Controllable electrochemical synthesis and magnetic behaviors of Mg–Mn–Fe–Co–Ni–Gd alloy films

Journal of Alloys and Compounds 598 (2014) 161–165

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Controllable electrochemical synthesis and magnetic behaviors of Mg–Mn–Fe–Co–Ni–Gd alloy films H. Li a, H. Sun a, C. Wang b, B. Wei a, C. Yao a,⇑, Y. Tong b,⇑, H. Ma a a

Department of Applied Chemistry, Yuncheng University, Yuncheng 044000, People’s Republic of China KLGHEI of Environment and Energy Chemistry, MOE of the Key Laboratory of Bioinorganic and Synthetic Chemistry, School of Chemistry and Chemical Engineering, Sun Yat-Sen University, Guangzhou 510275, People’s Republic of China b

a r t i c l e

i n f o

Article history: Received 9 January 2014 Received in revised form 11 February 2014 Accepted 11 February 2014 Available online 17 February 2014 Keywords: Electrodeposition Multi-component alloy Magnetic properties Mott–Anderson transition

a b s t r a c t Mg–Mn–Fe–Co–Ni–Gd alloy films were prepared by electrodeposition. The surface morphology can be controlled by deposition potential and solution composition. Hollow microspheres and core–shell microspheres can be obtained. The as-deposited alloy is amorphous. The ferromagnetism to diamagnetism transition can be observed when the contents of Mg is lower than 20%. The Anderson–Mott transition mechanism was proposed to explain the transition. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction As known to us, traditional alloys are mainly composed of one or two elements for primary properties. Also, other minor elements (<15%) are often added for acquiring specific microstructure and properties [1]. In recent years, alloy field has been extended to multi-component alloys (MCAs), which may contain at least five principal elements with the concentration of each element being between 35 and 5 at.% in equimolar or near equimolar ratios [2,3]. The MCAs are stable because of their large mixing entropies and the low free energy remarkably, and make random solid solution more stable than the ordered phases. The previous studies have shown that the MCAs might possess simple crystal structures, ease of nanoprecipitation, and promising properties in high hardness and superior resistance to temper softening, wear, oxidation and corrosion [4,5]. Some efforts have been devoted to study the magnetic property of the alloys. Andrzejewski et al. reported that the antiferromagnetic YbFe4Al8 exhibits a minus magnetic susceptibility when the temperature is below 50 K [6], as for their superconductivity, the authors pointed out that there are still some debates. Wang et al. found that the Pr60Al10Ni10Cu16Fe4 alloy shows an abnormal diamagnetic transition while cooling at 20 K in the zero field ⇑ Corresponding authors. Tel.: +86 359 8594394 (C. Yao), Tel.: +86 20 84110071 (Y. Tong). E-mail addresses: [email protected] (C. Yao), [email protected] (Y. Tong). http://dx.doi.org/10.1016/j.jallcom.2014.02.051 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

cooled (ZFC) curves, but the diamagnetic transition was not observed in the FC curves [7]. Our group have reported the magnetic properties of the MCAs based on TM–Ln (TM: transition metal; Ln: lanthanide) [8,9]. Alloys based on the Mg–TM–Ln system have attracted much attention due to their high energy density hydrogen storage capacity, high specific strength, excellent corrosion resistance, good bend ductility, light weight and low cost [10–16]. However, there are few reports on the MCAs based on Mg–TM–Ln. In this paper, we report the electrochemical synthesis of Mg–Mn–Fe–Co–Ni–Gd MCAs from dimethyl sulfoxide (DMSO) solution. The morphology and structure of Mg–Mn–Fe–Co–Ni–Gd alloys were readily controlled by adjusting the corresponding electrochemical parameters. It should be noted that the abnormal compositiondependent magnetic phase transitions were observed.

2. Experimental The electrochemical measurements were carried out on a CHI 750b electrochemical workstation (CHI Instrument Inc., USA). Three-electrode systems were employed, in which a Cu foil (10 mm  40 mm), a Pt foil and a Ag/AgCl electrode was used as working electrode, auxiliary electrode and reference electrode, respectively. The alloy films were deposited on a Cu foil by potentiostatic deposition in DMSO (AR) solution containing GdCl3, FeCl2, CoCl2, NiCl2, MnCl2, and MgCl2. All the salts were dehydrated in vacuum at 373 K. In addition, Vitamin C (as oxygen remover and stabilizer for ferrous ion) and (C2H5)4NPF6 (as supporting electrolyte) were also added in the electrolyte. DMSO–vitamin C(0.02 mol L1)–(C2H5)4NPF6(0.02 mol L1)–GdCl3(0.01 mol L1)–FeCl2(0.005 mol L1)– CoCl2(0.005 mol L1)– NiCl2(0.005 mol L1)–MnCl2(0.005 mol L1)–MgCl2(0.01 mol L1) was defined as

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mixed solution 1, and DMSO–vitamin C(0.02 mol L1)–(C2H5)4NPF6(0.02 mol L1)– GdCl3(0.02 mol L1–FeCl2(0.01 mol L1)–CoCl2(0.01 mol L1)–NiCl2(0.01 mol L1)– MnCl2(0.01 mol L1)–MgCl2(0.02 mol L1) was mixed solution 2. The electrochemical tests and electrodeposition were performed with the protection of Ar atmosphere. The surface morphology and the chemical composition of the alloy thin films were studied by scanning electron microscope (SEM, Quanta 400) equipped with energy dispersive X-ray spectrometer (EDX, OXFOXD-INCA). The structures of deposits were investigated by X-ray diffraction (XRD, RIGAKU, D/MAX 2200 VPC) with Cu Ka radiation (k = 1.5418 Å). The electrochemical properties of alloy thin films were investigated with CHI 750b electrochemical workstation. A superconducting quantum interference device (SQUID, QUANTUM DESIGN, MPMS XL-7) was used to study the magnetic behaviors of these alloy thin films.

3. Results and discussion 3.1. Cyclic voltammetry and I–t curves Fig. 1(a) shows the cyclic voltammograms (C–V) of Cu electrodes in different solutions of DMSO–vitamin C (0.02 mol L1)–(C2H5)4NPF6 (0.02 mol L1)–MClx, in which MClx =

GdCl3, NiCl2, CoCl2, FeCl2, MnCl2 and MgCl2 with a concentration of 0.01 mol L1 respectively. From the CV curves, it can be seen that the reduction peak potentials for Gd(III), Mn(II) and Mg(II) ions are at about 1.9 V, 2.3 V and 2.2 V, while Fe(II), Co(II) and Ni(II) are at 1.6 and 1.8 V, respectively. The six reduction peaks are corresponding to the reduction of these six ions. However, only one clear reduction peak was observed at about 1.5 V when the six metal ions were presented in mixed solution 1. This means that the induced codeposition may occur. Fig. 1(b) shows the current–time curves of samples deposited at 2.0 V in mixed solution 1 and mixed solution 2. The results show that current density in mixed solution 1 quickly dropped to the lowest value of 0.5 mA cm2 within about 60 s and then increased slowly to a nearly constant value of 0.8 mA cm2. However, the current decreases slightly and then keeps a constant value of about 1.3 mA cm2 in mixed solution 2. This indicates a faster electro-crystallization nucleation process and steady growth of the deposits occurred when the salts concentration was relatively

Fig. 1. (a) Cyclic voltammograms of Cu electrodes in different solutions of DMSO–vitamin C(0.02 mol L1)–(C2H5)4NPF6(0.02 mol L1)–MClx, MClx = GdCl3, NiCl2, CoCl2, FeCl2, MnCl2 and MgCl2 with a concentration of 0.01 mol L1 respectively and mixed solution 1: DMSO–GdCl3(0.01 mol L1)–FeCl2(0.005 mol L1)–CoCl2(0.005 mol L1)– NiCl2(0.005 mol L1)–MnCl2(0.005 mol L1)–MgCl2(0.01 mol L1), sweep rate: 50 m Vs1. (b) Current density curves of mixed solution 1 and mixed solution 2 (with double salts concentration of 1) as a function of time in DMSO organic system.

Fig. 2. The typical SEM images of the alloys deposited at 2.0 V for 1 h in mixed solution1 (a and b) and mixed solution 2 (c–f), (d–f) exhibiting the formation process of the open core–shell structures to closed ones.

H. Li et al. / Journal of Alloys and Compounds 598 (2014) 161–165

high. In the view of elctrochemsitry the different morphology and structure of the deposited Mg–Mn–Fe–Co–Ni–Gd alloys can be anticipated. 3.2. Morphology, structure and composition The morphology and microstructure of the as-prepared Mg– Mn–Fe–Co–Ni–Gd alloys were studied by SEM. The typical SEM images are given in Fig. 2. Fig. 2(a) is the SEM image of the sample obtained in mixed solution 1 at 2.0 V for 1 h, a lot of hollow hemispheres with uniform distribution were successfully grown on Cu substrates. The high resolution SEM images in Fig. 2(b) and the inset demonstrate that the diameter of the hollow hemispherical architectures varies from 500 nm to 1000 nm and the thickness of their shell is about 10–30 nm. Fig. 2(c–g) is the SEM images of the sample deposited in mixed solution 2 at 2.0 V for 1 h. The deposit consists of core–shell spheres, as shown in Fig. 2(c). The diameter of these spheres is about 500–1500 nm. Fig. 2(d and e) shows the SEM images of the unsealed or the broken core–shell spheres, clearly indicating the core consists several particles. The shell thickness of the spheres is about 140–190 nm and the average diameter of the particles is about 100 nm. The SEM images in Fig. 2(d–f) exhibit the formation process of the open core–shell structures to closed ones. So the morphology of the sample can be tuned by controlling the concentration of the electrolyte, i.e. the hollow spheres can be obtained in mix solution 1 and core– shell spheres can be obtained in mixed solution 2.

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The structures of hollow and core–shell spheres were studied by XRD, as shown in Fig. 3(a). The XRD patterns were recorded within the scanning range of 2h from 15° to 70°. Besides the diffraction peaks of Cu (1 1 1) and (2 0 0) coming from Cu substrates, no any other diffraction peaks were observed for both the two samples, this indicates that both the two samples were amorphous. To further investigate their structure, the transmission electron microscopy (TEM) was carried out and the corresponding selected area electron diffraction (SAED) patterns are shown in Fig. 3(b and c). The amorphous diffraction halos appared, this confirms their amorphous structure. Fig. 3(d) is the EDX spectrums of one of the chosen core–shell sphere, it can be found that the core–shell sphere contains Fe, Co, Ni, Mn, Mg, Gd, S and O elements, except Cu from the substrate. The element C, S and O should come from DMSO which was absorbed in the alloy because of the possible strong absorption caused by the particular hollow and core–shell nano-structures. The compositions of the alloys are averagely calculated from five different particles. After normalization, the hollow structure are Mg60.83Mn4.30Fe8.02Co3.51Ni5.52Gd17.82 (sample 1) and the core–shell nano-structure are is Mg12.03Mn11.46Fe30.10Co25.17Ni16.15Gd5.09 (sample 2). 3.3. Magnetic property The magnetic properties of the alloys were also studied, the results are shown in Fig. 4. Fig. 4(a and b) displays the hysteresis loops of the alloys at 5 K and 300 K respectively. At 5 K, both the

Fig. 3. (a) XRD patterns of hollow and core–shell spheres deposited at 2.0 V for 1 h in mixed solution 1 and solution 2 respectively, (b and c) are the SAED patterns of their corresponding TEM. (d) The EDX spectrums of the closed core–shell sphere.

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two samples exhibit ferromagnetic behavior with large saturation magnetization. The remanence (Mr) and coercivity (Hc) are about 9.9 emu/g and 1203 Oe for sample 1, respectively. However, Mr and Hc will reduce to 0.49 emu/g and 123 Oe for sample 2. At 300 K (Fig. 4(b)), sample 1 presents an excellent soft ferromagnetic behavior. It appears to be easy magnetized with a saturation magnetization of 12.6 emu/g at external field of 6 kOe. More interestingly, sample 2 showed the a negative slope of the magnetization as a function of external-field, which suggests a diamagnetic behavior property. Temperature dependence magnetization curves (M–T) were also given in Fig. 4(c and d). For sample 1, the magnetization decreases steadily with increasing temperature. but it is noteworthy that there is a translation from hard ferromagnetism (5 K) to softferromagnetism (300 K). This is consistent with the result shown in Fig. 4 (a and b). The M–T curves of sample 2 (Fig. 4(d)) shows a magnetic abrupt jumption, which indicates an intrinsic transition from ferromagnetism to diamagnetism is occurred with increasing temperature. The critical transition temperatures is 53 K at the external field of 5 kOe, while it is reduced to 38 K at 10 kOe. This similar phenonmenon was also observed for Mg19.09Mn7.83Fe29.09Co16.93Ni17.91Gd9.15 alloy (Fig. 5). The magnetic transition was is observed until the field is higher than 20 kOe. Moreover, and the critical transition temperature was decreases with increasing the external field. 3.4. Theory discussion In recent years, the magnetic transitions of the crystal materials have been studied extensively including their mechanism, applications and novel physical properties, for example, the particular electron transport ability accompanying the magnetic transition [17,18]. Unfortunately, it’s still unknown what exactly lead to the

magnetic transitions. Some researchers attributed the magnetic transitions to crystal structure transition [19,20]. Also, the magnetic transition led by electron–electron correlation without structure change was accepted by other researchers. It is obvious that the structure transition mechanism is not suitable to explain our experimental results, because the Mg–Mn–Fe–Co–Ni–Gd alloy films are amorphous without any structure transition at room temperature. The magnetic transition of the as-prepared alloy films may be related to the amorphous structure and composition. Taking into account this concept, we proposed a mechanism which is based on Anderson–Mott transition to explain the magnetic transition. Anderson pointed out that in amorphous systems the energy band could separate into two parts: the localized states in the band tail and extended states in the middle of the band. There exists boundary between the two states which is defined as ‘‘mobility edge’’, as shown in Fig. 6 [21]. This transition can be caused by electron–electron interaction without structure transition. The frequency pseudo gap of localized states are mainly from the remnant geometric Bragg resonances, as well as those small electron–hole pockets on the Fermi surface [22]. The existing of the anisotropic effective masses in some directions for electron pockets is low, which is responsible for the diamagnetic of our sample [23]. Anderson–Mott transition is driven by electron–electron strong interaction, the mobility edge and Fermi face are relative to the compositions and structure of the materials. According to Kubo’s theory, the Fermi face can be described as: 2

EF ¼

h 2=3 ð3p2 nÞ 2m

h is Planck constant, m is the mass of electron, and n is the average electron density.

Fig. 4. Hysteresis loops of sample 1 (black line) and sample 2 (red line) at: (a) 5 K and (b) 300 K; M–T curves of the alloy of (c) sample 1 and (d) sample 2. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

H. Li et al. / Journal of Alloys and Compounds 598 (2014) 161–165

Fig. 5. M–T curves for several external field, used to study the magnetic transition in Mg19.09Mn7.83Fe29.09Co16.93Ni17.91Gd9.15 amorphous alloy.

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spheres via a simple electrodeposition method, their composition is characterized as Mg60.83Mn4.30Fe8.02Co3.51 Ni5.52Gd17.82 and Mg12.03Mn11.46Fe30.10Co25.17Ni16.15Gd5.09 respectively. Their surface morphologies can be tuned by the metal ions concentration. Most importantly, the core–shell structured Mg12.03Mn11.46Fe30.10Co25.17Ni16.15Gd5.09 amorphous alloy exhibits an abnormal magnetic first order transition from ferromagnetism to diamagnetism at the critical temperature of 53 K with the external field of 5 kOe and 38 K with external field of 10 kOe. It is found that the transitions are triggered by the external field while the sample exhibited only ferromagnetic in low field and magnetic transition in high field. The further research for the Mg19.09Mn7.83Fe29.09Co16.93Ni17.91Gd9.15 amorphous alloy showed that the critical transition temperature is nonlinear relationship with the external field. Finally, the Anderson–Mott transition mechanism was proposed to explain the transition, and accorded with the phenomenon very well. Acknowledgements This work was supported by the Natural Science Foundations of China (Grant No. 51101138), the Natural Science Foundations of Shanxi Province (Grant Nos. 2012021020-3, 2012011007-1 and 2013021011-5), the S&T Project of Shanxi Province (Grant Nos. 20121024 and 2013152), the Outstanding Innovative Teams of Higher Learning Institutions of Shanxi, the State Key Laboratory of Solidification Project (Grant No. SKLSP201227), and the Young Teacher Starting-up Research of Yuncheng University.

Fig. 6. The sketch map of Anderson–Mott transition.

When the electron around Fermi face sweep across the mobility edge, the Anderson–Mott transitions are triggered. In our experiment, the ferromagnetism to diamagnetism transition was observed for the alloys Mg12.03Mn11.46Fe30.10Co25.17Ni16.15Gd5.09 and Mg19.09Mn7.83Fe29.09Co16.93Ni17.91Gd9.15, while no diamagnetism transition appeared in the alloy Mg60.83Mn4.30Fe8.02Co3.51Ni5.52Gd17.82. For the latter, the Mg content is much higher than that in the former two alloys. There are only two valence shell electrons of 2 s2 in Mg atom with small atomic radius, which result in the average electron density of the alloy is smaller than others. Moreover, the value of EF decreases and the Fermi face shifts to the middle in the energy band (the Fermi face is far away from the mobility edge). This is disadvantage to the occurrence of Anderson–Mott transition, so the ferromagnetism to diamagnetism transition was not observed in the alloy Mg60.83Mn4.30 Fe8.02Co3.51Ni5.52Gd17.82. The experiment results also show that the ferromagnetism to diamagnetism transition is effected by temperature and external field. It can be seen in Fig. 5, at the same external field, the materials exhibited ferromagnetism at a low temperature and the ferromagnetism to diamagnetism transition occurred when the temperature reached a specific value. 4. Conclusions In summary, we have successfully synthesized amorphous Mg–Mn–Fe–Co–Ni–Gd hollow microspheres and core–shell micro-

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