Accepted Manuscript Defect induced room-temperature ferromagnetism and enhanced dielectric property in nanocrystalline ZnO co-doped with Tb and Co Santanu Das, Atul Bandyopadhyay, Sukhen Das, Dipankar Das, Soumyaditya Sutradhar PII:
S0925-8388(17)33471-0
DOI:
10.1016/j.jallcom.2017.10.057
Reference:
JALCOM 43454
To appear in:
Journal of Alloys and Compounds
Received Date: 18 July 2017 Revised Date:
7 October 2017
Accepted Date: 9 October 2017
Please cite this article as: S. Das, A. Bandyopadhyay, S. Das, D. Das, S. Sutradhar, Defect induced room-temperature ferromagnetism and enhanced dielectric property in nanocrystalline ZnO co-doped with Tb and Co, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2017.10.057. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Defect induced room-temperature ferromagnetism and enhanced dielectric property in nanocrystalline ZnO co-doped with Tb and Co
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Department of Physics, Amity University, Rajarhat, Kolkata-700156, West Bengal, India
Department of Physics, Jadavpur University, Jadavpur, Kolkata-700032, West Bengal, India 3
Department of Physics, University of Gour Banga, Malda, West Bengal-732103, India
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UGC-DAE Consortium for Scientific Research, III/LB-8, Kolkata-700098, West Bengal, India
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Santanu Das1,2, Atul Bandyopadhyay3, Sukhen Das2, Dipankar Das4, Soumyaditya Sutradhar1*
Microstructural, magnetic and dielectric properties of hydrothermally prepared dual ions doped (Tb and Co) zinc oxide (Tb0.05Co0.05Zn0.90O) nanoparticles are investigated in the present work. The polycrystalline single phase of undoped and doped ZnO nanoparticles with hexagonal wurtzite structure is confirmed by X-ray diffraction pattern. Raman and Photoluminescence
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spectroscopy studies are recorded at room temperature and the measurements ruled out the presence of any impurity phase in both undoped and Tb, Co co-doped ZnO samples. Also, the Photoluminescence emission spectra for the entire samples confirm the presence of various
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defect centers inside co-doped ZnO wurtzite nanostructure. Magnetic and dielectric properties of
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co-doped ZnO system are enhanced remarkably as compared to undoped as well as some monodoped ZnO systems. The room temperature ferromagnetism with high coercive field and magnetization has been observed in the magnetic measurement and the outcome of the measurement will be fruitful for future spintronics device applications. The dielectric behavior has been explained on the basis of Maxwell-Wagner-Sillars polarization phenomenon. High value of dielectric constant was observed due to the presence of grain boundary defects and various others defect centers at the surface of the hexagonal rod like doped ZnO nanostructures.
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Keywords: Nanoparticles, diluted magnetic semiconductor, magnetic properties, bound magnetic polaron, space charge polarization. *Corresponding author email:
[email protected] (Dr. Soumyaditya Sutradhar), Fax: 1800
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1. INTRODUCTION In last few years, dilute magnetic semiconductors (DMSs) have drawn considerable attention due to their advanced technological applications in spintronics [1-2]. Room temperature
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ferromagnetism (RTFM) and high magnetic moment by the addition of a small fraction of magnetic ions inside the structure of the host semiconducting materials is the primary criteria for a material to be used as spintronic material. In this regard, ZnO is one of the most promising host
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semiconducting materials for spintronic applications with direct band gap energy (3.37 eV) and large exciton binding energy (60 meV). Various works have been reported in recent time on
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RTFM where a variety of transition metal (TM) ions (Mn2+, Co2+, Ni2+ etc.) doped ZnO systems have been investigated. In these reports the ferromagnetic contributions of the doped materials at RT have been appeared due to the presence of the partially filled 3d shell [3-5]. In comparison with TM dopants, rare earth (RE) dopants may offer larger magnetic moment due to the presence
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of partially filled 4f orbital and as a consequence it generates enhanced ferromagnetic ordering at RT in the doped ZnO semiconductors. The individual role of 3d ions (TM) and 4f ions (RE) for ZnO based dilute magnetic semiconductors (DMSs) has been investigated in previous literatures
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[6-8]. Several theoretical models have been developed so far to explain the coupling mechanism for RTFM in this type of materials [9-11]. To date, the most popular and effective mechanisms
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are the carrier mediated indirect exchange interaction and the bound magnetic polaron (BMP) model in connection with the intrinsic defects related to vacancies or interstitials [12-14]. The presence of such defects is responsible for the enhancement of the s-f/d coupling. The co-doped samples will be considered for special interest in near future due to the strong hybridization of the local 4f electrons with the ‘d’ conduction electrons. Hence, another new area of research has opened up for further exploration on ZnO based DMS system. S. Das et al. has also reported
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enhanced RTFM behavior in TM (Co) and RE (Gd) co-doped ZnO nanoparticles [15]. However, the paramagnetic behavior from ZnO co-doped with Gd/Sm and Mn also reported by J. Das et al. also raises some burning questions on RTFM possibility of co-doped system [16]. Therefore,
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synthesis process and proper choice of doping elements in co-doped nanocomposite are very much crucial to achieve RTFM. In addition, understanding of underlying physics behind the codoping is also very important. In this direction, we report the RTFM with enhanced
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magnetization and coercivity in the Tb and Co co-doped ZnO nanoparticles. 2. EXPERIMENTAL DETAILS
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2.1. Material and method
Terbium (Tb) and Cobalt (Co) ions co-doped ZnO nanoparticles were prepared by the hydrothermal method. Zinc acetate dehydrate Zn(O2CCH3)2·2H2O (Sigma Aldrich, 99%), cobalt acetate tetrahydrate Co(O2CCH3)2·4H2O, terbium nitrate pentahydrate Tb(NO3)3·5H2O (Sigma
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Aldrich, 99%) and ammonium hydroxide NH4OH (concentration ~ 25%) were taken with analytical grade and without any further purification. All the constituent salts were dissolved in milli-Q water (resistivity value 16-17 MΩ.cm at 25 oC, Heal force ASTM Type-I Easy Series
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model) with proper stoichiometric ratio. The mixed salts solution was then taken in a beaker and the beaker was placed over magnetic stirrer and the solution was stirring for 2 hrs at room
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temperature (RT) to make a uniform distribution of the ions. After that NH4OH solution was added dropwise in the salt solution and the final pH of the solution was kept at ~ 9. The asprepared particles were settled down at the bottom of the beaker which were collected properly and washed several times by using milli-Q water and ethyl alcohol. Washed precipitates were taken in a Teflon lined autoclave which was then placed in an oven at 160 oC for 24 hrs. After 24 hrs the solid precipitates at the bottom of the Teflon jacket were collected, washed and dried at
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room temperature (RT). Finally, dried powders of doped ZnO were sintered at 500 oC temperatures for 2 hrs in vacuum atmosphere and are named as COZ (Co0.05Zn0.95O), TBZ (Tb0.05Zn0.95O) and TBCOZ (Tb0.05Co0.05Zn0.90O), respectively. This subsequent heat treatment
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of all the as prepared samples at 500 oC is required to eliminate unreacted precursor waste from the final samples and to achieve better crystallographic phase of the samples. 2.2. Characterization techniques
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The XRD of all the samples were recorded using powder X-ray diffractometer, Model D8, BRUKER AXS, using Cu Kα radiation (λ = 1.5405 Å) in the range of 2θ from 20 to 80o. The
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RT Raman spectroscopy study was employed by using Newport RS 2000TM. The photoluminescence spectroscopy study was done by using a spectroflurometer, Perkin Elmer, Germany with an excitation wavelength (λex) of 300 nm. The various magnetic properties were recorded using SQUID magnetometer (MPMS XL 7, Quantum Design), where the maximum
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applied magnetic field was 6 T. Dielectric measurements were investigated using Agilent 4294A Precision Impedance Analyzer.
3. RESULT AND DISCUSSION
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3.1. XRD analysis
X-ray diffraction pattern of all the present samples has been given in Fig. 1. All the peaks
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have been assigned with the help of the JCPDS (file no. 36-1451). All the peaks in the XRD pattern of Fig. 1 are matched very well with that of the desired phase of the polycrystalline single phase of ZnO hexagonal wurtzite structure (space group P63mc). Also, no other peaks corresponding to any other binary or ternary phases are observed in the XRD pattern in any of the sample indicating the successful incorporation of Tb3+ and Co2+ in the lattice sites of Zn2+ in the matrix of ZnO nanoparticles. Thus, it is quite clear from the given Fig. 1 that the individual
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doping of the Tb3+ and Co2+ ions as well as the simultaneous doping of both and Tb3+ and Co2+ ions even upto 5% (TBCOZ) inside the structure of ZnO nanoparticles does not produce any unwanted phase except the desired phase of ZnO wurtzite lattice structure.
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Any impurity phase in the present samples with concentration of dopants below the detection limit of XRD is hardly contribute any magnetic moment to RTFM in the samples, as none of these possible binary or ternary phases are ferromagnetic in nature at RT. The significant
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change in the intensity of diffraction peak and also the corresponding change in full width at half maximum (FWHM) with the proper selection of the dopant cations suggest the nanocrystalline
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nature of the samples. The average crystallite size, determined from the broadening (FWHM) of diffraction peaks using Scherer formula, are found to be 49.8192, 43.6179, 35.9579 and 40.9237 nm for ZnO, COZ, TBZ and TBCOZ respectively.
The incorporation of dopant ions can also be verified by the change in lattice parameters
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and the cell volume as a function of doping ions concentration. The lattice parameters “a” and the cell volumes for different doped ZnO samples are found to change in a typical manner with the selection of the dopants and the values are given in Table I.
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The average nanocrystallite diameter of the COZ with respect to undoped ZnO is small and it appears due to the effect of distortion produced inside the host ZnO lattice structure causes
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by the substitution of Zn atoms by the Co atoms. The average nanocrystallite diameter of TBZ is smaller than both undoped ZnO and COZ. This large change in the average nanocrystallite diameter appears due to the substitution of large number of Zn2+ (0.74 Å) by Tb3+ (0.94 Å). This substitution induces large number of cationic vacancies in terms of electrical neutrality in the ZnO matrix and these large cationic vacancies are mainly responsible for the decrease in the average nanocrystallite diameter for TBZ. However, in case of TBCOZ the average
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nanocrystallite diameter is also less the average nanocrystallite diameter of the undoped ZnO. At the same time the value also lies in between COZ and TBZ which is in very good agreement with the values of COZ and TBZ. These substitutions actually reduce the nucleation process of the
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nanosystem and the subsequent growth rate of ZnO crystal [17]. Similar result has been reported in our earlier publication [15]. 3.2. Raman spectroscopy studies
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Fig. 2(a-d) shows the observations obtained from the Raman scattering experiment performed for entire sample ranging from 50-800 cm-1 in order to investigate the various
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microstructural informations such as, quality of crystallinity, structural disorderness, intrinsic defects and also the successful incorporation of Tb and Co ions at the sites of Zn in the ZnO lattice structure. Raman active zone-centre optical phonon modes were predicted using group theory A1+ 2E2+E1+2B1. Here B1 is the silent mode and it is not Raman active. Both A1 and E1
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modes are Raman and infrared active polar modes and therefore they can split into transverse optical (TO) and longitudinal optical (LO) phonon components. Two strong Raman bands were found nearly at 96-101 and 433-441 cm-1 for all the samples, which have been assigned to non-
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polar optical phonon modes corresponding to E2L (E2 low mode) and E2H (E2 high mode), respectively. The E2L mode is associated with the presence of oxygen vacancies and zinc
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interstitials [18-21]. Here E2H mode is related to the motion of oxygen atom in wurtzite structure of ZnO. Observed shift in the position of aforesaid modes further indicate the introduction of lattice distortion inside the wurtzite structure of doped ZnO due to the effect of doping, which was earlier explained with the help of XRD analysis. Some other regular modes were also detected which are second order vibration modes such as E2H-2L (324-336 cm-1) and A1TO (383 cm-1) respectively [22]. Sharp A1TO peak was observed nearly at 586 cm-1 for undoped ZnO and
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in doped ZnO samples this peak becomes broadened. It can be ascribed to the multiphonon scattering from structural disorders and various defects, caused due to the dopants in host lattice structures of ZnO such as oxygen vacancy (Vo), zinc interstitial (Zni) and their complex due to
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dependence of oxygen stoichiometry [23]. Therefore our Raman results consistent with XRD analysis and manifests implantation of Tb and Co ions inside wurtzite structure of ZnO nanoparticles.
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3.3. Photoluminescence study
Photoluminescence (PL) spectroscopy analysis provides quantitative and qualitative informations
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about various forms of defects inside the samples. Some factors such as preparation technique, stoichiometry and post sintering process etc. are responsible for the occurrence of the various microstructural defects. Different energy states are expected corresponding to each type of defect centre in between the valance band and conduction band of the doped ZnO semiconducting
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nanoparticles and all these are capable of inducing PL emission spectra. The Gaussian decomposed PL emission spectra of the undoped and doped ZnO nanoparticles recorded at the excitation wavelength of 300 nm are shown in Fig. 3. In Fig. 3 the UV emission bands were
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observed at 344-397 nm [(K1, L1), (K2, L2), (K3, L3) and (K4, L4)] which was formed due to the recombination of free exactions through exciton-exciton collision process [25]. In visible region,
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violet emissions ranges from 397-430 nm [M1, M2, M3 and M4] is attributed to electronic transition from zinc interstitial level (Zni) to top of the valance band [26]. Electronic transitions from interstitial zinc level (Zni) to zinc vacancy level (VZn) induced overlapping of blue emissions located near around 445-468 nm [N1, N2, N3 and N4] [27]. Observed blue-green emission nearly at 490-500 nm [O1, O2, O3 and O4] are assigned due to the surface defect in the ZnO nanoparticles corresponding to the transition between oxygen vacancy and oxygen
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interstitial defect. Yellow-green emission band has been observed nearly at 520-550 nm [P1, P2, P3 and P4] and these peaks are assigned due to the electronic transition from bottom of the conduction band to singly occupied oxygen vacancies or zinc interstitials [28]. The similar result
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has been published by Kaviyarasu et al. [24]. In this published report crystalline ZnO:CeO2 nanoplatelets have been prepared by solvothermal synthesis route and a strong green emission has been observed due to the presence of oxygen vacancies. The PL spectroscopy analysis thus
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confirms the presence of an adequate amount of oxygen vacancy defect in all the samples, which strongly support the presence of colossal magnetic moment in the Tb and Co doped ZnO
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nanoparticles. 3.4. dc magnetic studies
Figs. 4-7 show the variations of ZFC and FC magnetizations (M) as a function of temperature and the M-H curves recorded at RT and at various others low temperatures (100, 50
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and 10 K) for the samples TBZ, COZ and TBCOZ. From Fig. 4(a) it is evident that two magnetization curves are overlapped for TBZ which indicates the absence of any magnetic ordering or the presence of feeble magnetic ordering. Although, the systematic analysis of M-H
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curves of TBZ given in Fig. 5, it is found that the sample gives a clear hysteresis loop at 10 K [Fig. 5(a)]. Whereas, the M-H curves recorded at 300, 100 and at 50 K [Fig. 5(b-d)] are almost
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linear up to the maximum applied magnetic field of 6 T. The maximum magnetization (Mmax), coercive field (Hc) and remanance (Mr) of TBZ are extracted from the loop recorded at 10 K are 7.16 emu/gm, 22.5 Oe and 0.02 emu/gm, respectively. Thus, these small coercive field (Hc) and remanance even at low temperature (10 K) will make them unsuitable for any practical purpose. Our result is consistent with the previous result obtained by other researcher where they have obtained low temperature ferromagnetic loop of Tb3+ ions doped ZnO [10, 29].
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The FC and ZFC curves of COZ and TBCOZ are also shown in Figs. 4(b) and 4(c). In both the cases FC and ZFC curves bifurcate from RT which clearly demonstrates the presence of magnetic ordering/ferromagnetism (FM) at and below RT. The origin of the observed FM at RT
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could arise from a number of sources viz., intrinsic property of the doped ions, formation of some nanoscale co-related secondary phases, Co/Tb precipitation and CoO, Tb2O3 etc. However, the cause of the introduction of the ferromagnetism by the CoO phase can be easily ruled out,
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since CoO is antiferromagnetic with a Neél temperature of 293 K. Tb2O3 is also paramagnetic material down to our lowest measuring temperature. Secondly, the cause of the ferromagnetism
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observed in the samples by metallic Co is also an unlikely source of this ferromagnetism, as XRD, Raman and PL analysis show no metallic Co-clusters in the samples. Thus, FM is purely intrinsic in nature as per as the present doped ZnO sample is being concerned. It is to be mentioned here that the bifurcation is less in COZ (~ 0.062 emu/gm at 100 K) in compare to that
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of TBCOZ (~ 0.11 emu/gm at 100 K) which indicates TBCOZ poses stronger magnetic ordering than COZ. This fact is substantiated by M-H loop at RT of COZ [Fig. 6(d)] and TBCOZ [Fig. 7(d)] where saturation magnetization of TBCOZ is nearly ~ 2 times higher than COZ. The
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coercive field (Hc) and remanance (Mr) of TBCOZ is also remarkably higher than that of COZ. Hc between 350 and 4000 Oe is often desirable for the typical magnetic storage media with non-
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volatility feature [30]. Thus the observed Hc (~ 620 Oe at RT) of TBCOZ make them potential candidates for the data storage industry in future device application. Consequently, DMS system co-doped with 3d and 4f ion may be the best choice for spintronic application compared to that of mono-doped system. Some selected M-H loops of TBCOZ at different temperatures 10, 50 and 100 K are shown in Figs. 7(a), 5(b) and 5(c) respectively. From the displayed loops it is
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evident that the hysteresis curves deviate from saturation with the lowering of temperature due to increase of paramagnetic contribution at lower temperature [will be discussed later]. The reports on 3d and 4f ions co doped system are very limited hence the correct
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mechanism of intrinsic ferromagnetism is still under debate. As compared to d electrons of transition metals, f electrons of rare earth can couple strongly with the host s electrons which may lead to the possibility of electron-mediated ferromagnetism [Ruderman-Kittel-Kasuya-
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Yosida (RKKY)]. According to the RKKY theory, the presence of free carriers in the matrix has a substantial impact for the introduction of ferromagnetic behavior. In our case, the dielectric
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constant in co-doped system is appreciably high [will be discussed later]. Thus due to lack of free carriers, RKKY interaction can be ruled out for the presence of robust ferromagnetism in TBCOZ. The strong magnetization and coercivity of TBCOZ are the results of bound magnetic polarons (BMPs).
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BMP model deals with the ferromagnetism resulted by the formation of polarons of the shallow donors (Vo) with their surrounding magnetic ions. Each BMP is comprised of a polaron originating from a defect and a few magnetic dopant ions located within its strain field [31]. The
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presence of large number of defects in the sample is already substantiated by Raman and PL analysis. Polarons can initiate the long-range ferromagnetism in DMSs when several magnetic
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ions are located in the polarizing field of a polaron forms a BMP, an exchange interaction occurs between each polaron and the magnetic dopants within its radius. This interaction tends to parallel alignment of the spin of all magnetic dopants within its radius which can be expressed in the form of s-d (s-f for rare earth dopants) exchange parameters [32], |ψr| Ω J/
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(1)
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Where, S and Ω are the spin and volume of the magnetic ion respectively. By applying rigid sphere approximation for the integral part of equation (1) it can be shown that magnitude of s-f interaction is much higher than s-d interaction as per ab-initio studies of the other researchers
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[33, 34]. Thus doping of f ion will enhance the magnetization of each BMPs and leading to the
between BMPs can be described as M = M Lx + χ H
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high magnetization. According to Coey's model, the measured magnetization due to interaction
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The details of fitting and the symbols are explained in the previous paper [35, 36]. Thus net
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magnetization is the additive effect of both correlated [FM] and isolated [PM] spins. The fraction of doped 3d/4f ion which do not take part in the BMP interaction are expected to behave like an independent paramagnetic part and contribute towards the paramagnetism (isolated spin). We have analyzed the initial M-H curves of TBCOZ recorded at 300, 100, 50 and 10 K by BMP
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model and two respective fittings at 300 and 50 K are shown in Figs. 8(a) and 8(b). The value of the number of BMP/volume and paramagnetic contribution of TBCOZ are estimated from the fittings are respectively, 2.3×1012, 0.071 emu/gm (300 K), 5.3×1012, 0.227 emu/gm (100 K),
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1.1×1013, 0.339 emu/gm (50 K) and 9.0×1014, 0.552 emu/gm (10 K) whereas the number of BMP’s for COZ at RT is 1.78×1012. Thus the enhanced number of BMP’s with larger moment in
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TBCOZ compared to COZ substantiates the strong magnetization in TBCOZ. It is to be noted that M-H loop at 300 K [Fig. 7(d)], the magnetic moment at the highest measuring field of 6 T appears to saturate, because the contribution to net moment by the isolated non interacting BMPs (paramagnetic moments) is very low at RT. At lower temperature the magnetization become unsaturated, because at such low temperature the paramagnetic moments increases and give measurable magnetization. This is well reflected from the paramagnetic contribution derived
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from BMP fitting. Thus in case of TBCOZ, magnetization was remarkably enhanced compared to mono doped system and this robust magnetization is achieved by creation of defects (BMP) within the sample by co-doping.
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It is to be further mentioned here that the larger ratio of dopant cations related asymmetry in co-doped ZnO sample compared to mono doped one creates more structural/vacancy defects such as oxygen vacancies, zinc interstitials and zinc vacancies leading to the enhancement of
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electrons and holes concentration in the host matrix. This fact is also substantiated from the XRD analysis where co-doped system shows larger lattice strain (~ 9.669×10-4) as compared to
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undoped ZnO (~ 6.955×10-4). It appears due to lattice distortion by the successful incorporation of the dopants inside the structure of wurtzite ZnO. Dhakel et al. showed that the doping of rare earth ions creates oxygen vacancies in the ZnO crystalline structure which causes the decrease of the lattice parameter [37]. Such defects, particularly oxygen vacancies are responsible for the
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development of the ferromagnetic ordering [38]. In nanostructure system, most of the defects are present at the surface of nanoparticles. According to the literature, magnetic cations, carriers and defects can create bound magnetic polarons (BMPs). Thus magnetic polarons are the outcome of
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surface distortion. The surface defects causes overlapping of dopant cations as well as adjacent defects by inducing ferromagnetic coupling between dopants spin when the concentration of
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surface defects exceed the percolation threshold [39]. Thus surface defects and magnetic cationic distribution in the host lattice are the key factors for triggering magnetic ordering in the diluted magnetic semiconducting system. It is to be noted that as compare to Co-ions, Tb-ions have larger magnetic moments. Thus magnetic polarons in co-doped system are now comprised of large moments from both magnetic dopant ions (Tb and Co-ions). The BMP fitting in the present case indicates that the number of BMP is increased in co-doped system compared to mono doped
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one. Thus enhanced number of BMP with larger moment is the main reason behind the robust ferromagnetism in the present sample. It is well known that polaron in general form string like structure. These BMPs are connected with one another, they would form a chain and the
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magnetic moments between the adjacent particles are non-collinear. This kind of magnetization process can be described by the fanning mode [40]. This mechanism leads to a large value of coercive field in co-doped ZnO sample [41]. However, the increase of coercivity due to
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enhancement in the magnetic anisotropy by low symmetric surface atoms and lattice defects are also possible [42].
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3.5. Dielectric studies
The variation of real part of dielectric constant (ε′) and loss tangent (tanδ) have been depicted in Figs. 8(c) and 8(d), respectively. Maximum value of ε' were recorded for the samples COZ, TBCOZ and TBZ are 1698, 790 and 267 respectively. The observed large value of ε′ can
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be explained on the basis of electron hopping mechanism between Zn2+ and Co2+ and/or Tb3+ in the doped nanostructures, which act as permanent electric dipoles under the presence of external electric field. The observed frequency dependent dielectric behavior can be explained on the
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basis of Koop’s theory which is on the other hand depends on Maxwell-Wagner-Sillars model for the homogeneous structure [43]. According to this model, a dielectric medium is expected to
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be made of well conducting grains which are separated by poorly conducting grain boundaries. Under the influence of external electric field the free charge carriers can easily drift inside the conducting grains but they are accumulated at the grain boundaries. These accumulations of free charge carriers at the edge of the grain boundaries process large polarizability which in turn produces large polarization and high value of dielectric constant of the materials. It has been observed that ε′ decreases with the decreasing nanocrystalline diameter of the doped
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nanomaterials. Actually the surface energy of the nanoparticles improves with the decreasing nanocrystallite diameters. So, large surface energy comprises the large growth rate of the nanoparticles at a certain temperature. Thus the samples of small nanocrystallite diameter can
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produce large microstructures/grains of the material. The grain size of COZ is expected to be the lowest as compared to TBCOZ and TBZ and a gradual increment of the grain size is also expected for COZ, TBCOZ and TBZ. The dielectric constant is a function of the number of
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grains or the amount of effective grain boundaries in the sample, so the large grain structure of TBZ as compared to TBCOZ and COZ reduces the number of grains or the amount of effective
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grain boundaries in the samples and thereby it reduces the dielectric constant of the materials. Dielectric loss tangent decreases with the increase of frequency in similar manner to that of the ε′. High values for dielectric loss factor (εʹʹ) at lower frequencies are observed due to impurities, crystal defects and moisture effect within the samples.
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4. Conclusion
In summary, we are able to synthesize single phase Tb0.05Co0.05Zn0.90O using very simple cost effective hydrothermal method. No discontinuous change was found in both the ZFC and
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FC curves, confirming that the observed FM can be attributed to the single magnetic phase without any phase-separation. The high magnetization and coercivity are accredited to the
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correlation of the BMP which are generated by the oxygen vacancies present in the sample. The high magnetization and dielectric constant at RT of TBCOZ may have good prospect in magneto-electronic appliances. Acknowledgments
We gratefully acknowledge Dr. Achintya Singha of Bose Institute, Kolkata for Raman spectroscopy measurement.
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REFERENCE [1] A. S. Chambers, B. Gallagher, New J. Phys. 10 (2008) 055004. [2] T. Dietl, Nat. Mater. 9 (2010) 965.
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[3] C. J. Cong, J. H. Hong, Q. U. Liu, L. Liao, K. L. Zhang, Solid State Comm. 138 (2006) 511. [4] B. Loukya, D. S. Negi, K. Dileep, N. Kumar, J. Ghatak, R. Datta, J. Magn. Magn. Mater. 45 (2013) 159.
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[5] K. Omri, J. E. Ghoul, O. M. Lemine, M. Bououdina, B. Zhang, L. E. Mira, Superlatt. Microstruct. 60 (2013) 139.
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[6] J. Kennedy, G. V. M. Williams, P. P. Murmu, B. J. Ruck, PHYSICAL REVIEW B 88 (2013) 214423.
[7] P. P. Murmu, J. Kennedy, B. J. Ruck, S. Rubanov, Electronic Materials Letters 11 (2015) 998.
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[8] P. P. Murmu, J. Kennedy, G. V. M. Williams, B. J. Ruck, S. Rubanov, Appl. Phys. Lett. 101 (2012) 082408.
[9] M. H. Choi, T. Y. Ma, Materials Letters 62 (2008) 1835.
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[10] A. Bandyopadhyay, A. K. Deb, K. Mukhopadhyay, S. K. Roy, P. K. Chakrabarti, J. Mater. Sci. 47 (2012) 2284.
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[11] S. Singh, J. N. Divya Deepthi, B. Ramachandran, M.S. Ramachandra Rao, Materials Letters 65 (2011) 2930.
[12] M. Berciu, R. N. Bhat, Phys. Rev. Lett. 87 (2001) 108203. [13] S. D. Sarma, E. H. Hwang, A. Kaminski, Phys. Rev. B, 67 (2003) 155201. [14] J. M. D. Coey, A. P. Douvalis, C. B. Fitzgerald, M. Venkatesan, Appl. Phys. Lett. 84 (2004) 1332
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[15] S. Das, S. Das, D. Das, S. Sutradhar, J. Alloys Compds. 691 (2017) 739. [16] J. Das, D. K. Mishra, D. R. Sahu, B. K. Roul, Physica B 407 (2012) 3575.
Jayavel, T. Jimbo, J. Appl. Phys. 108 (2010) 053904.
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[17] M. Subramanian, P. Thakur, M. Tanemura, T. Hihara, V. Ganesan, T. Soga, K. H. Chae, R.
[18] P. P. Murmu, J. Kennedy, B. J. Ruck, G. V. M. Williams, A. Markwitz, S. Rubanov, A. A. Suvorova, J. Mater. Sci. 47 (2012) 1119.
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[19] R. P. Wang, G. Xu, P. Jin, Phys. Rev. B. 69 (2004) 113303.
[20] J. B. Wang, H. M. Zhong, Z. F. Li, W. Lu, J. Appl. Phys. 97 (2005) 086105.
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[21] J. D. Ye, S. L. Gu, S. M. Liu, Y. Shi, Q. Chen, Appl. Phys. Lett. 88 (2006) 101905. [22] J. Kennedy, B. Sundrakannan, R. S. Katiyar, A. Markwitz, Z. Li, W. Gao, Current Applied Physics 8 (2008) 291.
[23] C. Jayachandraiah, G. Krishnaiah, Adv. Mater. Lett. 6 (2015) 743.
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[24] K. Kaviyarasu, X. Fuku, G. T. Mola, E. Manikandan, J. Kennedy, M. Maaza, Material Letters 183 (2016) 351.
[25] J. Kennedy, P. P. Murmu, E. Manikandan, S. Y. Lee, J. Alloys Compds. 616 (2014) 614.
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[26] P. S. Venkatesh, V. Purushothaman, S. E. Muthu, S. Arumugam, V. Ramakrishnan, K. Jeganathan, K. Ramamurthid, Cryst. Eng. Comm. 14 (2012) 4713.
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[27] X. Wei, B. Man, J. Appl. Phys. 45 (2006) 8586. [28] Y. Gong, Nanoscale. Res. Lett. 2 (2007) 297. [29] G. R. Li, X. H. Lu, C. Y. Su, Y. X. Tong, J. Phys. Chem. C 112 (2008) 2927. [30] J. J. Lee, G. Z. Xing, J. B. Yi, T. Chen, M. Ionescu, S. Li, Appl. Phys. Lett. 104 (2014) 012405. [31] A. C. Durst, R. N. Bhatt, P. A. Wolff, Phys. Rev. B 65 (2002) 235205.
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[32] M. H. N. Assadi, Y. B. Zhang, S. Li, J. Phys: Condens. Matter 22 (2010) 486003. [33] K. J. Kim, Y. R. Park, Appl. Phys. Lett. 81 (2002) 1420. [34] N. J. C. Ingle, I. S. Elfimov, Phys. Rev. B. 77 (2008) 121202.
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[35] A. Bandyopadhyay, S. Sutradhar, B. J. Sarkar, A. K. Deb, P. K. Chakrabarti, Appl. Phys. Lett. 100 (2012) 252411.
[36] T. Debnath, S. Das, D. Das, S. Sutradhar, J. Alloys Compds. 696 (2017) 670.
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[37] A. A. Dakhel and M. El-Hilo, J. Appl. Phys. 107 (2010) 123905.
[38] J. M. D. Coey, M. Venkatesan, C. B. Fitzgerald, Nat. Mater. 4 (2005) 173.
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[39] R. N. Aljafi, F. Rahaman, K. N. Batoo, J. Magn. Magn. Mater. 332 (2013) 130. [40] I. S. Jacobs and C. P. Bean, Phys. Rev., 100 (1955) 1060.
[41] K. C. Verma, R. K. Kotnala, Phys. Chem. Chem. Phys. 18 (2016) 5647. [42] M. Respaud, J. M. Broto, H. Rakoto, A. R. Fert, L. Thomas, B. Barbara, M. Verelst, E.
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Snoeck, P. Lecante and A. Mosset, Phys. Rev. B 57 (1998) 2925. [43] A. S. Lanje, S. J. Sharma, R. S. Ningthoujam, J. S. Ahn, R. B. Pode, Adv Powder Technol.
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FIGURE CAPTION Fig. 1 XRD patterns of the samples (a) ZnO, (b) COZ, (c) TBZ, (d) TBCOZ. Fig. 2 Characteristic Raman vibrational modes of the samples (a) ZnO, (b) COZ, (c) TBZ, (d)
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TBCOZ.
Fig. 3 Photoluminescence spectra of the samples (a) ZnO, (b) COZ, (c) TBZ, (d) TBCOZ.
Fig. 4 Temperature dependence of FC and ZFC magnetization curves of samples (a) TBZ, (b)
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COZ, (c) TBCOZ.
Fig. 5 Hysteresis (M-H) loops of TBZ recorded at (a) 10 K, (b) 50 K, (c) 100 K, (d) 300 K.
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Fig. 6 Hysteresis (M-H) loops of COZ recorded at (a) 10 K, (b) 50 K, (c) 100 K, (d) 300 K. Fig. 7 Hysteresis (M-H) loops of TBCOZ recorded at (a) 10 K, (b) 50 K, (c) 100 K, (d) 300 K. Fig. 8 Initial Magnetization vs field curve fitted by BMP model of TBCOZ at (a) 300 K (b) 50 K and (c) Variation of Dielectric constant with frequency for samples COZ, TBZ and
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TBCOZ, (d) Plot of loss tangent (tanδ) vs frequency for samples COZ, TBZ and TBCOZ.
TABLE CAPTION
TABLE I. Structural parameters calculated from XRD data are lattice parameter (a, c), packing factor (c/a), unit cell volume (V), average nanocrystallite size (D), micro strain (ɛ).
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TABLE I
V (Ǻ)3
ɛ (10-4)
1.6012
48.5881
6.9552
5.1614
1.5982
46.6182
8.6600
5.1188
1.6001
45.3614
9.6373
5.1404
1.6009
45.8953
9.6696
a (Ǻ)
c (Ǻ)
c/a
ZnO
49.8192
3.2723
5.2397
COZ
43.6179
3.2295
TBZ
35.9538
3.1989
TBCOZ
40.9237
3.2109
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D (nm)
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FIG. 4
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FIG. 7
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FIG. 8
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Highlights: Nanoparticles of pure ZnO and Tb3+ and Co3+ ions doped ZnO were prepared by hydrothermal method.
Raman and PL spectra.
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The presence of various intrinsic defects in the doped samples was confirmed by
value of coercivity even at room temperature.
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This doped ZnO nanoparticles show a very strong magnetic interaction and a large
Enhanced dielectric value was obtained due to defect induced space charge
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polarization in the doped samples.
