition of various process parameters. Resistivities of the graded film deposited at Ts =300 °C exhibited higher values than resistivities of the films deposited at ambient, 500 °C, and at ambient temperature followed by annealing at 500 °C. These results indicate that Ni plays important role in tuning resistivities of the films. Conclusion: Fe-Co-Ni films with continuously varying compositions have been fabricated by magnetron sputtering and characterized by high-throughput experimentation methods. The results revealed the correlation between continuously varying alloy compositions and selected properties of the films fabricated at various process parameters. Acknowledgement: This work is supported by the AME Programmatic Fund by the Agency for Science, Technology and Research, Singapore under Grant No. A1898b0043

**References:**

[1] Sean W. Fackler, Vasileios Alexandrakis, Dennis König, A. Gilad Kusne, Tieren Gao, Matthew J. Kramer, Drew Stasak, Kenny Lopez, Brad Zayac, Apurva Mehta, Alfred Ludwig and Ichiro Takeuchi, Science and Technology of Advanced Materials 2017, Vol.18, No.1, 231–238 [2] H. Koinuma, I. Takeuchi: Nat. Mater. 3 (2004) 429 –438. [3] Alfred Ludwig, npj Computational Materials (2019) 5:70

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/a7c2b3186ab8ed446ef672dcbe850a63.jpg)
Fig. 1 Elemental compositions of the graded Fe-Co-Ni films

![](https://assets.underline.io/uploads/markdown_image/1/image/dcd0207cd6dc3ac30e5af139256a3be1.jpg)
Fig. 2 Resistivities of Fe-Co-Ni films with continuously varying alloy compositions fabricated at different process parameters

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