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Chemical solution deposition of magnetoelectric ZnO–La2CoMnO6 nanocomposite thin films using a single precursor solution
Materials Chemistry and Physics 236 (2019) 121762
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Chemical solution deposition of magnetoelectric ZnO–La2CoMnO6 nanocomposite thin films using a single precursor solution Mizuki Saito, Manabu Hagiwara *, Shinobu Fujihara Department of Applied Chemistry, Faculty of Science and Technology, Keio University, 3-14-1 Hiyoshi, Kohoku-ku, Yokohama, 223-8522, Japan
H I G H L I G H T S
� ZnO–La2CoMnO6 nanocomposite film was fabricated by chemical solution deposition. � A vertical structure composed of ZnO columns and La2CoMnO6 nanograins was achieved. � The film showed ferromagnetic properties originating from the La2CoMnO6 phase. � The converse magnetoelectric response was observed at 10 K. A R T I C L E I N F O
A B S T R A C T
Keywords: Nanocomposite film Chemical solution deposition ZnO Ferromagnetics Magnetoelectric effect
We report a chemical solution deposition method for fabricating magnetoelectric ZnO–La2CoMnO6 (LCMO) nanocomposite films on platinum coated silicon (Pt/Si) substrates. A single precursor coating solution containing all constituent cations (Zn2þ, La3þ, Co2þ, and Mn2þ) was prepared by dissolving acetate and nitrate raw materials in 2-metoxyethanol in the presence of monoethanolamine. Investigations of decomposition and crystallization behaviors of a dried powder prepared from the ZnO–LCMO precursor solution revealed that a crystalline phase of ZnO first appeared at a low temperature below 400 � C followed by crystallization of LCMO at a much higher temperature. A film coated on a ZnO-seeded Pt/Si substrate was found to have a characteristic vertical nano composite structure composed of c-axis oriented ZnO columns and LCMO nanograins. The ZnO–LCMO film exhibited ferromagnetic properties originating from the LCMO phase. A converse magnetoelectric measurement at a temperature of 10 K demonstrated that the magnetization of the film at 45 kOe was decreased by 9% by application of an external electric field of 100 kV/cm.
1. Introduction The magnetoelectric (ME) effect—a change of magnetization induced by an applied electric field or vice versa—is attracting growing interest for its potential applications as transducers, magnetic field sensors, and information storage devices since the pioneering works of the 1950s–1960s [1]. The development of ME materials can be classified into two main approaches: single-phase materials and two-phase com posite systems. For single phase materials, the maximum allowable linear ME coefficient of a material is known to depend on its dielectric permittivity and magnetic permeability [1]. Hence the concept of mul tiferroics (i.e., ferroelectric (anti)ferromagnetics) has been regarded as an important strategy to discover superior single-phase ME materials because they may simultaneously have both high dielectric permittivity and high magnetic permeability. As a result, some multiferroics, such as
BiFeO3 [2] and rare-earth hexagonal manganites [3], have been re ported so far. The ME effect observed in these single-phase materials are, however, generally too small to be used in practical applications. Ferroelectric–ferromagnetic two-phase composite systems, where the ME coupling is achieved via the mechanical system (piezoelectric and magnetostrictive effects), have been extensively studied as an alternative approach to enhance the ME effect [4]. Among many com posite structures including 0–3 particulate and 2-2 laminate structures (the numbers represent the dimensions of connectivity of components in a composite), vertically aligned nanocomposite (VAN) thin films are of particular interest because an external electric field can be effectively applied to a ferroelectric (piezoelectric) phase in this structure [5]. For example of ME thin films with the VAN structure, BaTiO3–CoFe2O4 [6, 7], Na0.5Bi0.5TiO3–CoFe2O4 [8], BiFeO3–CoFe2O4 [9–11], and Pb (Zr0.52Ti0.48)O3–NiFe2O4 [12] systems have been reported. Recently, Fix
* Corresponding author. E-mail address: [email protected] (M. Hagiwara). https://doi.org/10.1016/j.matchemphys.2019.121762 Received 22 March 2019; Received in revised form 14 June 2019; Accepted 18 June 2019 Available online 23 June 2019 0254-0584/© 2019 Elsevier B.V. All rights reserved.
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Materials Chemistry and Physics 236 (2019) 121762
et al. [13] reported that a ZnO–La2CoMnO6 (LCMO) VAN thin film showed a large (60%) reversible change in magnetization by applying an electric field at 120 K. LCMO is a ferromagnetic oxide having a double-perovskite-type structure with rock-salt ordered B-site Co2þ and Mn4þ cations [14,15]. The magnetic properties of LCMO are highly affected by the valence states and the ordering degree of the B-site cations [15]. Importantly, LCMO possesses an electrical insulating property at low temperatures [16], and hence a high electric field can be applied to this material. With regard to the above-mentioned ZnO–LCMO VAN film, Fix et al. [13] have attributed the large ME ef fect to a charge-mediated mechanism based on the valence states of Mn and Co cations rather than the conventional strain-mediated coupling, although ZnO is a well-known piezoelectric material. According to their explanation, electrons conducting through the vertical interfaces be tween the ZnO and LCMO phases get trapped or detrapped on oxygen vacancies, leading to changes in the valence states of Mn and Co cations in the LCMO phase [13]. Such a novel ME coupling mechanism via the chemical system in VAN films has a great potential for development of advanced ME materials with superior performance. Most of VAN films reported so far, including the ZnO–LCMO films, have fabricated on single-crystal substrates by the pulsed laser deposi tion (PLD) method. The deposition of VAN films by the PLD method involves three stages: (i) surface diffusion of atomic clusters, (ii) epitaxial nucleation and island growth of two phases, and (iii) columnar growth [17]. The lattice matching between the substrate and the pro duced crystalline phases is of great importance to achieve the controlled phase separation in the second stage of the deposition [7], and hence only single-crystal substrates (mostly SrTiO3) are applicable to this deposition method of VAN films. For this reason, the PLD method does not match the modern Si-based electronics although it is useful for searching the new composite systems. In contrast to the PLD method, the chemical solution deposition (CSD) method is a low cost, non-vacuum, high-speed deposition technique suitable for industrial fabrication of functional oxide films on large-area substrates [18]. The CSD method has so far been widely used for fabri cating high-quality oxide thin films including ZnO and perovskites. For instance, it has been reported that thin films of multiferroic BiFeO3 fabricated by the CSD method show excellent ferroelectric properties comparable to PLD-derived thin films [19,20]. However, there have been few reports on the fabrication of VAN thin films by the CSD method. In the CSD method, the nucleation and growth of crystalline phases occur in an amorphous solid. Because diffusion of ions in the amorphous film is much slower than the surface diffusion in the PLD method, it is a challenge to control the phase separated structure to obtain VAN films by the CSD method. In this paper, we report a CSD technique for fabricating ZnO–LCMO nanocomposite films on Si substrates. In our method, a single precursor solution containing the constitutive cations of both ZnO and LCMO is spin-coated on a Si substrate with a c-axis oriented ZnO seed layer. A large difference in the crystallization temperatures of the ZnO and LCMO phases results in the phase separation into these phases during heating, leading to the formation of a ZnO–LCMO nanocomposite film composed of vertical ZnO columns and LCMO nanograins. We also demonstrate that the resulting ZnO–LCMO nanocomposite film possesses a ferromagnetic property as well as a reversible change of magnetization by applying an electric field.
To study the crystallization behavior from the solution, the precursor so lution was first dried at 90 � C for 24 h and then further dried at 200 � C for 1 h. The resulting dried powder was calcinated at a temperature of 400, 500, 600 or 700 � C for 2 h in air at 10 � C/min to obtain powder samples. For the fabrication of film samples, a c-axis oriented ZnO seed layer was first fabricated on Pt(111)/Ti/SiO2/Si(100) substrates (hereafter abbreviated as Pt/Si substrates) by the following CSD method [21,22]. Zn (CH3COO)2⋅2H2O was dissolved in a mixed solvent of 2-MOE (9.70 mL) and MEA (0.30 mL) by stirring at room temperature to obtain a trans parent ZnO solution with a concentration of 0.5 M. The solution was dropped on a Pt/Si substrate (25 � 25 mm2) and spin-coated at 1000 rpm for 10 s and then 3000 rpm for 50 s. The substrate was then immediately subjected to heat treatment at 400 � C for 1 h in air to crystalize ZnO. Next, the ZnO–LCMO precursor solution was dropped on the ZnO-seeded Pt/Si substrate, and spin-coated under the same condition with the ZnO solu tion. The resulting film was dried at 150 � C for 1 min and then heated at 700 � C for 10 min in air. After repeating the coating of the ZnO–LCMO solution followed by the drying and heating 3 or 20 times, the films were lastly heated at 700 � C for 1 h in air. Hereafter a film sample fabricated by repeating the coating cycle n times will be referred to as the ZnO–LCMO (n) film (where, n ¼ 3 or 20). To study the effect of the ZnO seed layer on the film structure, a ZnO–LCMO (3) film was fabricated also on a bare Pt/Si substrate without the seed layer. Thermogravimetric–differential thermal analysis (TG–DTA) of the dried powder prepared from the ZnO–LCMO precursor solution was carried out from room temperature to 800 � C under flowing air using a SHIMADZU DTG-60 instrument. The crystal structure of the powder and film samples was identified by X-ray diffraction (XRD; Bruker AXS D8) analysis using Cu Kα radiation. The microstructure of the film samples was observed by field emission scanning electron microscopy (FE-SEM; JSM-7600F, JEOL). To investigate the composite structure of ZnO and LCMO, the ZnO–LCMO (20) film was immerged in glacial acetic acid (99.7%, Wako) for 5 min and then observed by FE-SEM. A magnetization (M)–Magnetic field (H) curve at a temperature (T) of 10 K and a field cooling M–H curve at H ¼ 1 kOe of the ZnO–LCMO (20) film were measured by a Quantum Design MPMS SQUID magnetometer. For the measurement of the converse ME response, an Au top electrode with an area of approximately 4 � 6 mm2 was sputtered on the surface of a small piece of the ZnO–LCMO (20) film (approximately 5 � 10 mm2 in area), and the bare area of the film was chemically etched by an acidic hydrogen peroxide solution (H2O: HCl: H2O2 ¼ 8 : 4: 1) to contact the Pt bottom electrode. Two Cu wires are connected to the top and bottom electrodes and an external DC voltage was applied between the electrodes using a Keithley 2400 source meter. M–H curves of the film at T ¼ 10 K were measured by the SQUID magnetometer under applying electric fields of 0, 100, and 0 kV/cm in turn. Considering the sample size for the magnetic and ME measurements, the experimental error in the measured magne tization values is assumed to be less than 3% [23,24]. 3. Results and discussion 3.1. Crystallization behavior A purple homogeneous ZnO–LCMO precursor coating solution with the total metal ion concentration of 0.6 M was obtained by using the mixed solvent composed of 2-MOE (9.64 mL) and MEA (0.36 mL). We found that solvents with a lower MEA/2-MOE ratio could not completely dissolve the Mn source. Fig. 1 shows the TG–DTA curves of the dried powder prepared from the ZnO–LCMO precursor solution. Two main weight losses with strong exothermic peaks are observed at 277 � C and 393 � C. The overall weight loss is 43%. It is known that Zn (CH3COO)2⋅2H2O starts to dehydrate at about 50 � C and then de composes in to ZnO at 270 � C upon heating in air [25,26]. The first weight loss at 277 � C is thus likely to be attributed to the thermal decomposition of dehydrated zinc acetate or its complex with MEA. The second weight loss at 393 � C would be due to the combustion of the
2. Experimental Zn(CH3COO)2⋅2H2O (99.9%, Wako Pure Chemical Industries; 5 mmol), Mn(CH3COO)2⋅4H2O (99.9%, Wako; 0.25 mmol), La (NO3)3⋅6H2O (99.9%, Wako; 0.50 mmol), and Co(CH3COO)2⋅4H2O (99.0%, Wako; 0.25 mmol) were dissolved in a mixed solvent of 2-methox yethanol (2-MOE; 9.64 mL) and monoethanolamine (MEA; 0.34 mL) by stirring at 60 � C for 1 h to obtain a ZnO–LCMO precursor coating solution. The total concentration of the metal ions in the solution was 0.6 M and the mole fraction of LCMO, namely 100(%) � LCMO/(ZnO þ LCMO), was 5%. 2
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Materials Chemistry and Physics 236 (2019) 121762
The above observation shows that the presence of La, Co, Mn ions in the precursor does not prevent the crystallization of ZnO. The powder heated at 500 � C also shows the diffraction peaks only from the ZnO phase. In the powders heated at 400 or 500 � C, the residual La, Co, Mn ions should exist as a form of an amorphous oxide rather than acetates or nitrates because no significant weight loss is observed above 400 � C as shown in Fig. 1. When the precursor is heated at 600 or 700 � C, additional peaks appear at around 2θ ¼ 32.7, 39.0, and 46.8� , which can be indexed to (112), (202), and (004) planes of a pseudo tetragonal lattice of LCMO with the double-perovskite-type structure [15]. Any other peaks due to the im purity phase is not observed in the XRD patterns. This result shows that the crystallization of LCMO in the powder occurs at a temperature be tween 500 and 600 � C, which is somewhat inconsistent with the result of DTA where the small exothermic peak without weight loss is found below 500 � C. This possibly arises from the difference in the atmosphere con ditions during the TG–DTA measurement and the calcination. The TG-DTA curves were corrected with flowing air, whereas the calcination was performed in a closed box furnace. The oxygen partial pressure during the calcination should be thus lowered due to the combustion of organic species. The lowered oxygen partial pressure delays the oxidation of Mn2þ into Mn4þ difficult, leading to a higher crystallization temper ature of LCMO compared to the TG-DTA measurement. Therefore, it is difficult to determine the exact crystallization temperature of LCMO from the TG-DTA and calcination experiments, but nevertheless, the results are clearly showing that there is a significant difference between the crys tallization temperatures of the ZnO and LCMO phases. Such a large gap in the crystallization temperatures is a key to obtain ZnO–LCMO composite films with a controlled nanostructure.
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residual organic compounds. The nitrate ion originating from the La source can be involved in the combustion process as an internal oxidizer. From the presented thermal behavior of the ZnO–LCMO dried powder, it is found that the decomposition of the powder is completed at a rela tively low temperature below 400 � C. It should be noted here that a smaller broad exothermic peak without a corresponding weight change is also found in a temperature range between 460 � C and 500 � C, as shown in the inset of Fig. 1, which is presumably due to the crystalli zation of LCMO as discussed below. The dried powder was calcined in air at a temperature of 400, 500, 600, or 700 � C for 2 h to study its crystallization behavior. Fig. 2 shows the XRD patterns of the resultant powders. The powder heated at 400 � C shows clear diffraction peaks at 2θ ¼ 31.2, 34.6, 35.9, and 47.3� . All these peaks can be indexed to the wurtzite-type ZnO phase (ICDD 089–0510). This indicates that the crystallization of ZnO accompanies the decom position and combustion of organic species observed below 400 � C in the TG-DTA curves. It has been reported that the crystallization of ZnO from a precursor solution without any other cations starts at round 400 � C [22].
3.2. Film structures Based on the crystallization behavior studied above, the ZnO–LCMO precursor solution was spin-coated on a ZnO-seeded Pt/Si substrate and
Fig. 3. (a) An XRD pattern of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C and (b) an enlarged image of (a) between 30 and 37� .
Fig. 2. XRD patterns of the powder samples prepared by heating the ZnO–LCMO dired powder at a temperature of 400, 500, 600, or 700 � C for 2 h. 3
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Materials Chemistry and Physics 236 (2019) 121762
heated at 700 � C to obtain ZnO–LCMO nanocomposite films. The XRD pattern of the ZnO–LCMO (20) film is presented in Fig. 3. All the diffraction peaks observed in the 2θ range between 10 and 60� can be attributed to ZnO (ICDD 089–0510), LCMO [15], or the substrate, confirming that a composite between ZnO and LCMO is formed and no significant reaction among these phases and the substrate occurs. It is seen in the magnified image around the strongest ZnO (002) peak that the ZnO (100) and (101) peaks are very weak, showing that the ZnO phase in the film has a strong c-axis orientation. The position of the ZnO (002) peak is found at 2θ ¼ 34.61� , which is slightly higher than the reference data (ICDD 089–0510; 2θ ¼ 34.43� ). This suggests the possi bility of the partial substitution of Zn2þ ions (60 p.m.) by the smaller Co2þ ions (56 p.m.) [27]. Fig. 4(a) and (b) are FE-SEM images of the surface and cross-section of the ZnO–LCMO (20) film, respectively. A characteristic microstruc ture is observed on the surface of the film, where small nanograins about 50 nm are filling up intergranular gaps of larger grains about 200 nm. The cross-sectional image shows a column-like dense structure with a thickness about 600 nm. With an aim to get further insight into the film structure, the film was immersed in glacial acetic acid and then observed by FE-SEM again because ZnO easily dissolves in acetic acid whereas LCMO does not. In the surface FE-SEM image shown in Fig. 4(c), it is seen that the large grains have disappeared by the acetic-acid treatment while the small nanograins have remained. This result clearly demon strates that the larger grains seen in the as-deposited film are ZnO whereas the nanograins are LCMO. It is also observed in the acetic-acid treated film that there are many vertical holes. These holes should be formed by the dissolution of column-like ZnO grains. Based on these observations together with the XRD data, the ZnO–LCMO (20) film is found to have a characteristic vertical nanostructure as illustrated in
Fig. 4(d), where the nanograins of LCMO fills the gaps between the c-axis oriented ZnO columns. It should be noted that the interface between the ZnO seed layer and the film is not clearly observed in the cross-sectional image (Fig. 4(b)) because the ZnO columns were homoepitaxially grown on the seed layer. To reveal the role of the ZnO seed layer on the formation of the vertical nanostructure, structures of the ZnO–LCMO (3) films fabricated on bare and ZnO-seeded Pt/Si substrates were compared. The coating cycle was repeated for only 3 times in order to observe the early stage of the film grows where the effect of the seed layer is expected to be sig nificance. The XRD patterns and the surface FE-SEM images of the films are shown in Fig. 5. The film fabricated on the ZnO-seeded Pt/Si sub strate shows a strong ZnO (002) peak and its microstructure is similar to that of the ZnO–LCMO (20) film (Fig. 4(a)). This confirms that the ZnO columns were directly grown on the ZnO seed layer. On the other hand, the intensity of the ZnO (002) peak of the film fabricated on the bare substrate is much smaller and comparable to that of the ZnO (110) peak, indicating that the ZnO phase in this film are randomly oriented. Furthermore, no large grain attributable to the ZnO column is found on the surface of this film. Thus, it is clearly demonstrated that the ZnO seed layer plays a critical role on the growth of c-axis oriented ZnO columns. From the observations shown above, the mechanism underlying the formation of the ZnO–LCMO nanocomposite film can be understood as follows. Upon heating after the spin-coating process, heterogeneous nucleation of ZnO first occurs on the c-axis oriented ZnO seed layer at a temperature less than 400 � C. The ZnO crystal nuclei also have the c-axis orientation to match the crystal lattice with the seed layer. The ZnO nuclei grow mainly in the direction toward the film surface (along the caxis) to form column-like grains during further heated up, and the amorphous oxide containing the residual ions of La, Mn, and Co is
Fig. 4. Microstructure of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C: (a) Surface and (b) cross-sectional FE-SEM images of the asdeposited film, (c) a surface FE-SEM image of the film after immerged in glacial acetic acid for 5 min, and (d) a schematic illustration of the film structure. 4
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Materials Chemistry and Physics 236 (2019) 121762
Fig. 5. (a, b) XRD patterns and (c, d) FE-SEM images of the ZnO–LCMO (3) film prepared on (a, c) ZnO-seeded and (b, d) bare Pt/Si substrates at 700 � C.
Magnetization, M/emu·cm-3
concentrated in regions between the ZnO columns. Then, the homoge neous nucleation of the LCMO phase occurs in the amorphous oxide at a temperature between 450 and 600 � C, leading to the formation of the LCMO nanograins between the ZnO columns. The similar phase sepa ration process, which starts from the heterogeneous nucleation of ZnO on the ZnO columns, is repeated in the subsequent coating cycles, resulting in the formation of the vertically aligned nanocomposite film composed of the ZnO columns and the LCMO nanograins. 3.3. Magnetic and magnetoelectric properties Fig. 6(a) shows the M–H curve of the ZnO–LCMO (20) film measured at 10 K. A clear hysteresis with a saturation magnetization (Ms) of 58 emu/cm3 is observed. The Ms value of LCMO at 5 K is reported to range between 3.36 and 4.85 μB/f.u., which is corresponding to 260–380 emu/ cm3, depending on fabrication conditions such as firing temperature and atmosphere [15]. Since the mole fraction of LCMO in the composite film is 5%, the volume fraction of LCMO is calculated as 21% using the unit cell volumes of ZnO and LCMO (47.7 Å3 and 236 Å3, respectively, where both unit cells include two formula units). Considering this volume fraction of LCMO, the observed Ms value of the ZnO–LCMO (20) film is well consistent with the reported values. The M–T curve of the film measured at 1 kOe upon cooling is shown in Fig. 6(b). The M–T curve shows that the Currie temperature (TC), above which the ferromagne tism is vanished, of the film is at round 200 K. It has been reported that TC of LCMO is also varied in a range between 130 and 230 K depending on the preparation condition [15]. The observed TC of the ZnO–LCMO nanocomposite film is within the reported TC range. It should be noted that Co-doped ZnO also shows ferromagnetic behaviors [28,29], and the XRD pattern of the ZnO–LCMO (20) film indicates a possibility of the partial substitution of Co2þ for Zn2þ in the ZnO phase. However, TC of Co-doped ZnO is reported much higher than room temperature [29], and its Ms should be much smaller than the observation in this study. It can be thus concluded that the ZnO–LCMO (20) film has ferromagnetic properties originating from the LCMO phase, which are not degraded by the coexistence of the ZnO phase. To observe the converse ME response of the ZnO–LCMO (20) film, an M–H curve under an out-of-plane external electric field was measured. Fig. 7 shows a part of M–H hysteresis curves measured with and without applying an electric field of 100 kV/cm. It is observed that the
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showed ferromagnetic properties originating from the LCMO phase as well as the magnetoelectric response to the external electric field. The reduction of magnetization was 9% at an applied electric field of 100 kV/cm. The CSD method developed in this study enables low-cost and large-area fabrication of ZnO-based VAN films on Si substrates.
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Acknowledgement
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The authors wish to thank Prof. T. Sato of Keio university for his help in the ME measurement. M. S. thanks the Kato Foundation for Promotion of Science for its financial support towards this work.
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[1] W. Eerenstein, N.D. Mathur, J.F. Scott, Multiferroic and magnetoelectric materials, Nature 442 (2006) 759–765, https://doi.org/10.1038/nature05023. [2] J. Wang, J.B. Neaton, H. Zheng, V. Nagarajan, S.B. Ogale, B. Liu, D. Viehland, V. Vaithyanathan, D.G. Schlom, U. V Waghmare, N.A. Spaldin, K.M. Rabe, M. Wuttig, R. Ramesh, Epitaxial BiFeO3 multiferroic thin film heterostructures, Science 299 (2003) 1719–1722, https://doi.org/10.1126/science.1080615. [3] Z.J. Huang, Y. Cao, Y.Y. Sun, Y.Y. Xue, C.W. Chu, Coupling between the ferroelectric and antiferromagnetic orders in YMnO3, Phys. Rev. B 56 (1997) 2623–2626, https://doi.org/10.1103/PhysRevB.56.2623. [4] J. Ma, J. Hu, Z. Li, C.W. Nan, Recent progress in multiferroic magnetoelectric composites: from bulk to thin films, Adv. Mater. 23 (2011) 1062–1087, https://doi. org/10.1002/adma.201003636. [5] C.-W. Nan, G. Liu, Y. Lin, H. Chen, Magnetic-field-induced electric polarization in multiferroic nanostructures, Phys. Rev. Lett. 94 (2005) 197203, https://doi.org/ 10.1103/PhysRevLett.94.197203. [6] H. Zheng, J. Wang, L. Mohaddes-Ardabili, M. Wuttig, L. Salamanca-Riba, D. G. Schlom, R. Ramesh, Three-dimensional heteroepitaxy in self-assembled BaTiO3–CoFe2O4 nanostructures, Appl, Phys. Lett. 85 (2004) 2035–2037, https:// doi.org/10.1063/1.1786653. [7] N. Dix, R. Muralidharan, J. Rebled, S. Estrad� e, F. Peir� o, M. Varela, J. Fontcuberta, F. S� anchez, Selectable spontaneous polarization direction and magnetic anisotropy in BiFeO3 CoFe2O4 Epitaxial Nanostructures, ACS Nano 4 (2010) 4955–4961, https://doi.org/10.1021/nn101546r. [8] R. Wu, A. Kursumovic, X. Gao, C. Yun, M.E. Vickers, H. Wang, S. Cho, J. L. MacManus-Driscoll, Design of a vertical composite thin film system with ultralow leakage to yield large converse magnetoelectric Effect, ACS Appl. Mater. Interfaces 10 (2018) 18237–18245, https://doi.org/10.1021/acsami.8b03837. [9] L. Yan, Z. Wang, Z. Xing, J. Li, D. Viehland, Magnetoelectric and multiferroic properties of variously oriented epitaxial BiFeO3–CoFe2O4 nanostructured thin films, J. Appl. Phys. 107 (2015), https://doi.org/10.1063/1.3359650, 064106. [10] C. Zhang, S.C. Huberman, S. Ning, J. Pelliciari, R.A. Duncan, B. Liao, S. Ojha, J. W. Freeland, K.A. Nelson, R. Comin, G. Chen, C.A. Ross, Thermal conductivity in self-assembled CoFe2O4/BiFeO3 vertical nanocomposite films, Appl. Phys. Lett. 113 (2018) 1–6, https://doi.org/10.1063/1.5049176. [11] T.C. Kim, S.H. Lee, H.K. Jung, H. Lee, J.H. Mun, I. Oh, D.H. Kim, Magnetic property modulation of sputter-grown BiFeO3–CoFe2O4 nanocomposite thin films, Ceram. Int. 45 (2019) 12182–12188, https://doi.org/10.1016/j.ceramint.2019.03.122. [12] H. Ryu, P. Murugavel, J.H. Lee, S.C. Chae, T.W. Noh, Y.S. Oh, H.J. Kim, K.H. Kim, J.H. Jang, M. Kim, C. Bae, J.-G. Park, Magnetoelectric effects of nanoparticulate Pb (Zr0.52Ti0.48)O3–NiFe2O4 composite films, Appl. Phys. Lett. 89 (2006) 102907, https://doi.org/10.1063/1.2338766. [13] T. Fix, E.-M. Choi, J.W.A. Robinson, S.B. Lee, A. Chen, B. Prasad, H. Wang, M. G. Blamire, J.L. MacManus-Driscoll, Electric-field control of ferromagnetism in a nanocomposite via a ZnO phase, Nano Lett. 13 (2013) 5886–5890, https://doi.org/ 10.1021/nl402775h. [14] C.L. Bull, D. Gleeson, K.S. Knight, Determination of B-site ordering and structural transformations in the mixed transition metal perovskites La2CoMnO6 and La2NiMnO6, J. Phys. Condens. Matter 15 (2003) 4927–4936, https://doi.org/ 10.1088/0953-8984/15/29/304. [15] R. Dass, J. Goodenough, Multiple magnetic phases of La2CoMnO6 δ (0�δ�0.05), Phys. Rev. B 67 (2003), https://doi.org/10.1103/PhysRevB.67.014401, 014401. [16] M.P. Singh, K.D. Truong, P. Fournier, Magnetodielectric effect in double perovskite La2CoMnO6 thin films, Appl. Phys. Lett. 91 (2007), https://doi.org/10.1063/ 1.2762292, 042504. [17] A. Chen, Z. Bi, Q. Jia, J.L. MacManus-Driscoll, H. Wang, Microstructure, vertical strain control and tunable functionalities in self-assembled, vertically aligned nanocomposite thin films, Acta Mater. 61 (2013) 2783–2792, https://doi.org/ 10.1016/j.actamat.2012.09.072. [18] R.W. Schwartz, T. Schneller, R. Waser, Chemical solution deposition of electronic oxide films, Compt. Rendus Chem. 7 (2004) 433–461, https://doi.org/10.1016/j. crci.2004.01.007. [19] S.K. Singh, Y.K. Kim, H. Funakubo, H. Ishiwara, Epitaxial BiFeO3 thin films fabricated by chemical solution deposition, Appl. Phys. Lett. 88 (2006) 86–89, https://doi.org/10.1063/1.2196477. [20] Q. Zhang, N. Valanoor, O, Standard, Chemical solution deposition derived (001)oriented epitaxial BiFeO3 thin films with robust ferroelectric properties using stoichiometric precursors, J. Appl. Phys. 116 (2014), https://doi.org/10.1063/ 1.4891311.
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Magnetic field, H/kOe Fig. 7. Magnetization–magnetic field curves of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C. The curves were measured before, under, and after application of an electric field of 100 kV/cm at 10 K. The inset shows the variation of the saturation magnetization at 45 kOe.
magnetization at high magnetic fields is slightly decreased by applying the electric field. The magnetization is switched back to the initial state when the electric field is removed. On the other hand, the application of the electric filed is found to have almost no impact on the coercive force (Hc). This implies that the reduction of magnetization is not due to the Joule heating of the film during the measurement because Hc should be significantly decreased by heating. The reduction of Ms is 9% at 45 kOe, which is much smaller than the reported value for the ZnO–LCMO nanocomposite film prepared by the PLD method (60% reduction at an electric field of 305 kV/cm) [13]. There are some possible reasons for the smaller ME response of the sample obtained in this study. The first reason is related to the magnitude of the applied electric field. For the ZnO–LCMO nanocomposite film by the PLD method, it has been re ported that the reduction of magnetization depends nonlinearly on the applied electric field and becomes significant when the field exceeds 305 kV/cm [13]. Indeed, the reduction is only 14% at 289 kV/cm [13], which is comparable to the observation of this study. We couldn’t apply electric fields higher than 100 kV/cm because of the electrical break down of the sample. The poorer electrical break down strength of the sample would be related to the somewhat porous microstructure in the LCMO region as seen in Fig. 4(a). The second possible reason is the interface structure between ZnO and LCMO. Because the LCMO grains are smaller than the gaps between the ZnO columns, the majority of the LCMO grains are out of contact with the ZnO columns. These LCMO grains should be hidden from the ZnO/LCMO interfaces where the electron trap/detrap occurs. Therefore, we consider that the further modification of the microstructure of the LCMO region is needed to improve the ME response of the film. Nevertheless, the result shown here is of great importance for demonstrating that the CSD method is effec tive for fabricating magnetoelectric ZnO–LCMO VAN films on Si sub strates. It is also believed that the CSD method developed in this study can be applied to other ZnO-based vertical nanocomposite systems for various applications such as the ZnO–La0.7Sr0.3MnO3 system showing a strong magnetoresistance effect [30,31]. 4. Conclusion The ZnO–LCMO vertical nanocomposite film was deposited on the Pt/Si substrate using the single precursor solution. It was found that the use of the ZnO seed layer and the large difference in the crystallization temperatures of ZnO and LCMO were important for obtaining the ver tical nanocomposite structure. The ZnO–LCMO nanocomposite film 6
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Materials Chemistry and Physics 236 (2019) 121762 [27] R.D. Shannon, Revised effective ionic radii and systematic studies of interatomic distances in halides and chalcogenides, Acta Crystallogr. A 32 (1976) 751–767, https://doi.org/10.1107/S0567739476001551. [28] H.-J. Lee, S.-Y. Jeong, C.R. Cho, C.H. Park, Study of diluted magnetic semiconductor: Co-doped ZnO, Appl. Phys. Lett. 81 (2002) 4020–4022, https:// doi.org/10.1063/1.1517405. [29] C. Song, K.W. Geng, F. Zeng, X.B. Wang, Y.X. Shen, F. Pan, Y.N. Xie, T. Liu, H. T. Zhou, Z. Fan, Giant magnetic moment in an anomalous ferromagnetic insulator: Co-doped ZnO, Phys. Rev. B 73 (2006), https://doi.org/10.1103/ PhysRevB.73.024405, 024405. [30] J.L. MacManus-Driscoll, P. Zerrer, H. Wang, H. Yang, J. Yoon, A. Fouchet, R. Yu, M.G. Blamire, Q. Jia, Strain control and spontaneous phase ordering in vertical nanocomposite heteroepitaxial thin films, Nat. Mater. 7 (2008) 314–320, https:// doi.org/10.1038/nmat2124. [31] A. Chen, Z. Bi, C.-F. Tsai, J. Lee, Q. Su, X. Zhang, Q. Jia, J.L. MacManus-Driscoll, H. Wang, Tunable low-field magnetoresistance in (La0.7Sr0.3MnO3)0.5:(ZnO)0.5 selfassembled vertically aligned nanocomposite thin films, Adv. Funct. Mater. 21 (2011) 2423–2429, https://doi.org/10.1002/adfm.201002746.
[21] S. Fujihara, C. Sasaki, T. Kimura, Effects of Li and Mg doping on microstructure and properties of sol-gel ZnO thin films, J. Eur. Ceram. Soc. 21 (2001) 2109–2112, https://doi.org/10.1016/S0955-2219(01)00182-0. [22] S. Fujihara, C. Sasaki, T. Kimura, Crystallization behavior and origin of c-axis orientation in sol–gel-derived ZnO:Li thin films on glass substrates, Appl. Surf. Sci. 180 (2001) 341–350, https://doi.org/10.1016/S0169-4332(01)00367-1. [23] P. Stamenov, J.M.D. Coey, Sample size, position, and structure effects on magnetization measurements using second-order gradiometer pickup coils, Rev. Sci. Instrum. 77 (2006), https://doi.org/10.1063/1.2149190, 015106. [24] M. Sawicki, W. Stefanowicz, A. Ney, Sensitive SQUID magnetometry for studying nanomagnetism, Semicond. Sci. Technol. 26 (2011), https://doi.org/10.1088/ 0268-1242/26/6/064006, 064006. [25] C.-C. Lin, Y.-Y. Li, Synthesis of ZnO nanowires by thermal decomposition of zinc acetate dihydrate, Mater. Chem. Phys. 113 (2009) 334–337, https://doi.org/ 10.1016/j.matchemphys.2008.07.070. [26] T. Arii, A. Kishi, The effect of humidity on thermal process of zinc acetate, Thermochim. Acta 400 (2003) 175–185, https://doi.org/10.1016/S0040-6031(02) 00487-2.
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Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Chemical solution deposition of magnetoelectric ZnO–La2CoMnO6 nanocomposite thin films using a single precursor solution Mizuki Saito, Manabu Hagiwara *, Shinobu Fujihara Department of Applied Chemistry, Faculty of Science and Technology, Keio University, 3-14-1 Hiyoshi, Kohoku-ku, Yokohama, 223-8522, Japan
H I G H L I G H T S
� ZnO–La2CoMnO6 nanocomposite film was fabricated by chemical solution deposition. � A vertical structure composed of ZnO columns and La2CoMnO6 nanograins was achieved. � The film showed ferromagnetic properties originating from the La2CoMnO6 phase. � The converse magnetoelectric response was observed at 10 K. A R T I C L E I N F O
A B S T R A C T
Keywords: Nanocomposite film Chemical solution deposition ZnO Ferromagnetics Magnetoelectric effect
We report a chemical solution deposition method for fabricating magnetoelectric ZnO–La2CoMnO6 (LCMO) nanocomposite films on platinum coated silicon (Pt/Si) substrates. A single precursor coating solution containing all constituent cations (Zn2þ, La3þ, Co2þ, and Mn2þ) was prepared by dissolving acetate and nitrate raw materials in 2-metoxyethanol in the presence of monoethanolamine. Investigations of decomposition and crystallization behaviors of a dried powder prepared from the ZnO–LCMO precursor solution revealed that a crystalline phase of ZnO first appeared at a low temperature below 400 � C followed by crystallization of LCMO at a much higher temperature. A film coated on a ZnO-seeded Pt/Si substrate was found to have a characteristic vertical nano composite structure composed of c-axis oriented ZnO columns and LCMO nanograins. The ZnO–LCMO film exhibited ferromagnetic properties originating from the LCMO phase. A converse magnetoelectric measurement at a temperature of 10 K demonstrated that the magnetization of the film at 45 kOe was decreased by 9% by application of an external electric field of 100 kV/cm.
1. Introduction The magnetoelectric (ME) effect—a change of magnetization induced by an applied electric field or vice versa—is attracting growing interest for its potential applications as transducers, magnetic field sensors, and information storage devices since the pioneering works of the 1950s–1960s [1]. The development of ME materials can be classified into two main approaches: single-phase materials and two-phase com posite systems. For single phase materials, the maximum allowable linear ME coefficient of a material is known to depend on its dielectric permittivity and magnetic permeability [1]. Hence the concept of mul tiferroics (i.e., ferroelectric (anti)ferromagnetics) has been regarded as an important strategy to discover superior single-phase ME materials because they may simultaneously have both high dielectric permittivity and high magnetic permeability. As a result, some multiferroics, such as
BiFeO3 [2] and rare-earth hexagonal manganites [3], have been re ported so far. The ME effect observed in these single-phase materials are, however, generally too small to be used in practical applications. Ferroelectric–ferromagnetic two-phase composite systems, where the ME coupling is achieved via the mechanical system (piezoelectric and magnetostrictive effects), have been extensively studied as an alternative approach to enhance the ME effect [4]. Among many com posite structures including 0–3 particulate and 2-2 laminate structures (the numbers represent the dimensions of connectivity of components in a composite), vertically aligned nanocomposite (VAN) thin films are of particular interest because an external electric field can be effectively applied to a ferroelectric (piezoelectric) phase in this structure [5]. For example of ME thin films with the VAN structure, BaTiO3–CoFe2O4 [6, 7], Na0.5Bi0.5TiO3–CoFe2O4 [8], BiFeO3–CoFe2O4 [9–11], and Pb (Zr0.52Ti0.48)O3–NiFe2O4 [12] systems have been reported. Recently, Fix
* Corresponding author. E-mail address: [email protected] (M. Hagiwara). https://doi.org/10.1016/j.matchemphys.2019.121762 Received 22 March 2019; Received in revised form 14 June 2019; Accepted 18 June 2019 Available online 23 June 2019 0254-0584/© 2019 Elsevier B.V. All rights reserved.
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et al. [13] reported that a ZnO–La2CoMnO6 (LCMO) VAN thin film showed a large (60%) reversible change in magnetization by applying an electric field at 120 K. LCMO is a ferromagnetic oxide having a double-perovskite-type structure with rock-salt ordered B-site Co2þ and Mn4þ cations [14,15]. The magnetic properties of LCMO are highly affected by the valence states and the ordering degree of the B-site cations [15]. Importantly, LCMO possesses an electrical insulating property at low temperatures [16], and hence a high electric field can be applied to this material. With regard to the above-mentioned ZnO–LCMO VAN film, Fix et al. [13] have attributed the large ME ef fect to a charge-mediated mechanism based on the valence states of Mn and Co cations rather than the conventional strain-mediated coupling, although ZnO is a well-known piezoelectric material. According to their explanation, electrons conducting through the vertical interfaces be tween the ZnO and LCMO phases get trapped or detrapped on oxygen vacancies, leading to changes in the valence states of Mn and Co cations in the LCMO phase [13]. Such a novel ME coupling mechanism via the chemical system in VAN films has a great potential for development of advanced ME materials with superior performance. Most of VAN films reported so far, including the ZnO–LCMO films, have fabricated on single-crystal substrates by the pulsed laser deposi tion (PLD) method. The deposition of VAN films by the PLD method involves three stages: (i) surface diffusion of atomic clusters, (ii) epitaxial nucleation and island growth of two phases, and (iii) columnar growth [17]. The lattice matching between the substrate and the pro duced crystalline phases is of great importance to achieve the controlled phase separation in the second stage of the deposition [7], and hence only single-crystal substrates (mostly SrTiO3) are applicable to this deposition method of VAN films. For this reason, the PLD method does not match the modern Si-based electronics although it is useful for searching the new composite systems. In contrast to the PLD method, the chemical solution deposition (CSD) method is a low cost, non-vacuum, high-speed deposition technique suitable for industrial fabrication of functional oxide films on large-area substrates [18]. The CSD method has so far been widely used for fabri cating high-quality oxide thin films including ZnO and perovskites. For instance, it has been reported that thin films of multiferroic BiFeO3 fabricated by the CSD method show excellent ferroelectric properties comparable to PLD-derived thin films [19,20]. However, there have been few reports on the fabrication of VAN thin films by the CSD method. In the CSD method, the nucleation and growth of crystalline phases occur in an amorphous solid. Because diffusion of ions in the amorphous film is much slower than the surface diffusion in the PLD method, it is a challenge to control the phase separated structure to obtain VAN films by the CSD method. In this paper, we report a CSD technique for fabricating ZnO–LCMO nanocomposite films on Si substrates. In our method, a single precursor solution containing the constitutive cations of both ZnO and LCMO is spin-coated on a Si substrate with a c-axis oriented ZnO seed layer. A large difference in the crystallization temperatures of the ZnO and LCMO phases results in the phase separation into these phases during heating, leading to the formation of a ZnO–LCMO nanocomposite film composed of vertical ZnO columns and LCMO nanograins. We also demonstrate that the resulting ZnO–LCMO nanocomposite film possesses a ferromagnetic property as well as a reversible change of magnetization by applying an electric field.
To study the crystallization behavior from the solution, the precursor so lution was first dried at 90 � C for 24 h and then further dried at 200 � C for 1 h. The resulting dried powder was calcinated at a temperature of 400, 500, 600 or 700 � C for 2 h in air at 10 � C/min to obtain powder samples. For the fabrication of film samples, a c-axis oriented ZnO seed layer was first fabricated on Pt(111)/Ti/SiO2/Si(100) substrates (hereafter abbreviated as Pt/Si substrates) by the following CSD method [21,22]. Zn (CH3COO)2⋅2H2O was dissolved in a mixed solvent of 2-MOE (9.70 mL) and MEA (0.30 mL) by stirring at room temperature to obtain a trans parent ZnO solution with a concentration of 0.5 M. The solution was dropped on a Pt/Si substrate (25 � 25 mm2) and spin-coated at 1000 rpm for 10 s and then 3000 rpm for 50 s. The substrate was then immediately subjected to heat treatment at 400 � C for 1 h in air to crystalize ZnO. Next, the ZnO–LCMO precursor solution was dropped on the ZnO-seeded Pt/Si substrate, and spin-coated under the same condition with the ZnO solu tion. The resulting film was dried at 150 � C for 1 min and then heated at 700 � C for 10 min in air. After repeating the coating of the ZnO–LCMO solution followed by the drying and heating 3 or 20 times, the films were lastly heated at 700 � C for 1 h in air. Hereafter a film sample fabricated by repeating the coating cycle n times will be referred to as the ZnO–LCMO (n) film (where, n ¼ 3 or 20). To study the effect of the ZnO seed layer on the film structure, a ZnO–LCMO (3) film was fabricated also on a bare Pt/Si substrate without the seed layer. Thermogravimetric–differential thermal analysis (TG–DTA) of the dried powder prepared from the ZnO–LCMO precursor solution was carried out from room temperature to 800 � C under flowing air using a SHIMADZU DTG-60 instrument. The crystal structure of the powder and film samples was identified by X-ray diffraction (XRD; Bruker AXS D8) analysis using Cu Kα radiation. The microstructure of the film samples was observed by field emission scanning electron microscopy (FE-SEM; JSM-7600F, JEOL). To investigate the composite structure of ZnO and LCMO, the ZnO–LCMO (20) film was immerged in glacial acetic acid (99.7%, Wako) for 5 min and then observed by FE-SEM. A magnetization (M)–Magnetic field (H) curve at a temperature (T) of 10 K and a field cooling M–H curve at H ¼ 1 kOe of the ZnO–LCMO (20) film were measured by a Quantum Design MPMS SQUID magnetometer. For the measurement of the converse ME response, an Au top electrode with an area of approximately 4 � 6 mm2 was sputtered on the surface of a small piece of the ZnO–LCMO (20) film (approximately 5 � 10 mm2 in area), and the bare area of the film was chemically etched by an acidic hydrogen peroxide solution (H2O: HCl: H2O2 ¼ 8 : 4: 1) to contact the Pt bottom electrode. Two Cu wires are connected to the top and bottom electrodes and an external DC voltage was applied between the electrodes using a Keithley 2400 source meter. M–H curves of the film at T ¼ 10 K were measured by the SQUID magnetometer under applying electric fields of 0, 100, and 0 kV/cm in turn. Considering the sample size for the magnetic and ME measurements, the experimental error in the measured magne tization values is assumed to be less than 3% [23,24]. 3. Results and discussion 3.1. Crystallization behavior A purple homogeneous ZnO–LCMO precursor coating solution with the total metal ion concentration of 0.6 M was obtained by using the mixed solvent composed of 2-MOE (9.64 mL) and MEA (0.36 mL). We found that solvents with a lower MEA/2-MOE ratio could not completely dissolve the Mn source. Fig. 1 shows the TG–DTA curves of the dried powder prepared from the ZnO–LCMO precursor solution. Two main weight losses with strong exothermic peaks are observed at 277 � C and 393 � C. The overall weight loss is 43%. It is known that Zn (CH3COO)2⋅2H2O starts to dehydrate at about 50 � C and then de composes in to ZnO at 270 � C upon heating in air [25,26]. The first weight loss at 277 � C is thus likely to be attributed to the thermal decomposition of dehydrated zinc acetate or its complex with MEA. The second weight loss at 393 � C would be due to the combustion of the
2. Experimental Zn(CH3COO)2⋅2H2O (99.9%, Wako Pure Chemical Industries; 5 mmol), Mn(CH3COO)2⋅4H2O (99.9%, Wako; 0.25 mmol), La (NO3)3⋅6H2O (99.9%, Wako; 0.50 mmol), and Co(CH3COO)2⋅4H2O (99.0%, Wako; 0.25 mmol) were dissolved in a mixed solvent of 2-methox yethanol (2-MOE; 9.64 mL) and monoethanolamine (MEA; 0.34 mL) by stirring at 60 � C for 1 h to obtain a ZnO–LCMO precursor coating solution. The total concentration of the metal ions in the solution was 0.6 M and the mole fraction of LCMO, namely 100(%) � LCMO/(ZnO þ LCMO), was 5%. 2
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The above observation shows that the presence of La, Co, Mn ions in the precursor does not prevent the crystallization of ZnO. The powder heated at 500 � C also shows the diffraction peaks only from the ZnO phase. In the powders heated at 400 or 500 � C, the residual La, Co, Mn ions should exist as a form of an amorphous oxide rather than acetates or nitrates because no significant weight loss is observed above 400 � C as shown in Fig. 1. When the precursor is heated at 600 or 700 � C, additional peaks appear at around 2θ ¼ 32.7, 39.0, and 46.8� , which can be indexed to (112), (202), and (004) planes of a pseudo tetragonal lattice of LCMO with the double-perovskite-type structure [15]. Any other peaks due to the im purity phase is not observed in the XRD patterns. This result shows that the crystallization of LCMO in the powder occurs at a temperature be tween 500 and 600 � C, which is somewhat inconsistent with the result of DTA where the small exothermic peak without weight loss is found below 500 � C. This possibly arises from the difference in the atmosphere con ditions during the TG–DTA measurement and the calcination. The TG-DTA curves were corrected with flowing air, whereas the calcination was performed in a closed box furnace. The oxygen partial pressure during the calcination should be thus lowered due to the combustion of organic species. The lowered oxygen partial pressure delays the oxidation of Mn2þ into Mn4þ difficult, leading to a higher crystallization temper ature of LCMO compared to the TG-DTA measurement. Therefore, it is difficult to determine the exact crystallization temperature of LCMO from the TG-DTA and calcination experiments, but nevertheless, the results are clearly showing that there is a significant difference between the crys tallization temperatures of the ZnO and LCMO phases. Such a large gap in the crystallization temperatures is a key to obtain ZnO–LCMO composite films with a controlled nanostructure.
393 °C Exo.
0
DTA
-20
277 °C
-30
400
500
600
-40 -50
0
100
200
300
400
500
600
700
Endo.
Weight (%)
-10
800
Temperature/°C Fig. 1. TG–DTA curves of the ZnO–LCMO dried powder. The inset is an enlarged image of the DTA curve between 400 and 650 � C.
residual organic compounds. The nitrate ion originating from the La source can be involved in the combustion process as an internal oxidizer. From the presented thermal behavior of the ZnO–LCMO dried powder, it is found that the decomposition of the powder is completed at a rela tively low temperature below 400 � C. It should be noted here that a smaller broad exothermic peak without a corresponding weight change is also found in a temperature range between 460 � C and 500 � C, as shown in the inset of Fig. 1, which is presumably due to the crystalli zation of LCMO as discussed below. The dried powder was calcined in air at a temperature of 400, 500, 600, or 700 � C for 2 h to study its crystallization behavior. Fig. 2 shows the XRD patterns of the resultant powders. The powder heated at 400 � C shows clear diffraction peaks at 2θ ¼ 31.2, 34.6, 35.9, and 47.3� . All these peaks can be indexed to the wurtzite-type ZnO phase (ICDD 089–0510). This indicates that the crystallization of ZnO accompanies the decom position and combustion of organic species observed below 400 � C in the TG-DTA curves. It has been reported that the crystallization of ZnO from a precursor solution without any other cations starts at round 400 � C [22].
3.2. Film structures Based on the crystallization behavior studied above, the ZnO–LCMO precursor solution was spin-coated on a ZnO-seeded Pt/Si substrate and
Fig. 3. (a) An XRD pattern of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C and (b) an enlarged image of (a) between 30 and 37� .
Fig. 2. XRD patterns of the powder samples prepared by heating the ZnO–LCMO dired powder at a temperature of 400, 500, 600, or 700 � C for 2 h. 3
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heated at 700 � C to obtain ZnO–LCMO nanocomposite films. The XRD pattern of the ZnO–LCMO (20) film is presented in Fig. 3. All the diffraction peaks observed in the 2θ range between 10 and 60� can be attributed to ZnO (ICDD 089–0510), LCMO [15], or the substrate, confirming that a composite between ZnO and LCMO is formed and no significant reaction among these phases and the substrate occurs. It is seen in the magnified image around the strongest ZnO (002) peak that the ZnO (100) and (101) peaks are very weak, showing that the ZnO phase in the film has a strong c-axis orientation. The position of the ZnO (002) peak is found at 2θ ¼ 34.61� , which is slightly higher than the reference data (ICDD 089–0510; 2θ ¼ 34.43� ). This suggests the possi bility of the partial substitution of Zn2þ ions (60 p.m.) by the smaller Co2þ ions (56 p.m.) [27]. Fig. 4(a) and (b) are FE-SEM images of the surface and cross-section of the ZnO–LCMO (20) film, respectively. A characteristic microstruc ture is observed on the surface of the film, where small nanograins about 50 nm are filling up intergranular gaps of larger grains about 200 nm. The cross-sectional image shows a column-like dense structure with a thickness about 600 nm. With an aim to get further insight into the film structure, the film was immersed in glacial acetic acid and then observed by FE-SEM again because ZnO easily dissolves in acetic acid whereas LCMO does not. In the surface FE-SEM image shown in Fig. 4(c), it is seen that the large grains have disappeared by the acetic-acid treatment while the small nanograins have remained. This result clearly demon strates that the larger grains seen in the as-deposited film are ZnO whereas the nanograins are LCMO. It is also observed in the acetic-acid treated film that there are many vertical holes. These holes should be formed by the dissolution of column-like ZnO grains. Based on these observations together with the XRD data, the ZnO–LCMO (20) film is found to have a characteristic vertical nanostructure as illustrated in
Fig. 4(d), where the nanograins of LCMO fills the gaps between the c-axis oriented ZnO columns. It should be noted that the interface between the ZnO seed layer and the film is not clearly observed in the cross-sectional image (Fig. 4(b)) because the ZnO columns were homoepitaxially grown on the seed layer. To reveal the role of the ZnO seed layer on the formation of the vertical nanostructure, structures of the ZnO–LCMO (3) films fabricated on bare and ZnO-seeded Pt/Si substrates were compared. The coating cycle was repeated for only 3 times in order to observe the early stage of the film grows where the effect of the seed layer is expected to be sig nificance. The XRD patterns and the surface FE-SEM images of the films are shown in Fig. 5. The film fabricated on the ZnO-seeded Pt/Si sub strate shows a strong ZnO (002) peak and its microstructure is similar to that of the ZnO–LCMO (20) film (Fig. 4(a)). This confirms that the ZnO columns were directly grown on the ZnO seed layer. On the other hand, the intensity of the ZnO (002) peak of the film fabricated on the bare substrate is much smaller and comparable to that of the ZnO (110) peak, indicating that the ZnO phase in this film are randomly oriented. Furthermore, no large grain attributable to the ZnO column is found on the surface of this film. Thus, it is clearly demonstrated that the ZnO seed layer plays a critical role on the growth of c-axis oriented ZnO columns. From the observations shown above, the mechanism underlying the formation of the ZnO–LCMO nanocomposite film can be understood as follows. Upon heating after the spin-coating process, heterogeneous nucleation of ZnO first occurs on the c-axis oriented ZnO seed layer at a temperature less than 400 � C. The ZnO crystal nuclei also have the c-axis orientation to match the crystal lattice with the seed layer. The ZnO nuclei grow mainly in the direction toward the film surface (along the caxis) to form column-like grains during further heated up, and the amorphous oxide containing the residual ions of La, Mn, and Co is
Fig. 4. Microstructure of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C: (a) Surface and (b) cross-sectional FE-SEM images of the asdeposited film, (c) a surface FE-SEM image of the film after immerged in glacial acetic acid for 5 min, and (d) a schematic illustration of the film structure. 4
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Fig. 5. (a, b) XRD patterns and (c, d) FE-SEM images of the ZnO–LCMO (3) film prepared on (a, c) ZnO-seeded and (b, d) bare Pt/Si substrates at 700 � C.
Magnetization, M/emu·cm-3
concentrated in regions between the ZnO columns. Then, the homoge neous nucleation of the LCMO phase occurs in the amorphous oxide at a temperature between 450 and 600 � C, leading to the formation of the LCMO nanograins between the ZnO columns. The similar phase sepa ration process, which starts from the heterogeneous nucleation of ZnO on the ZnO columns, is repeated in the subsequent coating cycles, resulting in the formation of the vertically aligned nanocomposite film composed of the ZnO columns and the LCMO nanograins. 3.3. Magnetic and magnetoelectric properties Fig. 6(a) shows the M–H curve of the ZnO–LCMO (20) film measured at 10 K. A clear hysteresis with a saturation magnetization (Ms) of 58 emu/cm3 is observed. The Ms value of LCMO at 5 K is reported to range between 3.36 and 4.85 μB/f.u., which is corresponding to 260–380 emu/ cm3, depending on fabrication conditions such as firing temperature and atmosphere [15]. Since the mole fraction of LCMO in the composite film is 5%, the volume fraction of LCMO is calculated as 21% using the unit cell volumes of ZnO and LCMO (47.7 Å3 and 236 Å3, respectively, where both unit cells include two formula units). Considering this volume fraction of LCMO, the observed Ms value of the ZnO–LCMO (20) film is well consistent with the reported values. The M–T curve of the film measured at 1 kOe upon cooling is shown in Fig. 6(b). The M–T curve shows that the Currie temperature (TC), above which the ferromagne tism is vanished, of the film is at round 200 K. It has been reported that TC of LCMO is also varied in a range between 130 and 230 K depending on the preparation condition [15]. The observed TC of the ZnO–LCMO nanocomposite film is within the reported TC range. It should be noted that Co-doped ZnO also shows ferromagnetic behaviors [28,29], and the XRD pattern of the ZnO–LCMO (20) film indicates a possibility of the partial substitution of Co2þ for Zn2þ in the ZnO phase. However, TC of Co-doped ZnO is reported much higher than room temperature [29], and its Ms should be much smaller than the observation in this study. It can be thus concluded that the ZnO–LCMO (20) film has ferromagnetic properties originating from the LCMO phase, which are not degraded by the coexistence of the ZnO phase. To observe the converse ME response of the ZnO–LCMO (20) film, an M–H curve under an out-of-plane external electric field was measured. Fig. 7 shows a part of M–H hysteresis curves measured with and without applying an electric field of 100 kV/cm. It is observed that the
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showed ferromagnetic properties originating from the LCMO phase as well as the magnetoelectric response to the external electric field. The reduction of magnetization was 9% at an applied electric field of 100 kV/cm. The CSD method developed in this study enables low-cost and large-area fabrication of ZnO-based VAN films on Si substrates.
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Acknowledgement
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The authors wish to thank Prof. T. Sato of Keio university for his help in the ME measurement. M. S. thanks the Kato Foundation for Promotion of Science for its financial support towards this work.
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[1] W. Eerenstein, N.D. Mathur, J.F. Scott, Multiferroic and magnetoelectric materials, Nature 442 (2006) 759–765, https://doi.org/10.1038/nature05023. [2] J. Wang, J.B. Neaton, H. Zheng, V. Nagarajan, S.B. Ogale, B. Liu, D. Viehland, V. Vaithyanathan, D.G. Schlom, U. V Waghmare, N.A. Spaldin, K.M. Rabe, M. Wuttig, R. Ramesh, Epitaxial BiFeO3 multiferroic thin film heterostructures, Science 299 (2003) 1719–1722, https://doi.org/10.1126/science.1080615. [3] Z.J. Huang, Y. Cao, Y.Y. Sun, Y.Y. Xue, C.W. Chu, Coupling between the ferroelectric and antiferromagnetic orders in YMnO3, Phys. Rev. B 56 (1997) 2623–2626, https://doi.org/10.1103/PhysRevB.56.2623. [4] J. Ma, J. Hu, Z. Li, C.W. Nan, Recent progress in multiferroic magnetoelectric composites: from bulk to thin films, Adv. Mater. 23 (2011) 1062–1087, https://doi. org/10.1002/adma.201003636. [5] C.-W. Nan, G. Liu, Y. Lin, H. Chen, Magnetic-field-induced electric polarization in multiferroic nanostructures, Phys. Rev. Lett. 94 (2005) 197203, https://doi.org/ 10.1103/PhysRevLett.94.197203. [6] H. Zheng, J. Wang, L. Mohaddes-Ardabili, M. Wuttig, L. Salamanca-Riba, D. G. Schlom, R. Ramesh, Three-dimensional heteroepitaxy in self-assembled BaTiO3–CoFe2O4 nanostructures, Appl, Phys. Lett. 85 (2004) 2035–2037, https:// doi.org/10.1063/1.1786653. [7] N. Dix, R. Muralidharan, J. Rebled, S. Estrad� e, F. Peir� o, M. Varela, J. Fontcuberta, F. S� anchez, Selectable spontaneous polarization direction and magnetic anisotropy in BiFeO3 CoFe2O4 Epitaxial Nanostructures, ACS Nano 4 (2010) 4955–4961, https://doi.org/10.1021/nn101546r. [8] R. Wu, A. Kursumovic, X. Gao, C. Yun, M.E. Vickers, H. Wang, S. Cho, J. L. MacManus-Driscoll, Design of a vertical composite thin film system with ultralow leakage to yield large converse magnetoelectric Effect, ACS Appl. Mater. Interfaces 10 (2018) 18237–18245, https://doi.org/10.1021/acsami.8b03837. [9] L. Yan, Z. Wang, Z. Xing, J. Li, D. Viehland, Magnetoelectric and multiferroic properties of variously oriented epitaxial BiFeO3–CoFe2O4 nanostructured thin films, J. Appl. Phys. 107 (2015), https://doi.org/10.1063/1.3359650, 064106. [10] C. Zhang, S.C. Huberman, S. Ning, J. Pelliciari, R.A. Duncan, B. Liao, S. Ojha, J. W. Freeland, K.A. Nelson, R. Comin, G. Chen, C.A. Ross, Thermal conductivity in self-assembled CoFe2O4/BiFeO3 vertical nanocomposite films, Appl. Phys. Lett. 113 (2018) 1–6, https://doi.org/10.1063/1.5049176. [11] T.C. Kim, S.H. Lee, H.K. Jung, H. Lee, J.H. Mun, I. Oh, D.H. Kim, Magnetic property modulation of sputter-grown BiFeO3–CoFe2O4 nanocomposite thin films, Ceram. Int. 45 (2019) 12182–12188, https://doi.org/10.1016/j.ceramint.2019.03.122. [12] H. Ryu, P. Murugavel, J.H. Lee, S.C. Chae, T.W. Noh, Y.S. Oh, H.J. Kim, K.H. Kim, J.H. Jang, M. Kim, C. Bae, J.-G. Park, Magnetoelectric effects of nanoparticulate Pb (Zr0.52Ti0.48)O3–NiFe2O4 composite films, Appl. Phys. Lett. 89 (2006) 102907, https://doi.org/10.1063/1.2338766. [13] T. Fix, E.-M. Choi, J.W.A. Robinson, S.B. Lee, A. Chen, B. Prasad, H. Wang, M. G. Blamire, J.L. MacManus-Driscoll, Electric-field control of ferromagnetism in a nanocomposite via a ZnO phase, Nano Lett. 13 (2013) 5886–5890, https://doi.org/ 10.1021/nl402775h. [14] C.L. Bull, D. Gleeson, K.S. Knight, Determination of B-site ordering and structural transformations in the mixed transition metal perovskites La2CoMnO6 and La2NiMnO6, J. Phys. Condens. Matter 15 (2003) 4927–4936, https://doi.org/ 10.1088/0953-8984/15/29/304. [15] R. Dass, J. Goodenough, Multiple magnetic phases of La2CoMnO6 δ (0�δ�0.05), Phys. Rev. B 67 (2003), https://doi.org/10.1103/PhysRevB.67.014401, 014401. [16] M.P. Singh, K.D. Truong, P. Fournier, Magnetodielectric effect in double perovskite La2CoMnO6 thin films, Appl. Phys. Lett. 91 (2007), https://doi.org/10.1063/ 1.2762292, 042504. [17] A. Chen, Z. Bi, Q. Jia, J.L. MacManus-Driscoll, H. Wang, Microstructure, vertical strain control and tunable functionalities in self-assembled, vertically aligned nanocomposite thin films, Acta Mater. 61 (2013) 2783–2792, https://doi.org/ 10.1016/j.actamat.2012.09.072. [18] R.W. Schwartz, T. Schneller, R. Waser, Chemical solution deposition of electronic oxide films, Compt. Rendus Chem. 7 (2004) 433–461, https://doi.org/10.1016/j. crci.2004.01.007. [19] S.K. Singh, Y.K. Kim, H. Funakubo, H. Ishiwara, Epitaxial BiFeO3 thin films fabricated by chemical solution deposition, Appl. Phys. Lett. 88 (2006) 86–89, https://doi.org/10.1063/1.2196477. [20] Q. Zhang, N. Valanoor, O, Standard, Chemical solution deposition derived (001)oriented epitaxial BiFeO3 thin films with robust ferroelectric properties using stoichiometric precursors, J. Appl. Phys. 116 (2014), https://doi.org/10.1063/ 1.4891311.
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Magnetic field, H/kOe Fig. 7. Magnetization–magnetic field curves of the ZnO–LCMO (20) film fabricated on a ZnO-seeded Pt/Si substrate at 700 � C. The curves were measured before, under, and after application of an electric field of 100 kV/cm at 10 K. The inset shows the variation of the saturation magnetization at 45 kOe.
magnetization at high magnetic fields is slightly decreased by applying the electric field. The magnetization is switched back to the initial state when the electric field is removed. On the other hand, the application of the electric filed is found to have almost no impact on the coercive force (Hc). This implies that the reduction of magnetization is not due to the Joule heating of the film during the measurement because Hc should be significantly decreased by heating. The reduction of Ms is 9% at 45 kOe, which is much smaller than the reported value for the ZnO–LCMO nanocomposite film prepared by the PLD method (60% reduction at an electric field of 305 kV/cm) [13]. There are some possible reasons for the smaller ME response of the sample obtained in this study. The first reason is related to the magnitude of the applied electric field. For the ZnO–LCMO nanocomposite film by the PLD method, it has been re ported that the reduction of magnetization depends nonlinearly on the applied electric field and becomes significant when the field exceeds 305 kV/cm [13]. Indeed, the reduction is only 14% at 289 kV/cm [13], which is comparable to the observation of this study. We couldn’t apply electric fields higher than 100 kV/cm because of the electrical break down of the sample. The poorer electrical break down strength of the sample would be related to the somewhat porous microstructure in the LCMO region as seen in Fig. 4(a). The second possible reason is the interface structure between ZnO and LCMO. Because the LCMO grains are smaller than the gaps between the ZnO columns, the majority of the LCMO grains are out of contact with the ZnO columns. These LCMO grains should be hidden from the ZnO/LCMO interfaces where the electron trap/detrap occurs. Therefore, we consider that the further modification of the microstructure of the LCMO region is needed to improve the ME response of the film. Nevertheless, the result shown here is of great importance for demonstrating that the CSD method is effec tive for fabricating magnetoelectric ZnO–LCMO VAN films on Si sub strates. It is also believed that the CSD method developed in this study can be applied to other ZnO-based vertical nanocomposite systems for various applications such as the ZnO–La0.7Sr0.3MnO3 system showing a strong magnetoresistance effect [30,31]. 4. Conclusion The ZnO–LCMO vertical nanocomposite film was deposited on the Pt/Si substrate using the single precursor solution. It was found that the use of the ZnO seed layer and the large difference in the crystallization temperatures of ZnO and LCMO were important for obtaining the ver tical nanocomposite structure. The ZnO–LCMO nanocomposite film 6
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Materials Chemistry and Physics 236 (2019) 121762 [27] R.D. Shannon, Revised effective ionic radii and systematic studies of interatomic distances in halides and chalcogenides, Acta Crystallogr. A 32 (1976) 751–767, https://doi.org/10.1107/S0567739476001551. [28] H.-J. Lee, S.-Y. Jeong, C.R. Cho, C.H. Park, Study of diluted magnetic semiconductor: Co-doped ZnO, Appl. Phys. Lett. 81 (2002) 4020–4022, https:// doi.org/10.1063/1.1517405. [29] C. Song, K.W. Geng, F. Zeng, X.B. Wang, Y.X. Shen, F. Pan, Y.N. Xie, T. Liu, H. T. Zhou, Z. Fan, Giant magnetic moment in an anomalous ferromagnetic insulator: Co-doped ZnO, Phys. Rev. B 73 (2006), https://doi.org/10.1103/ PhysRevB.73.024405, 024405. [30] J.L. MacManus-Driscoll, P. Zerrer, H. Wang, H. Yang, J. Yoon, A. Fouchet, R. Yu, M.G. Blamire, Q. Jia, Strain control and spontaneous phase ordering in vertical nanocomposite heteroepitaxial thin films, Nat. Mater. 7 (2008) 314–320, https:// doi.org/10.1038/nmat2124. [31] A. Chen, Z. Bi, C.-F. Tsai, J. Lee, Q. Su, X. Zhang, Q. Jia, J.L. MacManus-Driscoll, H. Wang, Tunable low-field magnetoresistance in (La0.7Sr0.3MnO3)0.5:(ZnO)0.5 selfassembled vertically aligned nanocomposite thin films, Adv. Funct. Mater. 21 (2011) 2423–2429, https://doi.org/10.1002/adfm.201002746.
[21] S. Fujihara, C. Sasaki, T. Kimura, Effects of Li and Mg doping on microstructure and properties of sol-gel ZnO thin films, J. Eur. Ceram. Soc. 21 (2001) 2109–2112, https://doi.org/10.1016/S0955-2219(01)00182-0. [22] S. Fujihara, C. Sasaki, T. Kimura, Crystallization behavior and origin of c-axis orientation in sol–gel-derived ZnO:Li thin films on glass substrates, Appl. Surf. Sci. 180 (2001) 341–350, https://doi.org/10.1016/S0169-4332(01)00367-1. [23] P. Stamenov, J.M.D. Coey, Sample size, position, and structure effects on magnetization measurements using second-order gradiometer pickup coils, Rev. Sci. Instrum. 77 (2006), https://doi.org/10.1063/1.2149190, 015106. [24] M. Sawicki, W. Stefanowicz, A. Ney, Sensitive SQUID magnetometry for studying nanomagnetism, Semicond. Sci. Technol. 26 (2011), https://doi.org/10.1088/ 0268-1242/26/6/064006, 064006. [25] C.-C. Lin, Y.-Y. Li, Synthesis of ZnO nanowires by thermal decomposition of zinc acetate dihydrate, Mater. Chem. Phys. 113 (2009) 334–337, https://doi.org/ 10.1016/j.matchemphys.2008.07.070. [26] T. Arii, A. Kishi, The effect of humidity on thermal process of zinc acetate, Thermochim. Acta 400 (2003) 175–185, https://doi.org/10.1016/S0040-6031(02) 00487-2.
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