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Defect-Mediated Magnetism in Pure CaO Nanopowders Daqiang Gao, Jinyun Li, Zhuoxin Li, Zhaohui Zhang, Jing Zhang, Huigang Shi, and Desheng Xue* Key Laboratory for Magnetism and Magnetic Materials of MOE, Lanzhou UniVersity, Lanzhou 730000, P. R. China ReceiVed: December 18, 2009; ReVised Manuscript ReceiVed: April 28, 2010
The intrinsic nature of ferromagnetism in CaO nanopowders has been established with the experimental observation of magnetic hysteresis loop at room temperature. On the basis of the results of X-ray diffraction, magnetic properties, positron annihilation lifetime spectra, and first-principles calculations, it is found that there are correlations among the lattice constant, the defect concentration, and the magnetization of CaO nanopowders, which suggests that Ca defects are the main reason for the magnetic order and the defect concentration is related with the ferromagnetism. Such a ferromagnet without the presence of any transition metal could be a very good option for a class of spintronics. Introduction Since observed room temperature ferromagnetism (RTFM) in pure HfO2 by Venkatesan et al.,1 there have been lots of reports about unexpected magnetization in pure wide-band-gap semiconductors, such as TiO2,2 HfO2,3 CeO2,4 and ZnO,5,6 where the observed RTFM were considered to originate from point defects on the surface or the interface of nanograins.7–9 Recently, RTFM has been shown theoretically and known experimentally in insulator of MgO by Gallego et al.10 and Hu et al.11 They suggested the ferromagnetism (FM) was induced by Mg defects. Gao et al. also indicated that neutral oxygen vacancies in the supercell of MgO induced no FM, while Mg defects could bring large magnetic moments based on first-principles calculations.12 More recently, Kenmochi et al. found that the FM can be induced by doped N or C in CaO through first-principles calculations.13 Elfimov et al. also suggested that cation vacancies in pure CaO have an open-shell electronic configuration and carry a nonzero local magnetic moment,14 which could be a path to a new ferromagnet, but so far without experimental proof. Osorio-Guille’n et al. suggested that a single Ca defect has a magnetic moment and the ferromagnetic interaction between two defects only extends to four neighbors, which means that to achieve magnetic percolation on an fcc lattice with such an interaction range one needs the defects concentration of 1.8 × 1021 cm-3. According to total-energy calculations for CaO, they demonstrated that a nonequilibrium defect enhancement factor of 103 is needed to achieve magnetism in such systems.15 It seems that it is difficult to observe FM in pure CaO in experiment. Actually, on the basis of the same theory, they suggest that it is more difficult to obtain FM in HfO216 for needing of higher defect concentration than CaO. But more recently, there are lots of reports about the observation of RTFM in HfO2 originated from point defects,3,17,18 although Rao et al. and Abraham et al. did not observe the FM in the pure HfO2 film.19,20 These results indicate that different preparation methods could produce crystals with different defect concentrations. Defects in CaO may have a similar interaction radius and the CaO lattice has a lower percolation threshold than HfO2 lattice,16 so it is possible to observe the FM in CaO by using a * Corresponding author: E-mail:
[email protected].
far-from-equilibrium preparation method to get more defects.18 In this work, CaO nanopowders were synthesized by the sol-gel method, and the RTFM is observed in CaO nanopowders. We suggested that the observed FM is induced by Ca defects. Experimental Section CaO powders were prepared by the sol-gel method. First, 4 g of Ca(NO)2 · 4H2O was dissolved in the mixture of 20 mL of 2-methoxyethanol and 0.4 mL of ethylenediamine. The dissolved solution was stirred for 4 h at 60 °C. Then, it was dried at 90 °C to form the precursor. The thermogravimetric analysis for precursor indicates that the crystallization temperature is about 650 °C, so we chose the annealing temperature of 700, 800, 900, and 1000 °C in order to get pure phase of CaO powders. All the samples were obtained by annealed the precursor 2 h in air, and we named them as S1, S2, S3, and S4, respectively. Any possibility of magnetic contaminations through accidental or trace impurities has been meticulously avoided during the sample preparation. High-purity quartz boat without any metallic contamination was used for annealing the samples in the furnace. The morphologies of the nanopowders were obtained by using transmission electron microscopy (TEM, JEM-2010). X-ray diffraction (XRD, X’ Pert PRO PHILIPS with Cu KR radiation) was employed to study the structure of the nanopowders. The doping levels and the bonding characteristics were determined by X-ray photoelectron spectroscopy (XPS, VG ESCALAB 210). Peak positions are referenced to the adventitious C 1s peak taken to be 285.0 eV. The measurements of magnetic properties were made using the Quantum Design MPMS magnetometer based on superconducting quantum interference device (SQUID) and the vibrating sample magnetometer (VSM, Lakeshore 7304). The thermogravimetry (TG) differential thermal analysis (Du Pont Instrument 1090B) was employed to obtain the variation of mass during the oxygen annealing. Positron annihilation spectroscopy (PAS) was used to investigate the variation of the defect concentration given by positron annihilation lifetime spectra. Results and Discussion Structural analysis of CaO nanopowders annealed at different temperatures was carried out by XRD, and the results are shown
10.1021/jp911957j 2010 American Chemical Society Published on Web 06/18/2010
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Figure 2. M-H curves for samples S1, S2, S3, and S4, which were investigated at room temperature by VSM under the maximum applied magnetic field of 8000 Oe. ZFC-FC curves for samples S1, S2, S3, and S4 in the temperature range from 2 to 330 K are shown in the inset.
Figure 1. (a) XRD pattens for samples S1, S2, S3, and S4. The left inset is the magnification of the XRD pattens, and the right inset shows the variation of the lattice parameter a of samples S1, S2, S3, and S4. (b) TEM image and (c) HRTEM image of sample S1.
in Figure 1a. All the diffraction peaks can be indexed to the rocksalt structure of CaO (JCPDS 82-1690). The absence of any other peaks suggests that there is no secondary phase present. Using the Scherrer formula for the full width at halfmaximum of the main peaks, the average crystalline size was estimated to be around 41, 48, 54, and 63 nm for the samples of S1, S2, S3, and S4, respectively. As can be seen, the diffraction peaks shift to the lower angle with the increasing of the annealed temperature, implying the extending of the lattice constant (a). Refinement of lattice constants for samples S1-S4 is shown in the inset of Figure 1a. The results indicate a increases with the increasing of annealing temperature. Figure 1b shows the TEM image of S1, which reveals that powders congregate together and the average size is about 180 nm. It can be clearly seen from the HRTEM of S1 in Figure 1c that as-grown CaO powders has a lattice spacing of 0.2408 nm, which is equal to the lattice constant of the standard CaO in (200) plane. The magnetization versus magnetic field (M-H) curves for CaO nanopowders are displayed in Figure 2, which were investigated at room temperature by VSM under the maximum applied magnetic field of 8000 Oe using a sample holder of high-purity capsules free from any metallic impurity. The same measurement procedure was repeated for empty sample holder, and the result shows it is paramagnetic (PM), and the PM signal of the holder was subtracted from the measured magnetic signal of CaO samples. The hysteresis loops indicate that all the samples have clearly RTFM, and the measured coercivities are 120, 87, 70, and 77 Oe for samples S1, S2, S3, and S4, respectively. It is worth noticing that the magnetism of the samples strongly depends on the annealing temperature: the saturation magnetization decreases from 0.031 to 0.007 emu/g with the increasing of the annealing temperature from 700 to 1000 °C. The inset of Figure 2 shows zero-field-cooled (ZFC) and field-cooled (FC) magnetization curves in the temperature range of 2-330 K at a dc field of 100 Oe of the CaO nanopowders, which is a further indication of FM for the samples. There is no blocking temperature in this temperature range for all samples, so the ferromagnetic contamination can
Figure 3. (a) M-H curves for samples S1, S2, S3, and S4, which were investigated at 10 K by SQUID under the maximum applied magnetic field of 10 000 Oe. (b) M-H curves for sample S2 at different temperatures.
be ruled out. At the same time, it can be concluded that the Curie temperature of this sample is above 330 K.21 The M-H curves of samples at 10 K also measured by SQUID, and the results are shown in Figure 3a, where the PM signal of the holder has been subtracted. It can be seen that all the samples at 10 K show FM and the saturation magnetization for all samples are larger that that of RT. Figure 3b shows the M-H curves for sample S2 at different temperatures from 10 to 250 K. The PM signal contribution due to the holder has been subtracted, and the magnetization starts to saturate in fields of about 5000 Oe. The magnetization curve shows considerable hysteresis, and the coercive field decreases in a monotonic fashion from a value of 214 Oe at 10 K to a value of about 80 Oe near RT. The M-H curves of the sample show saturation magnetization in the temperature range of 10-300 K, which is a typical behavior of a ferromagnetic material.
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Figure 4. XPS spectrum for sample S1. Ca 2p and O 1s core level XPS spectra are shown in the inset.
To explain the origin of FM in undoped semiconductors or oxides, a careful consideration whether the contamination is responsible for the FM has to be undertaken. In our experiments all the processes were carried out very carefully, and the capsules used to hold the samples during the magnetic measurements were also checked and showed no ferromagnetic signal. Since extreme precautions were taken during different steps involved in the experimental process to avoid any magnetic contamination, only probable suspect could be trace magnetic impurities present in the precursor materials. Iron was the only trace magnetic impurity present in the main precursor material Ca(NO)2 · 4H2O in proportion of 0.0002 at. %. However, magnetic measurement indicated that the precursor powder is PM. So the trace magnetic impurity present in the main precursor material can be ruled out. The further evidence for the purity and composition of the products was obtained by XPS, and the results show that the indexed peaks were corresponding to C, O, and Ca for all the CaO samples. A representative XPS spectrum of S1 is shown in Figure 4, and the Ca 2p and O 1s core level XPS spectra are shown in the inset. For the Ca 2p spectrum, peaks correspond to the core level 2p3/2 and 2p1/2 transitions of calcium at 346.6 and 351.1 eV,22 respectively, which reveals that calcium in the sample is bivalent. So far, the role of trace impurities in producing RTFM in CaO can be ruled out decisively. Therefore, we suggest that the observed RTFM is intrinsic for all CaO samples and reconsider the possibility of RTFM that was previously assumed for the other undoped oxides: FM may be due to vacancies and/or defects.7–9,23,24 According to the theoretical calculations of Elfimov et al. and Osorio-Guille’n et al., the observed FM in CaO having the rocksalt structure may be connected with the effect of Ca defects.17,18 In order to experimentally verify this prediction, we have examined the influence of the vacuum annealing on magnetism of CaO powders. At the same time, oxygen annealing was also done to compare with it. The results indicate that FM reduced after a subsequent annealing of S1 in vacuum (1 × 10-3 Pa) at 500 °C for 1 h, while oxygen annealing seems to have no influence on the FM for the samples (see inset in Figure 5a). The same postannealing was done for all the samples in vacuum (1 × 10-3 Pa) at 500 °C for 1 h. Usually, a vacuum annealing effectively introduce some oxygen vacancies in oxides while oxygen annealing reduces it. From the variation of magnetization (Figure 5a) we suggest that RTFM in CaO powders originates from Ca defects rather than oxygen vacancies.11 Here, another post-oxygen annealing result can rule out the existence of oxygen vacancies. Sample S2 was placed in the quartz boat in an oxygen atmosphere at a flux of 250 N mL/min in the temperature range of 50-800 °C (10 °C/min)
Figure 5. The positron annihilation parameters for samples S1, S2, S3 and S4, defect-specific lifetime originates from the annihilation of positrons at defect-type defects (a) τ1 within the grains and its intensity I1, (b) τ2 comes from grain boundary and its intensity I2, (c) the mean lifetime jτ.
and then cooled down (10 °C/min), and this progress was repeated about four times. Thermogravimetric (TG) differential thermal analysis was employed to obtain the variation of the masses during these annealing progresses. Generally, oxygen annealing can reduce the oxygen vacancies by bonding the oxygen atoms and cations together. So if oxygen vacancies exist in sample, the mass of the sample must increase through the long-time oxygen annealing. However, from the results shown in Figure 5b we can see that during these four progresses (except for first time which leads by decomposition of Ca(OH)2 on the surface of CaO powders) the mass of S2 has some decrease and the variation also decreases from first to fourth, which may be induced by some O atoms escaping from CaO to form high crystallization. This result further confirmed that the observed RTFM in CaO powders is induced by Ca defects, which is consistent with previous theoretical calculations.14 Positron annihilation spectroscopy (PAS) can be used to analyze vacancy-type defect in oxides.25 Positron annihilation lifetime spectroscopy (PALS) is based on measuring the lifetimes of positrons injected into a material to analyze the defect concentration of cation.26 Here, the CaO powders were pressed at the pressure of 20 MPa about 10 min for the measurement of PALS. The positron annihilation lifetime spectra of the samples S1-S4 with an experimental time resolution 182 ps and about 2 × 106 counts were collected and fitted into three lifetime components. Generally, τ3 represents the annihilation of positrons between the interspace of the powders, and the values for samples S1-S4 are basically the same. So we do not take it into account here. A defect-specific lifetime τ1 originates from the annihilation of positrons within the grains while τ2 comes from grain boundaries. Meanwhile, I1 or I2 is the intensity of each positron annihilation lifetime (shown in Figure 6a,b). It is clear seen that τ1 and τ 2 decreased as the annealed temperature increased, implying that the cation defect concentration whatever within the grains or grain boundaries
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Figure 7. (a) Calculated lattice constants of different Ca defect concentrations. (b) Variation of magnetization for S2 and postannealing S2 in step 1, step 2, and step 3. XRD results are shown in the inset.
Figure 6. Positron annihilation parameters for samples S1, S2, S3, and S4; defect-specific lifetime originates from the annihilation of positrons at defect-type defects. (a) τ1 within the grains and its intensity I1. (b) τ2 comes from grain boundary and its intensity I2. (c) The mean lifetime jτ.
decreases gradually. The mean lifetime jτ can reveal the cation defect concentration in the whole sample, which is defined as
τ¯ ) τ1
I1 I2 + τ2 I1 + I2 I1 + I2
As can be seen, jτ decreases from sample S1 to S4, which indicates that the Ca defect concentration decreases gradually from sample S1 to S4. At the same time, the variation of Ca defect concentration is consistent with that of the magnetization for CaO powders from S1 to S4. So we can conclude that the magnetization decreases as the decreasing of Ca defect concentration. On the basis of first-principles calculations, Khalid et al. showed that an increase in the concentration of Zn vacancies leads to an effective reduction of the equilibrium lattice constant a.27 Similarly, first-principles calculations were performed for CaO in our case. We apply the hybrid full-potential linearized augmented plane-wave plus local orbitals (FP-L/APW+lo) method within the density functional theory (DFT) as implemented in the Wien2k code 22, and the generalized gradient approximations 23 (GGA) were used to describe the exchangecorrelation functional. A 2 × 2 × 2 CaO supercell containing 64 atoms was chosen for the calculation. The cases of 0, 1, 2, and 3 neutral Ca vacancies, corresponding to concentrations of 0%, 3.125%, 6.25%, and 9.375%, respectively, were calculated with the convergence criteria of total energy less than 10-4 Ry.
Theoretical calculations indicate that a Ca deficiency results in a magnetic moment of 2 µB per vacancy, about 90% of which is concentrated on the six oxygen ions that are the nearest neighbors to the Ca vacancy. Meanwhile, the result shown in Figure 7a indicates that an increase in the concentration of Ca vacancies leads to an effective reduction of the equilibrium lattice constant a. In experiment, XRD results reveal the decrease of lattice constant a (shows in Figure 1a), and PALS results indicate Ca defect concentration increases in CaO samples from S4 to S1. So the experimental finding can be confirmed by our first-principles investigations. At the same time, we can also conclude that lattice constant a increases as the magnetization decreases. For further verifying this, the three-step postannealing process for S2 was undertaken to study the relationship between lattice constant a and magnetization. A sample of S2 was placed in a quartz boat for annealing in air in the temperature range of 50-800 °C with the heating rate about 10 °C/min and then natural cooling, and the process was named step 1. Step 2 and step 3 were the repeats of step 1. XRD results indicated that the diffraction peaks shift to lower angle gradually from S2 to progress step 3, implying the increase of the lattice constant (which means the decrease of the defect concentration, shown in Figure 7b). Meanwhile, the obvious decrease of magnetization is observed after the three-step annealing (see in Figure 7b), which is consistent with the expected result. Therefore, it is concluded that the creation of larger defects is related to the observed FM in pure CaO nanopowders. However, the real concentration of these defects in CaO powders and how defects induce FM in CaO powders remain to be determined, which is a challenge for future work. Conclusion In conclusion, we have investigated the magnetic properties of CaO nanopowders prepared by a simple sol-gel method. Magnetic measurements indicate that all as-prepared CaO nanopowders show the RTFM, and Ca defects are the reason
Magnetism in Pure CaO Nanopowders for the magnetic order. PALS results indicate that Ca vacancies decrease from S1 to S4, which combining with experiment results and first-principles calculations suggests that Ca defects are the reason for the magnetic order in CaO nanopowders. Further theoretic investigations on the defects introducing FM are expected, and our work is under way. Acknowledgment. This work is supported by National Science Fund for Distinguished Young Scholars (Grant No. 50925103), The Keygrant Project of Chinese Minisity of Education (Grant No. 309027), The Fundamental Research Funds for the Central Universities (Grant No. Lzujbky-2009162), and NSFC (Grant No. 50801033). References and Notes (1) Venkatesan, M.; Fitzgerald, C. B.; Coey, J. M. D. Nature (London) 2004, 430, 630. (2) Sudakar, C.; Kharel, P.; Suryanarayanan, R.; Thakur, J. S.; Naik, V. M.; Naik, R.; Lawes, G. J. Magn. Magn. Mater. 2008, 320, L31. (3) Hadacek, N.; Nosov, A.; Ranno, L.; Strobel, P.; Gal’era, R. M. J. Phys.: Condens. Matter 2007, 19, 486206. (4) Liu, Y. L.; Lockman, Z.; Aziz, A.; Driscoll, J. M. J. Phys: Condens. Matter. 2008, 20, 165201. (5) Banerjee, S.; Mandal, M.; Gayathri, N.; Sardar, M. Appl. Phys. Lett. 2007, 91, 182501. (6) Darshana, Y.; Inamdar; Amit, D. L.; Arjun, K. P.; Igor, D.; Naushad, A.; Shailaja, M. J. Phys. Chem. C 2010, 114, 1451. (7) Han, X. P.; Lee, J. C.; Yoo, H. I. Phys. ReV. B 2009, 79, 100403. (8) Coey, J. M. D.; Venkatesan, M.; Stamenov, P.; Fitzgerald, C. B.; Dorneles, L. S. Phys. ReV. B 2005, 72, 024450. (9) Liu, E. Z.; Jiang, J. Z. J. Phys. Chem. C 2009, 113, 16116. (10) Gallego, S.; Beltra´n, J. I.; Cerda´, J.; Mun˜oz, M. C. J. Phys: Condens. Matter 2005, 17, L451.
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