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A novel approach for fabricating Ni-coated FeSiAl soft magnetic composite via cold spraying
Journal of Alloys and Compounds 749 (2018) 523e533
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
A novel approach for fabricating Ni-coated FeSiAl soft magnetic composite via cold spraying Xinliang Xie a, Chaoyue Chen a, *, Yingchun Xie a, b, Zhongming Ren c, Eric Aubry d, Gang Ji e, Hanlin Liao a ICB UMR 6303, CNRS, Univ. Bourgogne Franche-Comt e, UTBM, F-90010 Belfort, France National Engineering Laboratory for Modern Materials Surface Engineering Technology, The Key Lab of Guangdong for Modern Surface Engineering Technology, Guangdong Institute of New Materials, Guangzhou 510651, PR China c Shanghai University & State Key Laboratory of Advanced Special Steel, 149 Yanchang Road, Shanghai 200072, PR China d Nipson Technology, 12 Avenue des Trois Ch^ enes, 90000 Belfort, France e Unit e Mat eriaux et Transformations, CNRS UMR 8207, Universit e Lille 1, Villeneuve d'Ascq, 59655, France a
b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 January 2018 Received in revised form 23 March 2018 Accepted 24 March 2018 Available online 27 March 2018
Currently, soft magnetic composites (SMCs) have attracted increasing attention due to their outstanding magnetic properties, and various methods have been developed and applied for their fabrication. As an emerging additive manufacturing technique, cold spraying (CS) can fabricate bulk material via solid-state deposition by avoiding oxidation and phase change. In this work, SMCs coating was first fabricated by Nicoated FeSiAl composite powder via CS. Two groups of Ni-coated FeSiAl composite particles (40 mm and 57 mm) were used as feedstocks for deposition under different processing parameters. No phase transformation can be detected from XRD analysis. The coating thickness increased with the increasing of propelling gas temperature and pressure. Higher deformation of FeSiAl particles and higher microhardness of the composites fabricated from the powders with larger size were obtained due to the enhanced in-situ peening effect of the rebounded particles during deposition. The magnetic property of cold sprayed SMCs showed a soft ferromagnetic characteristic with a coercivity of about 60 Oe. The investigation on the coating formation mechanism was carried out by single particle deposition, and the results showed that the Ni bonding layer with sufficient plastic deformation plays a significant role during the deposition of the composite coating. © 2018 Elsevier B.V. All rights reserved.
Keywords: Cold spraying Soft magnetic composites (SMCs) FeSiAl alloy Cladded powder
1. Introduction Due to their unique magnetic properties such as high permeability and low magnetic coercivity, the soft magnetic materials have been widely used in electrical and electronic devices [1,2]. Generally, the soft magnetic materials can be classified into soft magnetic alloys, soft magnetic oxides and soft magnetic composites (SMCs) [3]. Based on features like low core loss, higher saturation magnetization and electrical resistivity, SMCs have more competitive advantages in various applications and are attracting more and more attention [4e6]. For the moment, various material systems of SMCs have been studied in the literature, such as FeeSi [7], Fe-Ni [8], Fe-Si-Al [3,6,9], Fe-Ni-Mo [10] and Fe-Cu-Nb-Si-B [11].
* Corresponding author. E-mail address: [email protected] (C. Chen). https://doi.org/10.1016/j.jallcom.2018.03.306 0925-8388/© 2018 Elsevier B.V. All rights reserved.
Theoretically, ideal SMCs should not only have high saturation flux density but also stable permeability. Among the above-mentioned systems, FeSiAl-based SMCs have become attractive candidates in various electronic devices with their outstanding soft magnetic properties [9,12,13]. Up to now, various methods and techniques have been developed and used to fabricate the desired SMCs. Traditionally, the most common technique for the production of SMCs is the powder metallurgy [11,14]. In addition, different methods have been applied based on powder compaction following sintering or annealing treatments [15]. Recently, efforts have been made to fabricate FeSi based SMCs coatings by thermal spray technologies such as high velocity oxygen fuel (HVOF) [16,17] and atmospheric plasma spray (APS) [18]. In such thermal spray processes, the molten particles are deposited onto substrate through the acceleration of high-energy sources. The inevitable melting caused by high-energy source can lead to unexpected phase change and
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further affect the magnetic properties. Besides, the unavoidable high porosity and serious oxidation can also seriously weaken the magnetic performance of the thermal sprayed SMCs coatings. As an emerging technique, cold spraying (CS) provides an effective approach to produce high quality coatings of metallic, metal matrix composite even ceramic [19e22]. In this process, the feedstock particles are accelerated by supersonic carrying gas through a Laval nozzle, and then the coating is formed through the intensive plastic deformation of the deposited particles on the substrate. Owing to the unique cold feature, the solid-state particles undergo almost no melting during CS process. Unlike the melting and solidification process in conventional thermal spraying, the cold sprayed coating is formed through the solid-state consolidation via mechanical interlocking and metallurgical bonding [23,24]. Therefore, the cold sprayed coating can effectively avoid the serious oxidation, phase transformation and thermal residual stress that are commonly found in traditional thermal sprayed samples. As a result, the aforementioned benefits enable the CS technique to produce various high-quality coatings without the size limitation. Currently, CS technique has been widely considered as an effective method for additive manufacturing and damage repair [25,26]. Up to now, very few attempts have been carried out to fabricate soft magnetic coatings via CS. M. Cherigui et al. [27] fabricated FeSiBNbCu/Al composite coatings via CS of nanostructured powders mixed with Al powders. The as-sprayed coatings showed a soft ferromagnetic property despite the presence of Al, which was considered as a nonmagnetic material [27]. However, the addition of a nonmagnetic Al layer greatly reduced the overall magnetic permeability, which yielded a problematic balance of properties [5]. According to the reports by W.Y Li et al. [28], nanostructured Fe-Si coatings were deposited by using mechanically milled powders as the feedstocks. The fabricated Fe-Si coating exhibited a high coercivity (190 kA/m) and had the potential for the application to magnetic recording materials [28]. However, due to the introduction of oxygen and nanostructure to the feedstocks during mechanical alloying, the coercivity of the nanostructured Fe-Si coating was three to four times higher in magnitude than the maghemite iron oxide (20-30 kA/m) [28]. Thus, it should be noted that the choice of a proper feedstock for CS deposition of SMCs is of significant importance to the soft magnetic properties. In this study, the Ni/FeSiAl soft magnetic composite was fabricated by the CS technology. Because of their low ductility, FeSiAl particles are difficult to be directly deposited via CS. In order to improve the deposition efficiency of the hard FeSiAl particles, a novel composite powder where FeSiAl core is decorated by thin Ni layer was used as feedstock to fabricate SMCs coatings. Ni was selected for the matrix phase because of its soft magnetic property and its excellent ductility. In addition, Ni shows a relatively high resistance to corrosion and wear. The aim of this work is to examine the effect of powder size distribution range and processing parameters on the microstructure, microhardness and magnetic properties of SMCs coatings. The coating microstructure, particle deformation and retainability of the FeSiAl particles were characterized and estimated. The soft magnetic properties of the assprayed SMCs coatings were evaluated. At last, the deposition mechanism of the composite powder was discussed.
as powder 1 (35-87 mm) and powder 2 (28-58 mm), were used as feedstocks in this study. The detailed size distributions of these two feedstocks measured by a laser diffraction sizer are given in Fig. 1. The surface morphologies of the feedstock powders given in Fig. 2 exhibits a near spherical shape. As the cross-sectional morphologies given in Fig. 2 (b) and (d), the composite particles exhibit a core-shell structure. A thin Ni layer (about 3e5 mm) is uniformly decorated around a FeSiAl particle. As for some smaller FeSiAl particles, they are aggregated and covered with a thicker Ni layer varying from 6 to 14 mm. The average volume fractions of FeSiAl particle in the core-shell structured composite powder 1 and powder 2 were estimated as 62% and 60%, respectively. 2.2. Coating fabrication FeSiAl/Ni composite coatings and splats were produced by using a homemade CS system (LERMPS, UTBM, France) with an optimal de-Laval-type converging-diverging nozzle. High pressure compressed air was used as powder carrier gas and propelling gas. The powders were deposited onto stainless steel substrate under different propelling gas pressures and temperatures, which are listed in Table 1. Nozzle standoff distance and traverse speed were set as 30 mm and 100 mm/s, respectively. Nozzle trajectory was repeated for 20 times to obtain a thick coating. The spacing between adjacent nozzle passes was set as 2 mm. In addition, for understanding the coating formation mechanism, a single particle deposition onto a stainless steel substrate was also conducted at the nozzle traverse speed of 500 mm/s. The stainless steel substrates were polished or grit blasted in order to get individual splat deposition and full coating deposition. 2.3. Materials characterization In order to examine the phase transformation during CS process, the as-sprayed coatings as well as the feedstock powders were examined by an X-ray diffractometer (Siemens D500, Germany) with the Co (l ¼ 1.78897 Å) source at a current of 40 mA, voltage of 35 kV and scan step of 0.02 . A scanning electron microscopy (SEM JEOL-5900LV, Japan) was used to characterize the morphology and microstructure of the powders and the coatings. Prior to electron microscopy, the as-sprayed coating samples were prepared using standard metallographic procedures with the final polishing
2. Experimental procedure 2.1. Starting materials Commercial gas-atomized Fe-9.6%Si-5.4%Al particles were precoated with a Ni layer by hydrometallurgy process (Beijing General Research Institute of Nonferrous Metals, China). Two types of Ni-coated FeSiAl powders with different size distributions, referred
Fig. 1. Particle size distributions of two feedstock powders.
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Fig. 2. SEM micrographs of Ni/FeSiAl composite powders depicting top view and as-polished cross-section: (a) and (b) powder 1, (c) and (d) powder 2.
Table 1 Propelling gas pressure and gas temperature for deposition conditions 1 to 4. Deposition condition
Pressure (MPa)
Temperature ( C)
C1 C2 C3 C4
2.6 2.6 3.0 3.0
550 620 620 650
applied by 0.05 mm Al2O3 solution. The volume fraction analysis of FeSiAl particle in powders and coatings were performed on the cross-sectional micrographs with the 300 and 500 magnification using the imaging software (Image J), respectively. For each sample, five locations were selected from the cross-section of polished coating and the measured data were then averaged. The flattening ratio of FeSiAl particle was obtained by measuring the longest length of the particle over the shortest length of the particle. For flattening ratio measurement of the FeSiAl particles, an average value was obtained using a minimum of 40 particles per sample. Electron back scattered diffraction (EBSD) scans were recorded at electron energy of 25 KeV, and with a step size of 100 nm using the FEI-Sirion SEM equipped with a TSL-OIM EBSD camera. EBSD samples were sectioned from the CS deposition and prepared by standard metallographic techniques, and the final polishing was conducted using a 0.05 mm diameter colloidal silica suspension and vibratory polishing. The coating microhardness was measured by a Vickers hardness indenter (Leitz, Germany) with a load of 100 g for 10 s. 10 positions were randomly tested and averaged to determine the microhardness. The magnetic measurements for the powders and the SMCs coatings were carried out with a hysteresismeter Bull M 2000/2010 SIIS apparatus (Niposon technology) at ambient temperature. In order to prepare the samples for magnetic property test, SMCs coatings were deposited onto grit blasted Al plate with the propelling gas pressure and temperature of 3.0 MPa and 650 C (C4), respectively. Then, the 10 mm 10 mm cut samples were placed in
high-concentration sodium hydroxide solution for several hours to fully dissolve the Al substrate; while the SMCs coating can be well retained during corrosion. Before the magnetic test, the composite coatings with a thickness of 300e400 mm were grounded to remove the surface oxide films.
3. Results 3.1. Phase composition Fig. 3(a) shows the XRD patterns of the initial powder and CS coating fabricated at C4. Besides the main phases of Ni and Al0.3Si0.7Fe3, no other potential phases appear on the XRD patterns of cold sprayed SMCs coatings, indicating that neither significant chemical reaction nor phase transformation occurred during deposition. This is typically attributed to the low temperature characteristic of the CS process. One can observe that the (111) peek of Ni is much stronger than the (110) peek of Al0.3Si0.7Fe3 phase even though the volume fraction of Ni (40 vol%) is lower than that of Al0.3Si0.7Fe3 (60 vol%) in the powder. This is due to the depth limitation of X-ray diffraction and the core-shell structure of the feedstock powder, where most of the X-rays are diffracted by the Ni coating layers (several microns). Fig. 3(b) shows the detailed XRD spectra of the 2q angle between 50 and 55 . It can be observed that the Ni diffraction peaks of the coatings slightly shifted to the right in comparison with that of the initial powder. The Al0.3Si0.7Fe3 (110) peeks show the same trend despite their low intensities. These results indicate that the lattice parameters of both the Ni and Al0.3Si0.7Fe3 phases are somewhat smaller in comparison with those of the feedstock. This can be explained by the severe plastic deformation of particles during the high-velocity impact in CS [29], and further investigation will be carried out by EBSD characterization in the following sections.
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Fig. 3. X-ray diffraction pattern of cold sprayed SMCs coatings fabricated at C4 and feedstock powder.
3.2. Microstructure Fig. 4 shows the cross-sectional images of the Ni/FeSiAl SMCs coatings fabricated by different feedstock powders. At the coating/ substrate interface, the intimate bonding can be observed without obvious cracks. The coatings exhibit high densification rate and crack-free features. As indicated by the red arrows, a few microsized pores can be observed in the cross-sectional images, which have the similar shapes with the remaining FeSiAl particles. Such micro-sized pores should be attributed to the detachment of FeSiAl particles during the sample preparation rather than the CS process. Statistical analysis of the cross-sectional SEM images at 500 yields a porosity less than 1% for all the as-sprayed coatings, where no evident inter-particle cracks or pores can be observed. As shown in Fig. 4 (a), a thinner SMCs coating with high surface roughness even serious delamination can be seen at low processing parameters. It can be explained by the strong peening and rebounding of the unsuccessful deposition particles due to the low deposition efficiency. From C2 to C4, the improved deposition efficiency with the increasing propelling gas pressure and temperature significantly increases the coating thickness (see Fig. 4 (b) to (d)). It suggests that higher propelling gas pressure and temperature are beneficial for the deposition of composite particle. As for the SMCs coatings produced by the starting powder with a different size distribution, similar tendency of coating thickness can be observed from C1 to C4. The variation of coating thickness with the processing parameters of these two types of feedstock
powders is given in Fig. 5. While comparing the SMCs coatings at the same processing parameters, a thicker coating was obtained from the finer powder, indicating a much higher particle deposition efficiency. According to the deposition window in terms of particle size given by Schmidt et al., the particle obtained a higher in-flight velocity at a smaller particle size within a certain range [30]. Thus, a higher impact velocity of composite particle contributes to a thicker
Fig. 5. Variation of coating thickness with processing parameters.
Fig. 4. Cross sections of the SMCs coatings fabricated from powder 1 (aed) and powder 2 (eeh) under different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
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coating for the finer starting powder. It is commonly considered that the uniform distribution of the reinforced particles in the composite coating is the key to the desired coating properties. As shown in Fig. 4 (a) and (e), the FeSiAl particles were found concentrated on the top of the coating, while the bottom mainly consists of Ni layers. In the cases from C2 to C4, the FeSiAl particles can be observed uniformly distributed in the SMCs coatings as the increase of processing parameters. Besides, the volume fraction of FeSiAl particles was measured to evaluate the FeSiAl content in the SMCs coatings, and the results are given in Table 2. The volume fraction of FeSiAl particles in all the coatings were calculated to be less than 30%, which is half of the value in initial powder (~60 vol%). This indicates that many FeSiAl particles were lost during the coating build-up process. As the propelling gas pressure and temperature decrease, more FeSiAl particles rebound away from the coating surface, which results in less FeSiAl content in the coatings. For each condition, the average volume fraction of FeSiAl in the coatings obtained from powder 2 is slightly larger than that from powder 1. In addition, it can be noticed that only a few of large FeSiAl particles retain in the coatings. The results demonstrate that finer particles possess a higher deposition efficiency (see Table 3). Fig. 6 shows the magnified views of the cross-sectional morphologies of the cold sprayed coatings. The FeSiAl particles are intimately bonded with the ductile Ni binder phase. As shown in Fig. 6 (a) and (e) with relatively low processing parameters, the FeSiAl particles are highly deformed and elongated at the direction perpendicular to deposition. In the case of C1, the flattening ratio of FeSiAl particles were measured to be 2.52, showing an increment about 131.2% in comparison with the initial powder (flattening ratio ¼ 1.14). As the propelling gas pressure and temperature increase from C1 to C4, the flattening ratio of FeSiAl particle gradually reduced from 2.52 to 1.68, indicating lower particle deformation at a higher propelling gas pressure and temperature. Generally, the deformation of cold sprayed particles mainly depends on two major facts: the impact of deposited particle and the peening or tamping effect of subsequent particle impacts [31]. High processing parameters can lead to higher particle impact velocity, which is beneficial for particle deformation and can result in higher deposition efficiency. However, much higher particle deformation has been observed at C1 with relatively low processing parameters. The extremely low deposition efficiency of FeSiAl particles at such low processing parameters leads to the enhanced peening effects during deposition. Many FeSiAl particles rebounded from the coating surface owing to unsuccessful bonding. These particles were acting as in-situ shot peening particles, inducing hammering effect on the previous deposited layers. Thus, more prominent deformation of FeSiAl particles can be found in the cases of low processing parameters. Meanwhile, in the view of starting particle sizes, the average flattening ratio of FeSiAl particle in the coating obtained from powder 1 is larger than that from powder 2 under the same conditions. Because of the lower deposition efficiency and larger kinetic energy of the coarser particles, the enhanced in-situ peening effect directly leads to the severe deformation of FeSiAl particles during deposition. EBSD characterization was applied to further investigate the particle deformation behavior and microstructure of cold sprayed
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coatings. Fig. 7 shows the EBSD result of the SMCs coating obtained from powder 1 at C4. As it can be observed from Fig. 8 (a), the elongated FeSiAl particles are in large grain size, and they are surrounded by a plenty of randomly oriented ultrafine Ni grains. The distribution of Ni grain size is given in Fig. 7 (b) and its average grain size is measured to be 0.57 mm. Grain boundary distribution in Fig. 7 (c) shows that most of the Ni grains are in large grain boundaries, as described in the red lines, while small grain boundaries (blue lines) or even some subgrains can be observed within the FeSiAl particles. The formation of these nanometer-sized Ni grains is due to the dynamic recrystallization during the severe plastic deformation and thermal softening process [32,33]. Fig. 7 (d) is a kernel average misorientation (KAM) map, which is a measure of plastic strain. Higher plastic strain can be observed within the FeSiAl particles rather than that in Ni layers. Such phenomena demonstrate that considerable strain stress was generated and concentrated within the FeSiAl particles during the high-velocity particle impact. The crystal orientation evolution within a FeSiAl grain is shown in Fig. 8. It can be seen that the lattice is gradually rotated along the line ABCD and their orientation evolution is illustrated in the IPF image (Fig. 8 (b)). The results indicate that deformation induced strain stress may lead to the formation of subgrains or small grain boundaries in the hard FeSiAl particles.
3.3. Surface morphology Fig. 9 displays the top-view surface morphologies of the coatings obtained under different conditions. A common feature is the presence of craters for all the coatings. As for the coatings obtained from powder 1 (see Fig. 9(aed)), relatively flat surface with small craters can be observed. However, the coating surfaces obtained from powder 2 present several large craters (about 200 mm) (see Fig. 9(eeh)). The formation of such craters could be attributed to the erosion effect of the high velocity particles during deposition, especially for the coatings deposited from the finer powders. Once small craters were formed, particles became more and more difficult to deposit on the wall of the craters, due to the fact that the deposition efficiency decreases with the angle of the projection [34,35]. As a result, large craters were gradually formed with the growth in coating thickness. In addition, many traces of the shapes of their corresponding powder particles are observed on the coating surfaces. The magnified images in Fig. 10 show the representative surface morphologies in four different cases. In the first case, the entire particles rebounded away with only traces of particles remaining on a highly deformed surface. In the second case (see Fig. 10 (b)), a FeSiAl particle as the core part bounced off leaving a highly deformed Ni layer on the coating surface. In the third case, a successful deposit particle was eroded by the incoming particles with the upper Ni layer being fractured. These exposed FeSiAl particles will prevent the subsequent particles from deposition. In the final case, particles were successful bonded with previous deposited layers. These deposition morphologies allow us to understand the coating formation mechanisms, which will be discussed in the following section.
Table 2 Volume fraction of FeSiAl particles in the SMCs coatings. Volume fraction
Initial powder
C1
C2
C3
C4
Powder 1 (d0.5 ¼ 56 mm) Powder 2 (d0.5 ¼ 40 mm)
62.0 ± 1.5% 60.0 ± 1.2%
19.1 ± 0.9% 22.1 ± 1.8%
23.7 ± 2.4% 24.3 ± 1.6%
25.4 ± 2.0% 27.4 ± 1.5%
28.6 ± 2.3% 29.8 ± 2.8%
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Table 3 Flattening ratio variation of the FeSiAl particles under different processing parameters. Flattening ratio
Initial powder
C1
C2
C3
C4
Powder 1 (d0.5 ¼ 56 mm) Powder 2 (d0.5 ¼ 40 mm)
1.14 ± 0.12 1.18 ± 0.16
2.52 ± 0.96 (131.2%) 1.96 ± 0.71 (79.8%)
1.86 ± 0.55 (68.8%) 1.73 ± 0.58 (58.7%)
1.82 ± 0.50 (67.9%) 1.68 ± 0.59 (57.8%)
1.68 ± 0.51 (54.0%) 1.56 ± 0.44 (43.1%)
Fig. 6. Magnified views of the cross-sectional morphology of cold spray coatings fabricated from powder 1 (aed) and powder 2 (eeh) with different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 7. EBSD characterization of the cross-section of the as-sprayed coating fabricated at C4: (a) IPFþIQ map, (b) Ni grain size distribution, (c) Grain boundary distribution, (e) Kernel average misorientation (KAM) map.
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3.4. Microhardness
Fig. 8. A close view of a FeSiAl particle in Fig. 7: (a) Euler angle map with four points (A, B, C and D) marked in a line from left to right, (b) Inverse pole figure of the points in the line ABCD.
Fig. 11 shows the microhardness of SMCs coatings fabricated under different conditions. Clearly, higher coating microhardness is obtained from the coarser powder despite the fact that the volume fraction of FeSiAl particles is lower than that in the finer powder. In addition, coatings deposited with the coarser powder exhibite smaller variation in microhardness values compared with that of the finer powder. For the coarser powder, the microhardness value decreases slightly with the increasing of processing parameters, and C1 coating presents the highest value even though it has the lowest FeSiAl volume fraction. The SMCs coatings obtained from the finer powders demonstrate the similar tendency of microhardness as a function of processing parameters. Fig. 12 shows the morphologies of the Vickers indentations on the coating cross-
Fig. 9. Surface morphologies of the cold sprayed SMCs coatings fabricated from powder 1 (aed) and powder 2 (eeh) at different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 10. Typical surface morphologies of the Ni/FeSiAl SMCs coating in magnified views.
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particles. At low processing parameters, the low deposition efficiency leads to the unsuccessful bonding of the FeSiAl particles. The rebounding of FeSiAl particles can result in the enhanced peening effect. Due to the in-situ hammering by these shot peening particles, plastic deformation of previously deposited layers could be greatly enhanced and thereby the improvement of interparticle bonding [38,39]. In addition, the induced plastic deformation via the hammering effect can also lead to dislocation accumulation and grain refinement for the previous deposited layers [38], which in turn results in the improvement of coating hardness and toughness. 3.5. Magnetic property
Fig. 11. Microhardness of SMCs coatings under different conditions.
sections. The coatings fabricated from powder 1 present a smaller indentation size and the smallest is shown in C1. Cracks are found in the coatings after indentation test, especially for the coatings fabricated from the fine powder. These results indicate that an improved coating toughness can be obtained from the coarser powder due to the enhanced peening effect during deposition. Generally, it is considered that a higher content of reinforced particles can significantly improve the microhardness of composite coatings [36,37]. However, the highest value of microhardness was obtained in the coating with the lowest volume fraction of FeSiAl
Fig. 13 shows the hysteresis loops of the Ni/FeSiAl SMCs coatings. Their corresponding coercivity and saturation magnetization values are given in Table 4. These coatings obtained from powder 1 and powder 2 exhibit the similar soft magnetic properties with the coercivity of 61.5 Oe and 62.2 Oe, respectively, which have a slight increment compared to that of the feedstock powder (48 Oe); while the saturation magnetizations of the coatings exhibit a slight decrease compared with that of the feedstock powder. The increase in the coercivity value can be explained by the followed reasons: (i) As mentioned above, a large amount of FeSiAl particles was lost during deposition. Consequently, only half volume fraction of FeSiAl particles were retained in the coating; (ii) The generation of strain stress and dislocations in the plastic deformed FeSiAl particles during impact can lead to the increase of coercivity. (iii) It is well known that the crystallite size has a great effect on the soft magnetic performance [40]. The coercivity increases with the crystal size reduction [41]. Even though grain refinement within the FeSiAl particles was not observed, the nano-sized Ni grains
Fig. 12. Morphology of the Vickers indentations on the coating cross-sections fabricated from powder 1 (aed) and powder 2 (eeh) at different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 13. (a) Hysteresis loops of the cold sprayed coatings and (b) magnification of the circled area.
X. Xie et al. / Journal of Alloys and Compounds 749 (2018) 523e533 Table 4 Coercivity and saturation magnetization of initial powder and SMCs coatings. Sample
Coercivity Hc (Oe)
Saturation magnetization Ms (emu/g)
Powder 1 Powder 1-C4 Powder 2-C4
48.4 61.5 62.2
80.5 68.2 79.7
derived from dynamic recrystallization may increase the coating coercivity. In the future plan, to further improve the soft magnetic property of the composite coating, post heat treatment on the assprayed coatings will be carried out to relax strain stress and enlarge grain size.
4. Discussion In CS, the successful deposition of a single ductile metallic particle on the substrate or on previous deposited layers is largely attributed to the occurrence of adiabatic shear instability [42]. This is a phenomenon that occurs in the particle when it undergoes high strain rate deformation upon impact. It is known that hard reinforcements (e.g. BN [43], Al2O3 [44], diamond [45], WC-Co [37], SiC [36]) during the CS deposition of metal matrix composite coatings do not undergo plastic deformation. Due to the lack of sufficient plastic deformation, these hard particles tend to rebound away from the coating surface rather than successful bonding. Similarly, it is also found from the present work that the volume fraction of FeSiAl particle was decreased compared with the starting feedstocks, which indicates their rebounding during deposition. Meanwhile, the hard FeSiAl core experienced slight deformation and no metallurgical bonding occurred with the soft Ni layer (Figs. 4 and 7). To further understand the influencing factors of magnetic and mechanical properties, the formation mechanism of cold sprayed SMCs coatings is discussed through the single composite particle deposition. Fig. 14 displays the surface and crosssectional morphologies of the single Ni/FeSiAl particle onto SS substrate. Four typical morphologies are given indicating the successful bonding, particle rebounding and complete rebounding of a single composite particle. In the first case (Fig. 14 (a) and (e)), the successful deposition of a composite particle can be seen on the SS substrate with strong deformation of the Ni layer. In Fig. 14 (e), the metal jet can be observed at the rim of the deformed particle, which gives a direct proof for the metallurgical bonding between the composite particle and substrate. During the impact of the composite particle, the ductile Ni layer experienced severe plastic
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deformation and thermal softening. As a consequence, in this region as adiabatic shear instability occurred, Ni layer behaves like a viscous metal, which results in an extrusive outward metal jet at the rim. As shown in Fig. 14 (b) and (f), the Ni layer around FeSiAl was fractured with large cracks, and a rim of metal jet can be found around the deformed particle. Different from the deposition behavior of ductile Ni layer, more elastic nature of FeSiAl particle can be expected. Since the kinetic energy of FeSiAl particle was partly transformed into the plastic deformation, the remaining was stored as elastic energy that led to the rebounding of FeSiAl particle. Thus, the observed crack on Ni layer can be formed as the rebounding energy exceeding its strength limitation. As seen in Fig. 14 (c) and (g), after the rebounding of FeSiAl, part of Ni layer still remains in the substrate crater owing to the existence of metallurgical bonding between Ni and substrate. Due to the high kinetic energy and insufficient plastic, the rebounding of FeSiAl particle caused the complete fracture of Ni bonding layer in the composite powder. In the last case, due to the absence of metallic bonding with the substrate, the composite particle experienced a complete rebound off the substrate surface. As it can be seen in Fig. 14 (d), the substrate experienced a high degree of deformation, leaving deep indentations of the rebounded particle. It indicates that this thin layer of Ni sometimes cannot achieve metallic bonding with the substrate due to the insufficient particle deformation, for example, the already deposited particle can be removed from the substrate by the subsequent particles. Based on the deposition behavior of the single composite particle, the formation mechanism of the composite coating in this work is investigated and discussed. Fig. 15 shows the schematic of the deposition process for Ni/FeSiAl composite coating. It can be seen that the composite particle with a shell-core structure experienced three stages during coating build-up. As the composite particle impacts onto the previous deposited Ni layers with sufficient thickness in the first stage, the metallurgical bonding occurs between the Ni bonding layers of composite particles. The sufficient plastic deformation can effectively remove the broken oxide-films on Ni layer interface and promote the successful deposition process by introducing metallurgical bonding. In the second and last stages, FeSiAl core or the entire Ni/FeSiAl composite particle experience rebounding after impact due to the absence of high bonding strength between the Ni bonding layers. As for the FeSiAl core, the high kinetic energy as a result of high velocity is mostly stored as elastic energy rather than dissipated in the forms of plastic deformation. Thus, the rebounding of FeSiAl particles can be
Fig. 14. Surface and cross-sectional morphologies of a single splat deposited onto SS substrate in different cases, red arrows showing the material jet. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 15. Schematic of the deposition process for Ni-coated FeSiAl coating.
observed as the bonding between Ni layer is weak, which attributes to the decreasing content of FeSiAl in the as-sprayed SMCs coating. The metallurgical bonding during the composite coating built-up process only takes place at the interfaces between the Ni layers or between the Ni layer and substrate. It can be concluded that this soft Ni layer plays a crucial role on the formation process of composite coatings. At the same time, these rebounded particles act as in-situ peening particles which can further enhance the plastic deformation of previous deposited layers. This in-situ peening particle effect is related to particle deposition efficiency which depends on particle size distribution and the processing parameters. CS of coarse powders at relative low propelling gas pressure and temperature can lead to lower particle deposition efficiency, but stronger in-situ peening effect. As a consequence, the coating microstructure fabricated at C1 demonstrates thinnest deposition layer with the largest deformed FeSiAl particles (Fig. 4 (a)).
5. Conclusions
Acknowledgements The author, Xinliang XIE thanks the financial support from the program of China Scholarship Council. The authors would like to thank Dr. Long HOU from Shanghai University for EBSD characterization. References [1] J. Bas, J. Calero, M. Dougan, Sintered soft magnetic materials. Properties and applications, J. Magn. Magn Mater. 254 (2003) 391e398. [2] Z. Zhang, W. Xu, T. Guo, Y. Jiang, M. Yan, Effect of processing parameters on the magnetic properties and microstructures of molybdenum permalloy compacts made by powder metallurgy, J. Alloys Compd. 594 (2014) 153e157. [3] J. Li, X. Peng, Y. Yang, H. Ge, Preparation and characterization of MnZn/FeSiAl soft magnetic composites, J. Magn. Magn Mater. 426 (2017) 132e136. starsi [4] A. Hamler, V. Gori can, B. Su c, A. Sirc, The use of soft magnetic composite materials in synchronous electric motor, J. Magn. Magn Mater. 304 (2006) e816ee819. [5] K.J. Sunday, M.L. Taheri, Soft magnetic composites: recent advancements in the technology, Met. Powder Rep. 72 (6) (2017) 425e429. [6] D. Liu, C. Wu, M. Yan, J. Wang, Correlating the microstructure, growth mechanism and magnetic properties of FeSiAl soft magnetic composites fabricated via HNO3 oxidation, Acta Mater. 146 (2018) 294e303. r, M. Stre , R. Bures, J. Kova , [7] M. Lauda, J. Füzer, P. Kolla ckova c, M. Bat kova I. Ba t ko, Magnetic properties and loss separation in FeSi/MnZnFe 2 O 4 soft magnetic composites, J. Magn. Magn Mater. 411 (2016) 12e17. [8] W. Xu, C. Wu, M. Yan, Preparation of FeeSieNi soft magnetic composites with excellent high-frequency properties, J. Magn. Magn Mater. 381 (2015) 116e119. [9] X. Jin, Q. Wang, W. Khan, Y. Li, Z.H. Tang, FeSiAl/(Ni0. 5Zn0. 5) Fe2O4 magnetic sheet composite with tunable electromagnetic properties for enhancing magnetic field coupling efficiency, J Alloys Compd. 729 (2017) 277e284. r, D. Oleksa kov [10] J. Füzer, P. Kolla a, S. Roth, AC magnetic properties of the bulk FeeNi and FeeNieMo soft magnetic alloys prepared by warm compaction, J. Alloys Compd. 483 (2009) 557e559. [11] C. Wu, H. Chen, H. Lv, M. Yan, Interplay of crystallization, stress relaxation and magnetic properties for FeCuNbSiB soft magnetic composites, J. Alloys Compd. 673 (2016) 278e282. [12] M. Huang, C. Wu, Y. Jiang, M. Yan, Evolution of phosphate coatings during high-temperature annealing and its influence on the Fe and FeSiAl soft magnetic composites, J. Alloys Compd. 644 (2015) 124e130. [13] C. Wu, X. Gao, G. Zhao, Y. Jiang, M. Yan, Two growth mechanisms in one-step fabrication of the oxide matrix for FeSiAl soft magnetic composites, J. Magn. Magn Mater. 452 (2018) 114e119. [14] R. Bai, Z. Zhu, H. Zhao, S. Mao, Q. Zhong, The percolation effect and optimization of soft magnetic properties of FeSiAl magnetic powder cores, J. Magn. Magn Mater. 433 (2017) 285e291. [15] T. Gheiratmand, H.M. Hosseini, S.S. Reihani, Iron-borosilicate soft magnetic composites: the correlation between processing parameters and magnetic properties for high frequency applications, J. Magn. Magn Mater. 429 (2017) 241e250. [16] M. Cherigui, N. Fenineche, G. Ji, T. Grosdidier, C. Coddet, Microstructure and magnetic properties of FeeSi-based coatings produced by HVOF thermal spraying process, J Alloys Compd. 427 (2007) 281e290. [17] M. Cherigui, N. Fenineche, C. Coddet, Structural study of iron-based microstructured and nanostructured powders sprayed by HVOF thermal spraying, Surf. Coating. Technol. 192 (2005) 19e26. [18] N. Fenineche, M. Cherigui, H. Aourag, C. Coddet, Structure and magnetic properties study of iron-based thermally sprayed alloys, Mater. Lett. 58 (2004) 1797e1801. [19] A. Papyrin, Cold spray technology, Adv. Mater. Process. 159 (2001) 49e51. [20] T. Stoltenhoff, H. Kreye, H. Richter, An analysis of the cold spray process and 0
0
In this work, SMCs coating was fabricated through a novel CS technique via Ni-coated FeSiAl composite powder for soft magnetic applications. Two groups of Ni-coated FeSiAl composite powders were used as feedstocks in the CS process with different processing parameters. The phase analysis exhibited no phase transformations in the cold sprayed SMCs coatings due to the low processing temperature. The coating thickness increased with the increase of propelling gas temperature and pressure, while relatively thicker coatings were obtained by the finer powders. A highest retainability of FeSiAl particles with about 50% was observed in the cold sprayed coating with the highest processing parameters, indicating the loss of FeSiAl particles due to their detachment from Ni layer during deposition. Higher deformation of FeSiAl particles and higher micro-hardness of composite were obtained by powders with larger size due to the enhanced peening effect of the rebounded large particles during deposition. EBSD analysis of the composite coating showed that high strain stress retained within the FeSiAl particles while large grain refinement was observed in the severe deformed Ni layers. The magnetic property test showed that the cold sprayed Ni/FeSiAl composite coatings had a ferromagnetic characteristic with a coercivity of ~60 Oe. The coating formation mechanism was discussed by accounting the coating microstructure and the morphologies of the single composite particles. The Ni bonding layer with sufficient plastic deformation was found to play a crucial role during the deposition of composite coating. Our future work will focus on the improvement of the magnetic properties of this composite coating via post heat treatment.
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nanoindentation studies, Scripta Mater. 62 (2010) 395e398. [34] C. Li, W. Li, Y. Wang, H. Fukanuma, Effect of spray angle on deposition characteristics in cold spraying, in: Proceedings of the International Thermal Spray Conference, ASM Int., Materials Park, OH, USA, Orlando, 2003, pp. 5e8. [35] C. Chen, Y. Xie, S. Yin, M.-P. Planche, S. Deng, R. Lupoi, H. Liao, Evaluation of the interfacial bonding between particles and substrate in angular cold spray, Mater. Lett. 173 (2016) 76e79. [36] M. Yu, X. Suo, W. Li, Y. Wang, H. Liao, Microstructure, mechanical property and wear performance of cold sprayed Al5056/SiCp composite coatings: effect of reinforcement content, Appl. Surf. Sci. 289 (2014) 188e196. [37] S. Yin, E.J. Ekoi, T.L. Lupton, D.P. Dowling, R. Lupoi, Cold spraying of WC-Co-Ni coatings using porous WC-17Co powders: formation mechanism, microstructure characterization and tribological performance, Mater. Des. 126 (2017) 305e313. [38] X.-T. Luo, Y.-K. Wei, Y. Wang, C.-J. Li, Microstructure and mechanical property of Ti and Ti6Al4V prepared by an in-situ shot peening assisted cold spraying, Mater. Des. 85 (2015) 527e533. [39] Y.-K. Wei, X.-T. Luo, C.-X. Li, C.-J. Li, Optimization of in-situ shot-peeningassisted cold spraying parameters for full corrosion protection of Mg alloy by fully Dense Al-Based alloy coating, J. Therm. Spray Technol. 26 (2017) 173e183. [40] G. Herzer, Nanocrystalline soft magnetic alloys, in: Handbook of Magnetic Materials, vol. 10, 1997, pp. 415e462. [41] T. Ma, M. Yan, W. Wang, The evolution of microstructure and magnetic properties of FeeSieAl powders prepared through melt-spinning, Scripta Mater. 58 (2008) 243e246. [42] G. Bae, Y. Xiong, S. Kumar, K. Kang, C. Lee, General aspects of interface bonding in kinetic sprayed coatings, Acta Mater. 56 (2008) 4858e4868. [43] X.-T. Luo, C.-J. Li, Large sized cubic BN reinforced nanocomposite with improved abrasive wear resistance deposited by cold spray, Mater. Des. 83 (2015) 249e256. [44] W.-Y. Li, C. Zhang, H. Liao, J. Li, C. Coddet, Characterizations of cold-sprayed NickeleAlumina composite coating with relatively large Nickel-coated Alumina powder, Surf. Coating. Technol. 202 (2008) 4855e4860. [45] S. Yin, Y. Xie, J. Cizek, E.J. Ekoi, T. Hussain, D.P. Dowling, R. Lupoi, Advanced diamond-reinforced metal matrix composites via cold spray: properties and deposition mechanism, Composites Part B 113 (2017) 44e54.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
A novel approach for fabricating Ni-coated FeSiAl soft magnetic composite via cold spraying Xinliang Xie a, Chaoyue Chen a, *, Yingchun Xie a, b, Zhongming Ren c, Eric Aubry d, Gang Ji e, Hanlin Liao a ICB UMR 6303, CNRS, Univ. Bourgogne Franche-Comt e, UTBM, F-90010 Belfort, France National Engineering Laboratory for Modern Materials Surface Engineering Technology, The Key Lab of Guangdong for Modern Surface Engineering Technology, Guangdong Institute of New Materials, Guangzhou 510651, PR China c Shanghai University & State Key Laboratory of Advanced Special Steel, 149 Yanchang Road, Shanghai 200072, PR China d Nipson Technology, 12 Avenue des Trois Ch^ enes, 90000 Belfort, France e Unit e Mat eriaux et Transformations, CNRS UMR 8207, Universit e Lille 1, Villeneuve d'Ascq, 59655, France a
b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 January 2018 Received in revised form 23 March 2018 Accepted 24 March 2018 Available online 27 March 2018
Currently, soft magnetic composites (SMCs) have attracted increasing attention due to their outstanding magnetic properties, and various methods have been developed and applied for their fabrication. As an emerging additive manufacturing technique, cold spraying (CS) can fabricate bulk material via solid-state deposition by avoiding oxidation and phase change. In this work, SMCs coating was first fabricated by Nicoated FeSiAl composite powder via CS. Two groups of Ni-coated FeSiAl composite particles (40 mm and 57 mm) were used as feedstocks for deposition under different processing parameters. No phase transformation can be detected from XRD analysis. The coating thickness increased with the increasing of propelling gas temperature and pressure. Higher deformation of FeSiAl particles and higher microhardness of the composites fabricated from the powders with larger size were obtained due to the enhanced in-situ peening effect of the rebounded particles during deposition. The magnetic property of cold sprayed SMCs showed a soft ferromagnetic characteristic with a coercivity of about 60 Oe. The investigation on the coating formation mechanism was carried out by single particle deposition, and the results showed that the Ni bonding layer with sufficient plastic deformation plays a significant role during the deposition of the composite coating. © 2018 Elsevier B.V. All rights reserved.
Keywords: Cold spraying Soft magnetic composites (SMCs) FeSiAl alloy Cladded powder
1. Introduction Due to their unique magnetic properties such as high permeability and low magnetic coercivity, the soft magnetic materials have been widely used in electrical and electronic devices [1,2]. Generally, the soft magnetic materials can be classified into soft magnetic alloys, soft magnetic oxides and soft magnetic composites (SMCs) [3]. Based on features like low core loss, higher saturation magnetization and electrical resistivity, SMCs have more competitive advantages in various applications and are attracting more and more attention [4e6]. For the moment, various material systems of SMCs have been studied in the literature, such as FeeSi [7], Fe-Ni [8], Fe-Si-Al [3,6,9], Fe-Ni-Mo [10] and Fe-Cu-Nb-Si-B [11].
* Corresponding author. E-mail address: [email protected] (C. Chen). https://doi.org/10.1016/j.jallcom.2018.03.306 0925-8388/© 2018 Elsevier B.V. All rights reserved.
Theoretically, ideal SMCs should not only have high saturation flux density but also stable permeability. Among the above-mentioned systems, FeSiAl-based SMCs have become attractive candidates in various electronic devices with their outstanding soft magnetic properties [9,12,13]. Up to now, various methods and techniques have been developed and used to fabricate the desired SMCs. Traditionally, the most common technique for the production of SMCs is the powder metallurgy [11,14]. In addition, different methods have been applied based on powder compaction following sintering or annealing treatments [15]. Recently, efforts have been made to fabricate FeSi based SMCs coatings by thermal spray technologies such as high velocity oxygen fuel (HVOF) [16,17] and atmospheric plasma spray (APS) [18]. In such thermal spray processes, the molten particles are deposited onto substrate through the acceleration of high-energy sources. The inevitable melting caused by high-energy source can lead to unexpected phase change and
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further affect the magnetic properties. Besides, the unavoidable high porosity and serious oxidation can also seriously weaken the magnetic performance of the thermal sprayed SMCs coatings. As an emerging technique, cold spraying (CS) provides an effective approach to produce high quality coatings of metallic, metal matrix composite even ceramic [19e22]. In this process, the feedstock particles are accelerated by supersonic carrying gas through a Laval nozzle, and then the coating is formed through the intensive plastic deformation of the deposited particles on the substrate. Owing to the unique cold feature, the solid-state particles undergo almost no melting during CS process. Unlike the melting and solidification process in conventional thermal spraying, the cold sprayed coating is formed through the solid-state consolidation via mechanical interlocking and metallurgical bonding [23,24]. Therefore, the cold sprayed coating can effectively avoid the serious oxidation, phase transformation and thermal residual stress that are commonly found in traditional thermal sprayed samples. As a result, the aforementioned benefits enable the CS technique to produce various high-quality coatings without the size limitation. Currently, CS technique has been widely considered as an effective method for additive manufacturing and damage repair [25,26]. Up to now, very few attempts have been carried out to fabricate soft magnetic coatings via CS. M. Cherigui et al. [27] fabricated FeSiBNbCu/Al composite coatings via CS of nanostructured powders mixed with Al powders. The as-sprayed coatings showed a soft ferromagnetic property despite the presence of Al, which was considered as a nonmagnetic material [27]. However, the addition of a nonmagnetic Al layer greatly reduced the overall magnetic permeability, which yielded a problematic balance of properties [5]. According to the reports by W.Y Li et al. [28], nanostructured Fe-Si coatings were deposited by using mechanically milled powders as the feedstocks. The fabricated Fe-Si coating exhibited a high coercivity (190 kA/m) and had the potential for the application to magnetic recording materials [28]. However, due to the introduction of oxygen and nanostructure to the feedstocks during mechanical alloying, the coercivity of the nanostructured Fe-Si coating was three to four times higher in magnitude than the maghemite iron oxide (20-30 kA/m) [28]. Thus, it should be noted that the choice of a proper feedstock for CS deposition of SMCs is of significant importance to the soft magnetic properties. In this study, the Ni/FeSiAl soft magnetic composite was fabricated by the CS technology. Because of their low ductility, FeSiAl particles are difficult to be directly deposited via CS. In order to improve the deposition efficiency of the hard FeSiAl particles, a novel composite powder where FeSiAl core is decorated by thin Ni layer was used as feedstock to fabricate SMCs coatings. Ni was selected for the matrix phase because of its soft magnetic property and its excellent ductility. In addition, Ni shows a relatively high resistance to corrosion and wear. The aim of this work is to examine the effect of powder size distribution range and processing parameters on the microstructure, microhardness and magnetic properties of SMCs coatings. The coating microstructure, particle deformation and retainability of the FeSiAl particles were characterized and estimated. The soft magnetic properties of the assprayed SMCs coatings were evaluated. At last, the deposition mechanism of the composite powder was discussed.
as powder 1 (35-87 mm) and powder 2 (28-58 mm), were used as feedstocks in this study. The detailed size distributions of these two feedstocks measured by a laser diffraction sizer are given in Fig. 1. The surface morphologies of the feedstock powders given in Fig. 2 exhibits a near spherical shape. As the cross-sectional morphologies given in Fig. 2 (b) and (d), the composite particles exhibit a core-shell structure. A thin Ni layer (about 3e5 mm) is uniformly decorated around a FeSiAl particle. As for some smaller FeSiAl particles, they are aggregated and covered with a thicker Ni layer varying from 6 to 14 mm. The average volume fractions of FeSiAl particle in the core-shell structured composite powder 1 and powder 2 were estimated as 62% and 60%, respectively. 2.2. Coating fabrication FeSiAl/Ni composite coatings and splats were produced by using a homemade CS system (LERMPS, UTBM, France) with an optimal de-Laval-type converging-diverging nozzle. High pressure compressed air was used as powder carrier gas and propelling gas. The powders were deposited onto stainless steel substrate under different propelling gas pressures and temperatures, which are listed in Table 1. Nozzle standoff distance and traverse speed were set as 30 mm and 100 mm/s, respectively. Nozzle trajectory was repeated for 20 times to obtain a thick coating. The spacing between adjacent nozzle passes was set as 2 mm. In addition, for understanding the coating formation mechanism, a single particle deposition onto a stainless steel substrate was also conducted at the nozzle traverse speed of 500 mm/s. The stainless steel substrates were polished or grit blasted in order to get individual splat deposition and full coating deposition. 2.3. Materials characterization In order to examine the phase transformation during CS process, the as-sprayed coatings as well as the feedstock powders were examined by an X-ray diffractometer (Siemens D500, Germany) with the Co (l ¼ 1.78897 Å) source at a current of 40 mA, voltage of 35 kV and scan step of 0.02 . A scanning electron microscopy (SEM JEOL-5900LV, Japan) was used to characterize the morphology and microstructure of the powders and the coatings. Prior to electron microscopy, the as-sprayed coating samples were prepared using standard metallographic procedures with the final polishing
2. Experimental procedure 2.1. Starting materials Commercial gas-atomized Fe-9.6%Si-5.4%Al particles were precoated with a Ni layer by hydrometallurgy process (Beijing General Research Institute of Nonferrous Metals, China). Two types of Ni-coated FeSiAl powders with different size distributions, referred
Fig. 1. Particle size distributions of two feedstock powders.
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Fig. 2. SEM micrographs of Ni/FeSiAl composite powders depicting top view and as-polished cross-section: (a) and (b) powder 1, (c) and (d) powder 2.
Table 1 Propelling gas pressure and gas temperature for deposition conditions 1 to 4. Deposition condition
Pressure (MPa)
Temperature ( C)
C1 C2 C3 C4
2.6 2.6 3.0 3.0
550 620 620 650
applied by 0.05 mm Al2O3 solution. The volume fraction analysis of FeSiAl particle in powders and coatings were performed on the cross-sectional micrographs with the 300 and 500 magnification using the imaging software (Image J), respectively. For each sample, five locations were selected from the cross-section of polished coating and the measured data were then averaged. The flattening ratio of FeSiAl particle was obtained by measuring the longest length of the particle over the shortest length of the particle. For flattening ratio measurement of the FeSiAl particles, an average value was obtained using a minimum of 40 particles per sample. Electron back scattered diffraction (EBSD) scans were recorded at electron energy of 25 KeV, and with a step size of 100 nm using the FEI-Sirion SEM equipped with a TSL-OIM EBSD camera. EBSD samples were sectioned from the CS deposition and prepared by standard metallographic techniques, and the final polishing was conducted using a 0.05 mm diameter colloidal silica suspension and vibratory polishing. The coating microhardness was measured by a Vickers hardness indenter (Leitz, Germany) with a load of 100 g for 10 s. 10 positions were randomly tested and averaged to determine the microhardness. The magnetic measurements for the powders and the SMCs coatings were carried out with a hysteresismeter Bull M 2000/2010 SIIS apparatus (Niposon technology) at ambient temperature. In order to prepare the samples for magnetic property test, SMCs coatings were deposited onto grit blasted Al plate with the propelling gas pressure and temperature of 3.0 MPa and 650 C (C4), respectively. Then, the 10 mm 10 mm cut samples were placed in
high-concentration sodium hydroxide solution for several hours to fully dissolve the Al substrate; while the SMCs coating can be well retained during corrosion. Before the magnetic test, the composite coatings with a thickness of 300e400 mm were grounded to remove the surface oxide films.
3. Results 3.1. Phase composition Fig. 3(a) shows the XRD patterns of the initial powder and CS coating fabricated at C4. Besides the main phases of Ni and Al0.3Si0.7Fe3, no other potential phases appear on the XRD patterns of cold sprayed SMCs coatings, indicating that neither significant chemical reaction nor phase transformation occurred during deposition. This is typically attributed to the low temperature characteristic of the CS process. One can observe that the (111) peek of Ni is much stronger than the (110) peek of Al0.3Si0.7Fe3 phase even though the volume fraction of Ni (40 vol%) is lower than that of Al0.3Si0.7Fe3 (60 vol%) in the powder. This is due to the depth limitation of X-ray diffraction and the core-shell structure of the feedstock powder, where most of the X-rays are diffracted by the Ni coating layers (several microns). Fig. 3(b) shows the detailed XRD spectra of the 2q angle between 50 and 55 . It can be observed that the Ni diffraction peaks of the coatings slightly shifted to the right in comparison with that of the initial powder. The Al0.3Si0.7Fe3 (110) peeks show the same trend despite their low intensities. These results indicate that the lattice parameters of both the Ni and Al0.3Si0.7Fe3 phases are somewhat smaller in comparison with those of the feedstock. This can be explained by the severe plastic deformation of particles during the high-velocity impact in CS [29], and further investigation will be carried out by EBSD characterization in the following sections.
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Fig. 3. X-ray diffraction pattern of cold sprayed SMCs coatings fabricated at C4 and feedstock powder.
3.2. Microstructure Fig. 4 shows the cross-sectional images of the Ni/FeSiAl SMCs coatings fabricated by different feedstock powders. At the coating/ substrate interface, the intimate bonding can be observed without obvious cracks. The coatings exhibit high densification rate and crack-free features. As indicated by the red arrows, a few microsized pores can be observed in the cross-sectional images, which have the similar shapes with the remaining FeSiAl particles. Such micro-sized pores should be attributed to the detachment of FeSiAl particles during the sample preparation rather than the CS process. Statistical analysis of the cross-sectional SEM images at 500 yields a porosity less than 1% for all the as-sprayed coatings, where no evident inter-particle cracks or pores can be observed. As shown in Fig. 4 (a), a thinner SMCs coating with high surface roughness even serious delamination can be seen at low processing parameters. It can be explained by the strong peening and rebounding of the unsuccessful deposition particles due to the low deposition efficiency. From C2 to C4, the improved deposition efficiency with the increasing propelling gas pressure and temperature significantly increases the coating thickness (see Fig. 4 (b) to (d)). It suggests that higher propelling gas pressure and temperature are beneficial for the deposition of composite particle. As for the SMCs coatings produced by the starting powder with a different size distribution, similar tendency of coating thickness can be observed from C1 to C4. The variation of coating thickness with the processing parameters of these two types of feedstock
powders is given in Fig. 5. While comparing the SMCs coatings at the same processing parameters, a thicker coating was obtained from the finer powder, indicating a much higher particle deposition efficiency. According to the deposition window in terms of particle size given by Schmidt et al., the particle obtained a higher in-flight velocity at a smaller particle size within a certain range [30]. Thus, a higher impact velocity of composite particle contributes to a thicker
Fig. 5. Variation of coating thickness with processing parameters.
Fig. 4. Cross sections of the SMCs coatings fabricated from powder 1 (aed) and powder 2 (eeh) under different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
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coating for the finer starting powder. It is commonly considered that the uniform distribution of the reinforced particles in the composite coating is the key to the desired coating properties. As shown in Fig. 4 (a) and (e), the FeSiAl particles were found concentrated on the top of the coating, while the bottom mainly consists of Ni layers. In the cases from C2 to C4, the FeSiAl particles can be observed uniformly distributed in the SMCs coatings as the increase of processing parameters. Besides, the volume fraction of FeSiAl particles was measured to evaluate the FeSiAl content in the SMCs coatings, and the results are given in Table 2. The volume fraction of FeSiAl particles in all the coatings were calculated to be less than 30%, which is half of the value in initial powder (~60 vol%). This indicates that many FeSiAl particles were lost during the coating build-up process. As the propelling gas pressure and temperature decrease, more FeSiAl particles rebound away from the coating surface, which results in less FeSiAl content in the coatings. For each condition, the average volume fraction of FeSiAl in the coatings obtained from powder 2 is slightly larger than that from powder 1. In addition, it can be noticed that only a few of large FeSiAl particles retain in the coatings. The results demonstrate that finer particles possess a higher deposition efficiency (see Table 3). Fig. 6 shows the magnified views of the cross-sectional morphologies of the cold sprayed coatings. The FeSiAl particles are intimately bonded with the ductile Ni binder phase. As shown in Fig. 6 (a) and (e) with relatively low processing parameters, the FeSiAl particles are highly deformed and elongated at the direction perpendicular to deposition. In the case of C1, the flattening ratio of FeSiAl particles were measured to be 2.52, showing an increment about 131.2% in comparison with the initial powder (flattening ratio ¼ 1.14). As the propelling gas pressure and temperature increase from C1 to C4, the flattening ratio of FeSiAl particle gradually reduced from 2.52 to 1.68, indicating lower particle deformation at a higher propelling gas pressure and temperature. Generally, the deformation of cold sprayed particles mainly depends on two major facts: the impact of deposited particle and the peening or tamping effect of subsequent particle impacts [31]. High processing parameters can lead to higher particle impact velocity, which is beneficial for particle deformation and can result in higher deposition efficiency. However, much higher particle deformation has been observed at C1 with relatively low processing parameters. The extremely low deposition efficiency of FeSiAl particles at such low processing parameters leads to the enhanced peening effects during deposition. Many FeSiAl particles rebounded from the coating surface owing to unsuccessful bonding. These particles were acting as in-situ shot peening particles, inducing hammering effect on the previous deposited layers. Thus, more prominent deformation of FeSiAl particles can be found in the cases of low processing parameters. Meanwhile, in the view of starting particle sizes, the average flattening ratio of FeSiAl particle in the coating obtained from powder 1 is larger than that from powder 2 under the same conditions. Because of the lower deposition efficiency and larger kinetic energy of the coarser particles, the enhanced in-situ peening effect directly leads to the severe deformation of FeSiAl particles during deposition. EBSD characterization was applied to further investigate the particle deformation behavior and microstructure of cold sprayed
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coatings. Fig. 7 shows the EBSD result of the SMCs coating obtained from powder 1 at C4. As it can be observed from Fig. 8 (a), the elongated FeSiAl particles are in large grain size, and they are surrounded by a plenty of randomly oriented ultrafine Ni grains. The distribution of Ni grain size is given in Fig. 7 (b) and its average grain size is measured to be 0.57 mm. Grain boundary distribution in Fig. 7 (c) shows that most of the Ni grains are in large grain boundaries, as described in the red lines, while small grain boundaries (blue lines) or even some subgrains can be observed within the FeSiAl particles. The formation of these nanometer-sized Ni grains is due to the dynamic recrystallization during the severe plastic deformation and thermal softening process [32,33]. Fig. 7 (d) is a kernel average misorientation (KAM) map, which is a measure of plastic strain. Higher plastic strain can be observed within the FeSiAl particles rather than that in Ni layers. Such phenomena demonstrate that considerable strain stress was generated and concentrated within the FeSiAl particles during the high-velocity particle impact. The crystal orientation evolution within a FeSiAl grain is shown in Fig. 8. It can be seen that the lattice is gradually rotated along the line ABCD and their orientation evolution is illustrated in the IPF image (Fig. 8 (b)). The results indicate that deformation induced strain stress may lead to the formation of subgrains or small grain boundaries in the hard FeSiAl particles.
3.3. Surface morphology Fig. 9 displays the top-view surface morphologies of the coatings obtained under different conditions. A common feature is the presence of craters for all the coatings. As for the coatings obtained from powder 1 (see Fig. 9(aed)), relatively flat surface with small craters can be observed. However, the coating surfaces obtained from powder 2 present several large craters (about 200 mm) (see Fig. 9(eeh)). The formation of such craters could be attributed to the erosion effect of the high velocity particles during deposition, especially for the coatings deposited from the finer powders. Once small craters were formed, particles became more and more difficult to deposit on the wall of the craters, due to the fact that the deposition efficiency decreases with the angle of the projection [34,35]. As a result, large craters were gradually formed with the growth in coating thickness. In addition, many traces of the shapes of their corresponding powder particles are observed on the coating surfaces. The magnified images in Fig. 10 show the representative surface morphologies in four different cases. In the first case, the entire particles rebounded away with only traces of particles remaining on a highly deformed surface. In the second case (see Fig. 10 (b)), a FeSiAl particle as the core part bounced off leaving a highly deformed Ni layer on the coating surface. In the third case, a successful deposit particle was eroded by the incoming particles with the upper Ni layer being fractured. These exposed FeSiAl particles will prevent the subsequent particles from deposition. In the final case, particles were successful bonded with previous deposited layers. These deposition morphologies allow us to understand the coating formation mechanisms, which will be discussed in the following section.
Table 2 Volume fraction of FeSiAl particles in the SMCs coatings. Volume fraction
Initial powder
C1
C2
C3
C4
Powder 1 (d0.5 ¼ 56 mm) Powder 2 (d0.5 ¼ 40 mm)
62.0 ± 1.5% 60.0 ± 1.2%
19.1 ± 0.9% 22.1 ± 1.8%
23.7 ± 2.4% 24.3 ± 1.6%
25.4 ± 2.0% 27.4 ± 1.5%
28.6 ± 2.3% 29.8 ± 2.8%
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Table 3 Flattening ratio variation of the FeSiAl particles under different processing parameters. Flattening ratio
Initial powder
C1
C2
C3
C4
Powder 1 (d0.5 ¼ 56 mm) Powder 2 (d0.5 ¼ 40 mm)
1.14 ± 0.12 1.18 ± 0.16
2.52 ± 0.96 (131.2%) 1.96 ± 0.71 (79.8%)
1.86 ± 0.55 (68.8%) 1.73 ± 0.58 (58.7%)
1.82 ± 0.50 (67.9%) 1.68 ± 0.59 (57.8%)
1.68 ± 0.51 (54.0%) 1.56 ± 0.44 (43.1%)
Fig. 6. Magnified views of the cross-sectional morphology of cold spray coatings fabricated from powder 1 (aed) and powder 2 (eeh) with different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 7. EBSD characterization of the cross-section of the as-sprayed coating fabricated at C4: (a) IPFþIQ map, (b) Ni grain size distribution, (c) Grain boundary distribution, (e) Kernel average misorientation (KAM) map.
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3.4. Microhardness
Fig. 8. A close view of a FeSiAl particle in Fig. 7: (a) Euler angle map with four points (A, B, C and D) marked in a line from left to right, (b) Inverse pole figure of the points in the line ABCD.
Fig. 11 shows the microhardness of SMCs coatings fabricated under different conditions. Clearly, higher coating microhardness is obtained from the coarser powder despite the fact that the volume fraction of FeSiAl particles is lower than that in the finer powder. In addition, coatings deposited with the coarser powder exhibite smaller variation in microhardness values compared with that of the finer powder. For the coarser powder, the microhardness value decreases slightly with the increasing of processing parameters, and C1 coating presents the highest value even though it has the lowest FeSiAl volume fraction. The SMCs coatings obtained from the finer powders demonstrate the similar tendency of microhardness as a function of processing parameters. Fig. 12 shows the morphologies of the Vickers indentations on the coating cross-
Fig. 9. Surface morphologies of the cold sprayed SMCs coatings fabricated from powder 1 (aed) and powder 2 (eeh) at different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 10. Typical surface morphologies of the Ni/FeSiAl SMCs coating in magnified views.
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particles. At low processing parameters, the low deposition efficiency leads to the unsuccessful bonding of the FeSiAl particles. The rebounding of FeSiAl particles can result in the enhanced peening effect. Due to the in-situ hammering by these shot peening particles, plastic deformation of previously deposited layers could be greatly enhanced and thereby the improvement of interparticle bonding [38,39]. In addition, the induced plastic deformation via the hammering effect can also lead to dislocation accumulation and grain refinement for the previous deposited layers [38], which in turn results in the improvement of coating hardness and toughness. 3.5. Magnetic property
Fig. 11. Microhardness of SMCs coatings under different conditions.
sections. The coatings fabricated from powder 1 present a smaller indentation size and the smallest is shown in C1. Cracks are found in the coatings after indentation test, especially for the coatings fabricated from the fine powder. These results indicate that an improved coating toughness can be obtained from the coarser powder due to the enhanced peening effect during deposition. Generally, it is considered that a higher content of reinforced particles can significantly improve the microhardness of composite coatings [36,37]. However, the highest value of microhardness was obtained in the coating with the lowest volume fraction of FeSiAl
Fig. 13 shows the hysteresis loops of the Ni/FeSiAl SMCs coatings. Their corresponding coercivity and saturation magnetization values are given in Table 4. These coatings obtained from powder 1 and powder 2 exhibit the similar soft magnetic properties with the coercivity of 61.5 Oe and 62.2 Oe, respectively, which have a slight increment compared to that of the feedstock powder (48 Oe); while the saturation magnetizations of the coatings exhibit a slight decrease compared with that of the feedstock powder. The increase in the coercivity value can be explained by the followed reasons: (i) As mentioned above, a large amount of FeSiAl particles was lost during deposition. Consequently, only half volume fraction of FeSiAl particles were retained in the coating; (ii) The generation of strain stress and dislocations in the plastic deformed FeSiAl particles during impact can lead to the increase of coercivity. (iii) It is well known that the crystallite size has a great effect on the soft magnetic performance [40]. The coercivity increases with the crystal size reduction [41]. Even though grain refinement within the FeSiAl particles was not observed, the nano-sized Ni grains
Fig. 12. Morphology of the Vickers indentations on the coating cross-sections fabricated from powder 1 (aed) and powder 2 (eeh) at different processing parameters: (a) and (e) C1, (b) and (f) C2, (c) and (g) C3, (d) and (h) C4.
Fig. 13. (a) Hysteresis loops of the cold sprayed coatings and (b) magnification of the circled area.
X. Xie et al. / Journal of Alloys and Compounds 749 (2018) 523e533 Table 4 Coercivity and saturation magnetization of initial powder and SMCs coatings. Sample
Coercivity Hc (Oe)
Saturation magnetization Ms (emu/g)
Powder 1 Powder 1-C4 Powder 2-C4
48.4 61.5 62.2
80.5 68.2 79.7
derived from dynamic recrystallization may increase the coating coercivity. In the future plan, to further improve the soft magnetic property of the composite coating, post heat treatment on the assprayed coatings will be carried out to relax strain stress and enlarge grain size.
4. Discussion In CS, the successful deposition of a single ductile metallic particle on the substrate or on previous deposited layers is largely attributed to the occurrence of adiabatic shear instability [42]. This is a phenomenon that occurs in the particle when it undergoes high strain rate deformation upon impact. It is known that hard reinforcements (e.g. BN [43], Al2O3 [44], diamond [45], WC-Co [37], SiC [36]) during the CS deposition of metal matrix composite coatings do not undergo plastic deformation. Due to the lack of sufficient plastic deformation, these hard particles tend to rebound away from the coating surface rather than successful bonding. Similarly, it is also found from the present work that the volume fraction of FeSiAl particle was decreased compared with the starting feedstocks, which indicates their rebounding during deposition. Meanwhile, the hard FeSiAl core experienced slight deformation and no metallurgical bonding occurred with the soft Ni layer (Figs. 4 and 7). To further understand the influencing factors of magnetic and mechanical properties, the formation mechanism of cold sprayed SMCs coatings is discussed through the single composite particle deposition. Fig. 14 displays the surface and crosssectional morphologies of the single Ni/FeSiAl particle onto SS substrate. Four typical morphologies are given indicating the successful bonding, particle rebounding and complete rebounding of a single composite particle. In the first case (Fig. 14 (a) and (e)), the successful deposition of a composite particle can be seen on the SS substrate with strong deformation of the Ni layer. In Fig. 14 (e), the metal jet can be observed at the rim of the deformed particle, which gives a direct proof for the metallurgical bonding between the composite particle and substrate. During the impact of the composite particle, the ductile Ni layer experienced severe plastic
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deformation and thermal softening. As a consequence, in this region as adiabatic shear instability occurred, Ni layer behaves like a viscous metal, which results in an extrusive outward metal jet at the rim. As shown in Fig. 14 (b) and (f), the Ni layer around FeSiAl was fractured with large cracks, and a rim of metal jet can be found around the deformed particle. Different from the deposition behavior of ductile Ni layer, more elastic nature of FeSiAl particle can be expected. Since the kinetic energy of FeSiAl particle was partly transformed into the plastic deformation, the remaining was stored as elastic energy that led to the rebounding of FeSiAl particle. Thus, the observed crack on Ni layer can be formed as the rebounding energy exceeding its strength limitation. As seen in Fig. 14 (c) and (g), after the rebounding of FeSiAl, part of Ni layer still remains in the substrate crater owing to the existence of metallurgical bonding between Ni and substrate. Due to the high kinetic energy and insufficient plastic, the rebounding of FeSiAl particle caused the complete fracture of Ni bonding layer in the composite powder. In the last case, due to the absence of metallic bonding with the substrate, the composite particle experienced a complete rebound off the substrate surface. As it can be seen in Fig. 14 (d), the substrate experienced a high degree of deformation, leaving deep indentations of the rebounded particle. It indicates that this thin layer of Ni sometimes cannot achieve metallic bonding with the substrate due to the insufficient particle deformation, for example, the already deposited particle can be removed from the substrate by the subsequent particles. Based on the deposition behavior of the single composite particle, the formation mechanism of the composite coating in this work is investigated and discussed. Fig. 15 shows the schematic of the deposition process for Ni/FeSiAl composite coating. It can be seen that the composite particle with a shell-core structure experienced three stages during coating build-up. As the composite particle impacts onto the previous deposited Ni layers with sufficient thickness in the first stage, the metallurgical bonding occurs between the Ni bonding layers of composite particles. The sufficient plastic deformation can effectively remove the broken oxide-films on Ni layer interface and promote the successful deposition process by introducing metallurgical bonding. In the second and last stages, FeSiAl core or the entire Ni/FeSiAl composite particle experience rebounding after impact due to the absence of high bonding strength between the Ni bonding layers. As for the FeSiAl core, the high kinetic energy as a result of high velocity is mostly stored as elastic energy rather than dissipated in the forms of plastic deformation. Thus, the rebounding of FeSiAl particles can be
Fig. 14. Surface and cross-sectional morphologies of a single splat deposited onto SS substrate in different cases, red arrows showing the material jet. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 15. Schematic of the deposition process for Ni-coated FeSiAl coating.
observed as the bonding between Ni layer is weak, which attributes to the decreasing content of FeSiAl in the as-sprayed SMCs coating. The metallurgical bonding during the composite coating built-up process only takes place at the interfaces between the Ni layers or between the Ni layer and substrate. It can be concluded that this soft Ni layer plays a crucial role on the formation process of composite coatings. At the same time, these rebounded particles act as in-situ peening particles which can further enhance the plastic deformation of previous deposited layers. This in-situ peening particle effect is related to particle deposition efficiency which depends on particle size distribution and the processing parameters. CS of coarse powders at relative low propelling gas pressure and temperature can lead to lower particle deposition efficiency, but stronger in-situ peening effect. As a consequence, the coating microstructure fabricated at C1 demonstrates thinnest deposition layer with the largest deformed FeSiAl particles (Fig. 4 (a)).
5. Conclusions
Acknowledgements The author, Xinliang XIE thanks the financial support from the program of China Scholarship Council. The authors would like to thank Dr. Long HOU from Shanghai University for EBSD characterization. References [1] J. Bas, J. Calero, M. Dougan, Sintered soft magnetic materials. Properties and applications, J. Magn. Magn Mater. 254 (2003) 391e398. [2] Z. Zhang, W. Xu, T. Guo, Y. Jiang, M. Yan, Effect of processing parameters on the magnetic properties and microstructures of molybdenum permalloy compacts made by powder metallurgy, J. Alloys Compd. 594 (2014) 153e157. [3] J. Li, X. Peng, Y. Yang, H. Ge, Preparation and characterization of MnZn/FeSiAl soft magnetic composites, J. Magn. Magn Mater. 426 (2017) 132e136. starsi [4] A. Hamler, V. Gori can, B. Su c, A. Sirc, The use of soft magnetic composite materials in synchronous electric motor, J. Magn. Magn Mater. 304 (2006) e816ee819. [5] K.J. Sunday, M.L. Taheri, Soft magnetic composites: recent advancements in the technology, Met. Powder Rep. 72 (6) (2017) 425e429. [6] D. Liu, C. Wu, M. Yan, J. Wang, Correlating the microstructure, growth mechanism and magnetic properties of FeSiAl soft magnetic composites fabricated via HNO3 oxidation, Acta Mater. 146 (2018) 294e303. r, M. Stre , R. Bures, J. Kova , [7] M. Lauda, J. Füzer, P. Kolla ckova c, M. Bat kova I. Ba t ko, Magnetic properties and loss separation in FeSi/MnZnFe 2 O 4 soft magnetic composites, J. Magn. Magn Mater. 411 (2016) 12e17. [8] W. Xu, C. Wu, M. Yan, Preparation of FeeSieNi soft magnetic composites with excellent high-frequency properties, J. Magn. Magn Mater. 381 (2015) 116e119. [9] X. Jin, Q. Wang, W. Khan, Y. Li, Z.H. Tang, FeSiAl/(Ni0. 5Zn0. 5) Fe2O4 magnetic sheet composite with tunable electromagnetic properties for enhancing magnetic field coupling efficiency, J Alloys Compd. 729 (2017) 277e284. r, D. Oleksa kov [10] J. Füzer, P. Kolla a, S. Roth, AC magnetic properties of the bulk FeeNi and FeeNieMo soft magnetic alloys prepared by warm compaction, J. Alloys Compd. 483 (2009) 557e559. [11] C. Wu, H. Chen, H. Lv, M. Yan, Interplay of crystallization, stress relaxation and magnetic properties for FeCuNbSiB soft magnetic composites, J. Alloys Compd. 673 (2016) 278e282. [12] M. Huang, C. Wu, Y. Jiang, M. Yan, Evolution of phosphate coatings during high-temperature annealing and its influence on the Fe and FeSiAl soft magnetic composites, J. Alloys Compd. 644 (2015) 124e130. [13] C. Wu, X. Gao, G. Zhao, Y. Jiang, M. Yan, Two growth mechanisms in one-step fabrication of the oxide matrix for FeSiAl soft magnetic composites, J. Magn. Magn Mater. 452 (2018) 114e119. [14] R. Bai, Z. Zhu, H. Zhao, S. Mao, Q. Zhong, The percolation effect and optimization of soft magnetic properties of FeSiAl magnetic powder cores, J. Magn. Magn Mater. 433 (2017) 285e291. [15] T. Gheiratmand, H.M. Hosseini, S.S. Reihani, Iron-borosilicate soft magnetic composites: the correlation between processing parameters and magnetic properties for high frequency applications, J. Magn. Magn Mater. 429 (2017) 241e250. [16] M. Cherigui, N. Fenineche, G. Ji, T. Grosdidier, C. Coddet, Microstructure and magnetic properties of FeeSi-based coatings produced by HVOF thermal spraying process, J Alloys Compd. 427 (2007) 281e290. [17] M. Cherigui, N. Fenineche, C. Coddet, Structural study of iron-based microstructured and nanostructured powders sprayed by HVOF thermal spraying, Surf. Coating. Technol. 192 (2005) 19e26. [18] N. Fenineche, M. Cherigui, H. Aourag, C. Coddet, Structure and magnetic properties study of iron-based thermally sprayed alloys, Mater. Lett. 58 (2004) 1797e1801. [19] A. Papyrin, Cold spray technology, Adv. Mater. Process. 159 (2001) 49e51. [20] T. Stoltenhoff, H. Kreye, H. Richter, An analysis of the cold spray process and 0
0
In this work, SMCs coating was fabricated through a novel CS technique via Ni-coated FeSiAl composite powder for soft magnetic applications. Two groups of Ni-coated FeSiAl composite powders were used as feedstocks in the CS process with different processing parameters. The phase analysis exhibited no phase transformations in the cold sprayed SMCs coatings due to the low processing temperature. The coating thickness increased with the increase of propelling gas temperature and pressure, while relatively thicker coatings were obtained by the finer powders. A highest retainability of FeSiAl particles with about 50% was observed in the cold sprayed coating with the highest processing parameters, indicating the loss of FeSiAl particles due to their detachment from Ni layer during deposition. Higher deformation of FeSiAl particles and higher micro-hardness of composite were obtained by powders with larger size due to the enhanced peening effect of the rebounded large particles during deposition. EBSD analysis of the composite coating showed that high strain stress retained within the FeSiAl particles while large grain refinement was observed in the severe deformed Ni layers. The magnetic property test showed that the cold sprayed Ni/FeSiAl composite coatings had a ferromagnetic characteristic with a coercivity of ~60 Oe. The coating formation mechanism was discussed by accounting the coating microstructure and the morphologies of the single composite particles. The Ni bonding layer with sufficient plastic deformation was found to play a crucial role during the deposition of composite coating. Our future work will focus on the improvement of the magnetic properties of this composite coating via post heat treatment.
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