- Email: [email protected]
Micro-Raman and XPS studies of pure ZnO ceramics
ARTICLE IN PRESS Physica B 405 (2010) 2492–2497
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
Micro-Raman and XPS studies of pure ZnO ceramics J. Das a,n, S.K. Pradhan b, D.R. Sahu c, D.K. Mishra d, S.N. Sarangi e, B.B. Nayak d, S. Verma e, B.K. Roul b a
Department of Physics, Silicon Institute of Technology, Bhubaneswar 751024, Orissa, India Institute of Materials Science, Acharya Vihar, Bhubaneswar 751013, Orissa, India School of Physics, University of the Witwatersrand, Private Bag 3, Wits 2050, Johannesburg, South Africa d Advanced Materials Technology Department, Institute of Minerals and Materials Technology (CSIR), Bhubaneswar 751013, Orissa, India e Institute of Physics, Sachivalaya Marg, Bhubaneswar 751004, Orissa, India b c
a r t i c l e in f o
a b s t r a c t
Article history: Received 10 February 2010 Received in revised form 5 March 2010 Accepted 8 March 2010
We observed that ZnO bulk ceramics prepared by solid state reaction route at 1300 1C exhibited unexpected room temperature ferromagnetic (RTFM) property. In the absence of any magnetic impurity the cause of room temperature ferromagnetic signal in the undoped system is certainly attributed to various kinds of native defects such as oxygen vacancies (VO) or zinc interstitials (IZn) and their clusters created inside the bulk ceramics during heating by slow step sintering schedule (SSSS). The microRaman investigation and X-ray photoelectron spectroscopy studies on the ZnO sample sintered at high temperature confirm the presence of such lattice defects. & 2010 Elsevier B.V. All rights reserved.
Keywords: Room temperature ferromagnetism Lattice defect X-ray photoelectron spectroscopy Zinc interstitial
1. Introduction Over the past decade, ZnO, an n-type semiconductor of II–IV group, has become the centre of attraction of many researchers all over the globe for its unique optical and electrical properties. Until now, it has found wide range of applications in micro- and opto-electronic fields owing to its wide band gap of 3.437 eV at 2 K, large excitation energy of 60 meV and its piezo-electric nature. However, any magnetic and semi-magnetic properties associated with this material system was not established until the theoretical prediction of possible room temperature ferromagnetism (RTFM) in transition metal (TM) doped ZnO by Dietl et al. [1]. Following the prediction, many research groups have been encouraged to put uninterrupted efforts in searching for hightemperature (preferably around room temperature) ferromagnetism in doped and undoped ZnO in real practice, which would make this already well-studied material a futuristic potential candidate for multifunctional device applications. Unfortunately, there have been a wide range of controversies over the exact kind of magnetic property observed in ZnO-based materials reported till date [2–9], which give rise to difficulty in the establishment of a confirmed ferromagnetic behavior in the material. Apart from this, the cause of the magnetic properties, is also not properly understood and is still under debate. There have been
n
Corresponding author. Tel.: + 91 674 2725448/649; fax: + 91 674 2725312. E-mail address: [email protected] (J. Das).
0921-4526/$ - see front matter & 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2010.03.020
speculations [10–13] that the RTFM in the thin films and bulk samples of some non-magnetic metal oxides observed by some researchers may be caused by intrinsic defects present or created during the processing of the material, other than the TM dopants. While most of the previous studies have been focused on ZnObased diluted magnetic semiconductors (DMS), in which the magnetic property was understood to be extrinsic, i.e., induced by doping of 3dn TM ions into the host semiconductor [14,15], the reported observation of room temperature ferromagnetism in the thin film of the band insulator HFO2 [16], which do not contain partially filled d or f shell, created high enthusiasm among the researchers to reinvestigate the mechanism of magnetism in the light of a new phenomena called ‘‘d0 ferromagnetism’’. It is speculated that this type of so called d0 magnetism is also possible in many other undoped diamagnetic metal oxides like TiO2, ZnO, In2O3 and SnO2, etc., where the origin of RTFM is assumed to be the exchange interactions between localized electrons spin moments resulting from oxygen vacancies in the system. As such, ZnO, which is strongly n-type, the dominant donor is a native defect, i.e., either the O vacancy (Ov) or the Zn interstitial [17]. In fact, it is clear that a lack of proper understanding of the possible mechanism of magnetism inside the system has hindered the full potential application of this material. However, the theory of defect-induced ferromagnetism in undoped ZnO has been unequivocally established by several research groups [18,19], which provides a new opportunity for fabricating high-temperature ZnO-based devices with enhanced magnetic property for practical applications [20].
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
Since bulk materials are important, both because of their potential application and also to understand the intrinsic physics behind the various properties reported for the material, we have attempted to prepare ZnO bulk ceramics, which would be very much useful in the form of targets for the subsequent thin film preparation for various kinds of practical applications. In this paper, room temperature ferromagnetism in ZnO has been reported with necessary supportive explanations, which are obtained from our micro-Raman and X-ray photoelectron spectroscopy analysis.
2493
2. Experimental ZnO bulk ceramics were synthesized by standard solid state reaction route. Appropriate proportion of highly pure zinc oxide (ZnO) powder (Puraterm-99.999%) was taken in an agate mortar and was ground properly for 2 h. The powder was initially heated at 400 1C for 6 h followed by room temperature quenching and grinding. This procedure was repeated five times in order to achieve homogeneous mixture of green powder with smaller particle size. Green powder was palletized after mixing with freshly prepared Poly Vinyl Alcohol (PVA) as binder. Cylindrical pellets with approximately 2 mm thickness and 10 mm diameter were prepared by hydraulic press having 30 ton base capacity with a pressure of 10 ton/cm2. Pressed pellets were slowly heated (50 1C/h) to 600 1C and kept for 8 h for slow release of PVA binder from the pellets using high-temperature programmable (Eurotherm controller, Model: 2404) vacuum furnace. A slow step sintering schedule (Fig. 1) was adopted up to 1300 1C to obtain highly dense ZnO bulk ceramics. X-ray diffraction (XRD) and scanning electron microscopy studies are carried out by Bruker D8 Advance X-ray Diffractometer and SEM (Philips FEG Xl’30), respectively for structural and morphological analyses. X-ray photoelectron spectroscopy and micro-Raman studies are carried out for the compositional analysis and also to confirm the presence of defects. All the hysteresis measurements of the samples were taken at low temperature (10 K) and room temperatures (300 K) using superconducting quantum interference device (SQUID) magnetometer with maximum field of 1 T.
3. Results and discussions Fig. 1. Slow step sintering schedule (SSSS) of pure ZnO.
Fig. 2 shows the XRD pattern of ZnO bulk ceramics as prepared and subsequently sintered at 850 and 1300 1C, respectively. The
Fig. 2. X-ray diffraction pattern of (a) ZnO as prepared, (b) ZnO sintered at 850 1C and (c) ZnO sintered at 1300 1C.
Table 1 Crystallite sizes, lattice parameter of ZnO (virgin) and ZnO bulk samples sintered at 850 1C and 1300 1C. Sample
Avg. crystallite size (nm)
Lattice parameters
c/a
ZnO (virgin) ZnO (850 1C) ZnO (1300 1C)
107.954 103.1786 97.53,414
a¼ 3.2450 c¼ 5.2000 a¼ 3.2410 c¼ 5.2019 a¼ 3.26 c¼ 5.255
1.60,246 well matched to the value of the reported 1.6025 (O) 1.6050 1.6119
ARTICLE IN PRESS 2494
J. Das et al. / Physica B 405 (2010) 2492–2497
peaks in the figure are indexed to hexagonal wurtzite structure of ZnO belonging to space group P63 mc (JCPDF # 890510). Preferred orientation is along (1 0 1) as expected for bulk ZnO. No evidence of any other phases is clear in the pattern. The average crystallite sizes are calculated using Scherrer’s formula and presented in Table 1. The XRD pattern showed a distinct decrease in crystallite size with an increase in crystalline order with sintered temperature. Lattice parameters of ZnO as prepared and sintered at 850 and 1300 1C, respectively, are calculated using the formula 1 4 h2 þ hk þk2 l2 þ 2 ¼ 3 a2 c d2hkl
Fig. 3. A plot showing the variation in c/a ratio with sintering temperature.
which are presented in Table 1. It is evident that the c/a ratio increases with an increase in sintering temperature (shown in Fig. 3), which is in agreement with the value of pure ZnO [Powder Diffraction File 36-1451 for hexagonal ZnO (1997)]. The crystallite size decreases with an increase in lattice parameter might result from the vacancies associated with the structure [21]. To study the physical densification of bulk ZnO ceramics processed by SSSS, micrograph study of the bulk surface was performed by SEM. Figs. 4a and b show the microstructure photograph of pure ZnO bulk ceramics sintered at 850 and 1300 1C for 24 h. It is seen that average microcrystallites of 8 mm dimension are grown with triangular grain boundary junction including the presence of multiple sub-micron pores of irregular structure (shown in Fig. 4c) in ZnO bulk ceramics sintered at 850 1C. The dimension of the pores is of the order of E1 mm. During grain growth process at elevated temperature (i.e., at around 1300 1C), mass flow to respective pores at atomic level occurs decreasing the number of pores and hence enhancing the physical densification. This is evident in Figs. 4b and d, which showed clean and clear triangular grain junction relating to three well-shaped uniform microcrystallites. This particular sample showed highest density as well as ferromagnetism at room temperature. A wide survey scan of XPS spectra is taken in the range 0–1100 eV as shown in Fig. 5a. No extra peak corresponding to any magnetic impurities other than Zn and O is observed in the figure. Fig. 5b shows the XPS spectra of C 1s. The 284.6 eV binding energy of C 1s is usually used as an internal reference in the spectrum [22]. The spectra corresponding to O 1s core and Zn 2p core level of ZnO bulk sample sintered at 1300 1C is calibrated by binding energy of C 1s (284.6 eV) as the internal reference spectra [22] and is shown in Figs. 5c and d. The O 1s spectrum is fitted with three Gaussian peaks having three binding energy components as per the literature [23]. The lower binding energy component at 529.44 eV is attributed to the wurtzite structure of
Fig. 4. SEM photographs of ZnO bulk ceramics sintered at (a) 850 1C, (b) 1300 1C, (c) presence of multiple pores in 850 1C-sintered ZnO ceramics and (d) the decrease in no. of pores due to grain growth in 1300 1C-sintered ZnO ceramics.
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
2495
Fig. 5. (a) Wide survey X-ray photoelectron spectrum of ZnO, (b) C 1s spectrum of ZnO, (c) O 1s spectrum of ZnO and (d) Zn 2p spectrum of ZnO.
hexagonal Zn2 + ion of the metal oxide [24]. The medium and higher binding energy components at 530.85 and 531.37 eV, respectively, are assigned to the oxygen deficient region within the ZnO matrix [24]. The asymmetry nature of the curve reflects the variation of oxidation number of oxygen. The Zn 2p spectrum in Fig. 6d shows a doublet whose binding energies are 1022.54 and 1045.69 eV and can be identified as Zn2p3/2 and Zn2p1/2 lines, respectively. The binding energy difference between the two lines is 23.15 eV, which is well lying within the standard reference value of ZnO [22]. The binding energy and the binding energy difference value calculated from the XPS study show that Zn atoms are in +2 oxidation state. At this stage, it is not possible to predict the presence of Zn interstitial defects from Zn 2p spectra, which may be established only by studying the Auger peak of Zn. The result obtained from micro-Raman analysis of pure ZnO bulk ceramics sintered at 1300 1C by SSSS for 24 h is shown in Figs. 6a and b. We observed slightly different spectra by focusing the beam at the grain boundary junctions and at the centre of the grain, respectively, with an aim to study the presence of defects attributed to intrinsic impurity and vacancy of anions and cations. 4 As known, the wurtzite semiconductor ZnO belongs to C3v (P63 mc) space group with two formula units per primitive cell,
where all atoms occupy C3n sites. The Raman active zone-centre optical phonons predicted by the group theory are A1+2E2 + E1+ 2B1. The B1 modes are silent modes, the A1 and E1 modes are polar and exhibit different frequencies for the transverse-optical (TO) and longitudinal-optical (LO) phonon modes. Both A1 and E1 are Raman and infrared active, whereas the E2 modes are nonpolar having two frequencies: E2 (high) associated with oxygen displacement and E2 (low) associated with Zn sub-lattice and are Raman active only [25,26]. As seen in the figure, the intensity of the peak corresponding to E2 (high) vibration mode, which is the band characteristic of the wurzite phase of ZnO and strongly depend on the isotopic composition of ZnO, is high and is centered at 436 cm 1, when the laser beam is focused at the centre of the grain. But its intensity is comparatively low and found to be centered at 433 cm 1 when the beam is focused at the grain boundary junction. Red shift in E2 (high) phonon mode of about 3–6 cm 1 is observed in both the spectra as compared to 439 cm 1 of ZnO bulk single crystal. The reason of such red shift due to the local heating effect [27] by the laser source can be ruled out as the excitation laser is too weak. We emphasize that the shift is only due to the defects produced during the sample processing by our SSSS method. Nonpolar E2 (high) phonon mode is usually associated with
ARTICLE IN PRESS 2496
J. Das et al. / Physica B 405 (2010) 2492–2497
Fig. 6. Raman spectra of ZnO bulk ceramic sample sintered at 1300 1C (a) Beam focussed at the centre of the grain and (b) beam focussed at the grain boundary.
Fig. 7. Magnetization vs. magnetic field (M–H) curves (a) at T¼ 10 K and (b) at T¼ 300 K for ZnO bulk samples sintered at 1300 1C.
oxygen atom displacement [28]. Hence, we assume that oxygen vacancy and host disorders are produced in ZnO during processing. The oxygen vacancies produced decreases the average atomic mass and causes mass fluctuation at oxygen position, resulting in the shifting of E2 line centre towards lower frequency. Again, a red shift of around 6 cm 1 in the E2 peak has been noted in the case of grain boundary junction as compared to the red shift of 3 cm 1 in case of beam focused at the centre of the grain. The larger red shift in E2 (high) phonon mode in case of laser beam focused at the grain boundary suggests that vacancies created during physical densification process while sintering, are mostly populated at the grain boundary (GB). Apart from this, distinct broader peaks along with shoulder peaks (SP) of considerable intensities are detected at around 575 cm 1 range in both the spectra, which correspond to 1LO (A1) mode. Normally, Raman peaks observed in between 570 and
590 cm 1 are considered to be associated with structural disorders, such as oxygen vacancy, Zn interstitial and their combination, due to the strong dependence on the oxygen stoichiometry [29,30]. Chen et al. [31] regarded this peak as a clear defect-induced mode. Our observation of broad shoulder peaks of almost same intensity associated with this 1LO peak clearly revealed the presence of defects or clusters of defects in the specimen, which might be created due to the outgasing of Zn during slow step heating of the sample. We speculate that vacancies created inside the bulk ceramics interacted with other nearby vacancies forming clusters, which is also the cause of the appearance of the above-mentioned broad SP associated with 1LO peak. Hysteresis behavior (M–H) of sintered bulk ZnO ceramics (processed at 1300 1C for 24 h) at 10 and 300 K is presented in Fig. 7 after correcting necessary background diamagnetic
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
subtraction. A clear M–H loop (inset) with saturation magnetization Ms ¼1.88 10 3 emu/g, coercivity Hc ¼34.44 Oe and remanent magnetization ¼8.07 10 5 emu/g is observed at 300 K (Fig. 7), indicating that pure ZnO bulk ceramics processed at 1300 1C by our SSSS method is ferromagnetic at RT, but becomes antiferromagnetic at low temperature (i.e., at 10 K) as clearly observed in the figure. The antiferromagnetic behavior is only due to the isolation of defects at low temperature whereas at room temperature there is a substantial enhancement in the ferromagnetic contribution in the sample, which is possibly due to the formation of clusters of defects in the sample when sintered at an elevated temperature of 1300 1C. As we have not incorporated any kind of magnetic impurities, this unconventional ferromagnetism observed here can definitely be assigned to the defects generated during material processing of this monoelemental (Zn) oxide. As known, contribution of defectinduced magnetization in the ZnO system is limited to oxygen vacancy (VO), oxygen interstitial (IO), Zn vacancy (VZn) and Zn interstitials (IZn), which are all non-magnetic, it is argued that only oxygen vacancy or zinc vacancy is responsible for RTFM. However, it has been explained by Banerjee et al. [10] that, in case of single O vacancy, at each of the next four Zn ions one sp3 hybrid of 4s and 4p orbital are found, which points towards the vacancy. As per their calculation a neutral O vacancy has two electrons in one state out of the four (one a1 and three degenerate t2), which is a singlet with zero spin, because of the splitting between a1 and t2 levels greater than Hund’s coupling (1 eV). Hence a single neutral oxygen vacancy cannot have a net magnetic moment associated with it. On the other hand, for a cluster of more than three O vacancies, the splitting is lower than Hund’s coupling, giving rise to a considerable net moment for the cluster. It is speculated by the researchers that heating at elevated temperature, may cause the O vacancies to migrate within the material and form clusters, favoring the decrease in the strain field energy ([10] and the Refs therein). Oxygen deficiency is manifested in the form of oxygen vacancy (deep donors) and zinc interstitials (shallow donors). It is suggested that the unpaired electron trapped in oxygen vacancies (F centre) are polarized to give room temperature ferromagnetism. Upon heating at elevated temperature, the ferromagnetic property was enhanced. It is also shown that for oxygen nonstoichiometry of 0.1% the magnetic moment contribution is considerably significant. It is understood that native defects play an essential role in understanding the behavior as well as for the successful application of any semiconductor and as such, in ZnO, specific defects have long been believed to play a more important role as its n-type conductivity has often been explained by the native point defects like oxygen vacancies and zinc interstitials [32]. We emphasize that during the processing, unreacted zinc outgases and annihilates the oxygen octahedron and tetrahedron networks and finally produces the above-mentioned defects and cluster of defect molecules, which are responsible to induce ferromagnetism in the slow step sintered specimen. Our microRaman as well as XPS study also clearly supported the presence of such defect structures in the sintered ceramics. However, the exact explanation of the ferromagnetic contribution by vacancy clusters needs to be studied more extensively both experimentally and theoretically.
4. Conclusion In conclusion, we emphasize that our experimental observation presents an alternate material preparation process for achieving ferromagnetic character in pure ZnO semiconductor
2497
without any magnetic cation substitution. Observation of defectinduced ferromagnetism with only small oxygen vacancies in bulk ZnO ceramics processed by SSSS at high temperature by us is definitely a challenge to explore and find the real explanation, which would be propitious for the development of multifunctional devices using ZnO. However, we also strongly urge that more research work in processing pure ZnO bulk ceramics (especially for bulk target for thin film materials) is essential along with satisfactory theoretical interpretation in this wellstudied material, which would open up a different avenue for magneto-optical applications.
Acknowledgement Authors are highly grateful to the ex-Director IMS, Prof. S.N. Behera for his constant guidance and encouragement to carry out research work at IMS, BBSR. We are also highly thankful to Mr. Santosh Choudhury to carry out the XPS study of our sample.
References [1] T. Dietl, H. Ohno, F. Matsukura, Science 287 (2000) 1019. [2] X.M. Cheng, C.L. Chien, J. Appl. Phys. 93 (2003) 7876. [3] S.J. Han, T.H. Jang, Y.B. Kim, B.G. Park, J.H. Park, Y.H. Jeonga, Appl. Phys. Lett. 83 (2003) 920. [4] T. Fukumura, Z. Jin, M. Kawasaki, T. Shono, T. Hasegawa, S. Koshihara, H. Koinuma, Appl. Phys. Lett. 78 (2001) 958. [5] A. Tiwari, C. Jin, A. Kvit Kumar, J.F. Muth, J. Narayana, Solid State Commun. 121 (2002) 371. [6] S.W. Jung, S.J. An, G.C. Yi, C.U. Jung, S. Lee, S. Cho, Appl. Phys. Lett. 80 (2002) 4561. [7] P. Sharma, A. Gupta, K.V. Rao, F.J. Owens, R. Sharma, R. Ahuja, J.M. Osorio Gullien, B. Johansson, G.A. Gehring, Nat. Mater. 2 (2003) 673. [8] J. Luo, J.K. Liang, Q.I. Liu, F.S. Liu, Y. Zhang, B.J. Sun, G.H. Rao, J. Appl. Phys. 97 (2005) 086106. [9] W. Chen, L.F. Zhao, T.Q. Wang, J.H. Miao, S. Liu, Z.C. Xia, S.L. Yuan, Solid State Commun. 134 (2005) 827. [10] S. Banerjee, M. Mandal, N. Gayathri, M. Sardar, Appl. Phys. Lett. 91 (2007) 182501. [11] K. Potzer, Z. Shengqiang, J. Grenzer, M. Helm, Fassbender, Appl. Phys. Lett. 92 (2008) 182504. [12] Q. Xu, H. Schmidt, S. Zhou, K. Potzger, M. Helm, H. Hochmuth, M. Lorenz, A. Setzer, P. Esquinazi, C. Meinecke, M. Grandmann, Appl. Phys. Lett. 92 (2008) 082508. [13] X. Zuo, S.D. Yoon, A. Yang, W.H. Duan, C. Vittoria, V.G. Harris, J. Appl. Phys. 105 (2009) 07C508. [14] H. Ohno, Science 281 (1998) 951. [15] T. Jungwirth, J. Sinova, J. Masek, J. Kucera, A.H. MacDonald, Rev. Mod. Phys. 78 (2006) 809. [16] J.M.D. Coey, Solid State Sci. 7 (2005) 660. ¨ [17] For a review, see G. Heiland, E. Mollwo, and F.Stockmann, in Solid State Physics, edited by F. Seitz and D. Turnbull, Academic, New York, vol 8, 1959. p. 191. [18] D.A. Schwartz, D.R. Gamelin, Adv. Mater. 16 (2004) 2115. [19] R. Radovanovic, D.R. Gamelan, Phys. Rev. Lett. 91 (2003) 157202. [20] S.A. Wolf, D.D. Awschalom, R.A. Buhraman, Science 294 (2001) 1488. [21] A. Sundaresan, R. Bhargavi, N. Rangarajan, U. Siddesh, C.N.R. Rao, Phys. Rev. B 74 (2006) 161306. [22] C.D. Wagner, W.M. Riggs, L.E. Davis, J.F. Moulder, G.E. Muilenberg, in: Handbook of X-ray Photoelectron Spectroscopy, Perkin Elmer, Eden Prairie, 1979. [23] M. Chen, X. Wang, Y.H. Yu, Z.L. Pei, X.D. Bai, C. Sun, R.F. Huang, L.S. Wen, Appl. Surf. Sci. 158 (2000) 134. [24] J.C.C. Fan, J.B. Goodenough, J. Appl. Phys. 48 (1977) 3524. [25] N. Ashkelon, et al., J. Appl. Phys. 93 (2003) 126. [26] J.F. Scott, Phys. Rev. B 2 (1970) 1209. [27] K.A. Alim, V.A. Fonoberov, A.A. Balandin, Appl. Phys. Lett. 86 (2005) 053103. [28] J. Serrano, F.J. Manjon, A.H. Romero, F. Widulle, R. Lauck, M. Cardona, Phys. Rev. Lett. 90 (2003) 055510. [29] R. Cusco, E.A. Llado, et al., Phys. Rev. B 75 (2007) 165202. [30] S.H. Jeong, J.K. Kim, B.T. Lee, J. Phys. D: Appl. Phys. 36 (2003) 2017. [31] Z.Q. Chen, et al., J. Appl. Phys. 97 (2005) 013528. [32] Q. Wang, Q. Sun, G. Chen, Y. KawaZoe, Puru Jena, Phys. Rev. B 77 (2008) 205411.
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
Micro-Raman and XPS studies of pure ZnO ceramics J. Das a,n, S.K. Pradhan b, D.R. Sahu c, D.K. Mishra d, S.N. Sarangi e, B.B. Nayak d, S. Verma e, B.K. Roul b a
Department of Physics, Silicon Institute of Technology, Bhubaneswar 751024, Orissa, India Institute of Materials Science, Acharya Vihar, Bhubaneswar 751013, Orissa, India School of Physics, University of the Witwatersrand, Private Bag 3, Wits 2050, Johannesburg, South Africa d Advanced Materials Technology Department, Institute of Minerals and Materials Technology (CSIR), Bhubaneswar 751013, Orissa, India e Institute of Physics, Sachivalaya Marg, Bhubaneswar 751004, Orissa, India b c
a r t i c l e in f o
a b s t r a c t
Article history: Received 10 February 2010 Received in revised form 5 March 2010 Accepted 8 March 2010
We observed that ZnO bulk ceramics prepared by solid state reaction route at 1300 1C exhibited unexpected room temperature ferromagnetic (RTFM) property. In the absence of any magnetic impurity the cause of room temperature ferromagnetic signal in the undoped system is certainly attributed to various kinds of native defects such as oxygen vacancies (VO) or zinc interstitials (IZn) and their clusters created inside the bulk ceramics during heating by slow step sintering schedule (SSSS). The microRaman investigation and X-ray photoelectron spectroscopy studies on the ZnO sample sintered at high temperature confirm the presence of such lattice defects. & 2010 Elsevier B.V. All rights reserved.
Keywords: Room temperature ferromagnetism Lattice defect X-ray photoelectron spectroscopy Zinc interstitial
1. Introduction Over the past decade, ZnO, an n-type semiconductor of II–IV group, has become the centre of attraction of many researchers all over the globe for its unique optical and electrical properties. Until now, it has found wide range of applications in micro- and opto-electronic fields owing to its wide band gap of 3.437 eV at 2 K, large excitation energy of 60 meV and its piezo-electric nature. However, any magnetic and semi-magnetic properties associated with this material system was not established until the theoretical prediction of possible room temperature ferromagnetism (RTFM) in transition metal (TM) doped ZnO by Dietl et al. [1]. Following the prediction, many research groups have been encouraged to put uninterrupted efforts in searching for hightemperature (preferably around room temperature) ferromagnetism in doped and undoped ZnO in real practice, which would make this already well-studied material a futuristic potential candidate for multifunctional device applications. Unfortunately, there have been a wide range of controversies over the exact kind of magnetic property observed in ZnO-based materials reported till date [2–9], which give rise to difficulty in the establishment of a confirmed ferromagnetic behavior in the material. Apart from this, the cause of the magnetic properties, is also not properly understood and is still under debate. There have been
n
Corresponding author. Tel.: + 91 674 2725448/649; fax: + 91 674 2725312. E-mail address: [email protected] (J. Das).
0921-4526/$ - see front matter & 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2010.03.020
speculations [10–13] that the RTFM in the thin films and bulk samples of some non-magnetic metal oxides observed by some researchers may be caused by intrinsic defects present or created during the processing of the material, other than the TM dopants. While most of the previous studies have been focused on ZnObased diluted magnetic semiconductors (DMS), in which the magnetic property was understood to be extrinsic, i.e., induced by doping of 3dn TM ions into the host semiconductor [14,15], the reported observation of room temperature ferromagnetism in the thin film of the band insulator HFO2 [16], which do not contain partially filled d or f shell, created high enthusiasm among the researchers to reinvestigate the mechanism of magnetism in the light of a new phenomena called ‘‘d0 ferromagnetism’’. It is speculated that this type of so called d0 magnetism is also possible in many other undoped diamagnetic metal oxides like TiO2, ZnO, In2O3 and SnO2, etc., where the origin of RTFM is assumed to be the exchange interactions between localized electrons spin moments resulting from oxygen vacancies in the system. As such, ZnO, which is strongly n-type, the dominant donor is a native defect, i.e., either the O vacancy (Ov) or the Zn interstitial [17]. In fact, it is clear that a lack of proper understanding of the possible mechanism of magnetism inside the system has hindered the full potential application of this material. However, the theory of defect-induced ferromagnetism in undoped ZnO has been unequivocally established by several research groups [18,19], which provides a new opportunity for fabricating high-temperature ZnO-based devices with enhanced magnetic property for practical applications [20].
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
Since bulk materials are important, both because of their potential application and also to understand the intrinsic physics behind the various properties reported for the material, we have attempted to prepare ZnO bulk ceramics, which would be very much useful in the form of targets for the subsequent thin film preparation for various kinds of practical applications. In this paper, room temperature ferromagnetism in ZnO has been reported with necessary supportive explanations, which are obtained from our micro-Raman and X-ray photoelectron spectroscopy analysis.
2493
2. Experimental ZnO bulk ceramics were synthesized by standard solid state reaction route. Appropriate proportion of highly pure zinc oxide (ZnO) powder (Puraterm-99.999%) was taken in an agate mortar and was ground properly for 2 h. The powder was initially heated at 400 1C for 6 h followed by room temperature quenching and grinding. This procedure was repeated five times in order to achieve homogeneous mixture of green powder with smaller particle size. Green powder was palletized after mixing with freshly prepared Poly Vinyl Alcohol (PVA) as binder. Cylindrical pellets with approximately 2 mm thickness and 10 mm diameter were prepared by hydraulic press having 30 ton base capacity with a pressure of 10 ton/cm2. Pressed pellets were slowly heated (50 1C/h) to 600 1C and kept for 8 h for slow release of PVA binder from the pellets using high-temperature programmable (Eurotherm controller, Model: 2404) vacuum furnace. A slow step sintering schedule (Fig. 1) was adopted up to 1300 1C to obtain highly dense ZnO bulk ceramics. X-ray diffraction (XRD) and scanning electron microscopy studies are carried out by Bruker D8 Advance X-ray Diffractometer and SEM (Philips FEG Xl’30), respectively for structural and morphological analyses. X-ray photoelectron spectroscopy and micro-Raman studies are carried out for the compositional analysis and also to confirm the presence of defects. All the hysteresis measurements of the samples were taken at low temperature (10 K) and room temperatures (300 K) using superconducting quantum interference device (SQUID) magnetometer with maximum field of 1 T.
3. Results and discussions Fig. 1. Slow step sintering schedule (SSSS) of pure ZnO.
Fig. 2 shows the XRD pattern of ZnO bulk ceramics as prepared and subsequently sintered at 850 and 1300 1C, respectively. The
Fig. 2. X-ray diffraction pattern of (a) ZnO as prepared, (b) ZnO sintered at 850 1C and (c) ZnO sintered at 1300 1C.
Table 1 Crystallite sizes, lattice parameter of ZnO (virgin) and ZnO bulk samples sintered at 850 1C and 1300 1C. Sample
Avg. crystallite size (nm)
Lattice parameters
c/a
ZnO (virgin) ZnO (850 1C) ZnO (1300 1C)
107.954 103.1786 97.53,414
a¼ 3.2450 c¼ 5.2000 a¼ 3.2410 c¼ 5.2019 a¼ 3.26 c¼ 5.255
1.60,246 well matched to the value of the reported 1.6025 (O) 1.6050 1.6119
ARTICLE IN PRESS 2494
J. Das et al. / Physica B 405 (2010) 2492–2497
peaks in the figure are indexed to hexagonal wurtzite structure of ZnO belonging to space group P63 mc (JCPDF # 890510). Preferred orientation is along (1 0 1) as expected for bulk ZnO. No evidence of any other phases is clear in the pattern. The average crystallite sizes are calculated using Scherrer’s formula and presented in Table 1. The XRD pattern showed a distinct decrease in crystallite size with an increase in crystalline order with sintered temperature. Lattice parameters of ZnO as prepared and sintered at 850 and 1300 1C, respectively, are calculated using the formula 1 4 h2 þ hk þk2 l2 þ 2 ¼ 3 a2 c d2hkl
Fig. 3. A plot showing the variation in c/a ratio with sintering temperature.
which are presented in Table 1. It is evident that the c/a ratio increases with an increase in sintering temperature (shown in Fig. 3), which is in agreement with the value of pure ZnO [Powder Diffraction File 36-1451 for hexagonal ZnO (1997)]. The crystallite size decreases with an increase in lattice parameter might result from the vacancies associated with the structure [21]. To study the physical densification of bulk ZnO ceramics processed by SSSS, micrograph study of the bulk surface was performed by SEM. Figs. 4a and b show the microstructure photograph of pure ZnO bulk ceramics sintered at 850 and 1300 1C for 24 h. It is seen that average microcrystallites of 8 mm dimension are grown with triangular grain boundary junction including the presence of multiple sub-micron pores of irregular structure (shown in Fig. 4c) in ZnO bulk ceramics sintered at 850 1C. The dimension of the pores is of the order of E1 mm. During grain growth process at elevated temperature (i.e., at around 1300 1C), mass flow to respective pores at atomic level occurs decreasing the number of pores and hence enhancing the physical densification. This is evident in Figs. 4b and d, which showed clean and clear triangular grain junction relating to three well-shaped uniform microcrystallites. This particular sample showed highest density as well as ferromagnetism at room temperature. A wide survey scan of XPS spectra is taken in the range 0–1100 eV as shown in Fig. 5a. No extra peak corresponding to any magnetic impurities other than Zn and O is observed in the figure. Fig. 5b shows the XPS spectra of C 1s. The 284.6 eV binding energy of C 1s is usually used as an internal reference in the spectrum [22]. The spectra corresponding to O 1s core and Zn 2p core level of ZnO bulk sample sintered at 1300 1C is calibrated by binding energy of C 1s (284.6 eV) as the internal reference spectra [22] and is shown in Figs. 5c and d. The O 1s spectrum is fitted with three Gaussian peaks having three binding energy components as per the literature [23]. The lower binding energy component at 529.44 eV is attributed to the wurtzite structure of
Fig. 4. SEM photographs of ZnO bulk ceramics sintered at (a) 850 1C, (b) 1300 1C, (c) presence of multiple pores in 850 1C-sintered ZnO ceramics and (d) the decrease in no. of pores due to grain growth in 1300 1C-sintered ZnO ceramics.
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
2495
Fig. 5. (a) Wide survey X-ray photoelectron spectrum of ZnO, (b) C 1s spectrum of ZnO, (c) O 1s spectrum of ZnO and (d) Zn 2p spectrum of ZnO.
hexagonal Zn2 + ion of the metal oxide [24]. The medium and higher binding energy components at 530.85 and 531.37 eV, respectively, are assigned to the oxygen deficient region within the ZnO matrix [24]. The asymmetry nature of the curve reflects the variation of oxidation number of oxygen. The Zn 2p spectrum in Fig. 6d shows a doublet whose binding energies are 1022.54 and 1045.69 eV and can be identified as Zn2p3/2 and Zn2p1/2 lines, respectively. The binding energy difference between the two lines is 23.15 eV, which is well lying within the standard reference value of ZnO [22]. The binding energy and the binding energy difference value calculated from the XPS study show that Zn atoms are in +2 oxidation state. At this stage, it is not possible to predict the presence of Zn interstitial defects from Zn 2p spectra, which may be established only by studying the Auger peak of Zn. The result obtained from micro-Raman analysis of pure ZnO bulk ceramics sintered at 1300 1C by SSSS for 24 h is shown in Figs. 6a and b. We observed slightly different spectra by focusing the beam at the grain boundary junctions and at the centre of the grain, respectively, with an aim to study the presence of defects attributed to intrinsic impurity and vacancy of anions and cations. 4 As known, the wurtzite semiconductor ZnO belongs to C3v (P63 mc) space group with two formula units per primitive cell,
where all atoms occupy C3n sites. The Raman active zone-centre optical phonons predicted by the group theory are A1+2E2 + E1+ 2B1. The B1 modes are silent modes, the A1 and E1 modes are polar and exhibit different frequencies for the transverse-optical (TO) and longitudinal-optical (LO) phonon modes. Both A1 and E1 are Raman and infrared active, whereas the E2 modes are nonpolar having two frequencies: E2 (high) associated with oxygen displacement and E2 (low) associated with Zn sub-lattice and are Raman active only [25,26]. As seen in the figure, the intensity of the peak corresponding to E2 (high) vibration mode, which is the band characteristic of the wurzite phase of ZnO and strongly depend on the isotopic composition of ZnO, is high and is centered at 436 cm 1, when the laser beam is focused at the centre of the grain. But its intensity is comparatively low and found to be centered at 433 cm 1 when the beam is focused at the grain boundary junction. Red shift in E2 (high) phonon mode of about 3–6 cm 1 is observed in both the spectra as compared to 439 cm 1 of ZnO bulk single crystal. The reason of such red shift due to the local heating effect [27] by the laser source can be ruled out as the excitation laser is too weak. We emphasize that the shift is only due to the defects produced during the sample processing by our SSSS method. Nonpolar E2 (high) phonon mode is usually associated with
ARTICLE IN PRESS 2496
J. Das et al. / Physica B 405 (2010) 2492–2497
Fig. 6. Raman spectra of ZnO bulk ceramic sample sintered at 1300 1C (a) Beam focussed at the centre of the grain and (b) beam focussed at the grain boundary.
Fig. 7. Magnetization vs. magnetic field (M–H) curves (a) at T¼ 10 K and (b) at T¼ 300 K for ZnO bulk samples sintered at 1300 1C.
oxygen atom displacement [28]. Hence, we assume that oxygen vacancy and host disorders are produced in ZnO during processing. The oxygen vacancies produced decreases the average atomic mass and causes mass fluctuation at oxygen position, resulting in the shifting of E2 line centre towards lower frequency. Again, a red shift of around 6 cm 1 in the E2 peak has been noted in the case of grain boundary junction as compared to the red shift of 3 cm 1 in case of beam focused at the centre of the grain. The larger red shift in E2 (high) phonon mode in case of laser beam focused at the grain boundary suggests that vacancies created during physical densification process while sintering, are mostly populated at the grain boundary (GB). Apart from this, distinct broader peaks along with shoulder peaks (SP) of considerable intensities are detected at around 575 cm 1 range in both the spectra, which correspond to 1LO (A1) mode. Normally, Raman peaks observed in between 570 and
590 cm 1 are considered to be associated with structural disorders, such as oxygen vacancy, Zn interstitial and their combination, due to the strong dependence on the oxygen stoichiometry [29,30]. Chen et al. [31] regarded this peak as a clear defect-induced mode. Our observation of broad shoulder peaks of almost same intensity associated with this 1LO peak clearly revealed the presence of defects or clusters of defects in the specimen, which might be created due to the outgasing of Zn during slow step heating of the sample. We speculate that vacancies created inside the bulk ceramics interacted with other nearby vacancies forming clusters, which is also the cause of the appearance of the above-mentioned broad SP associated with 1LO peak. Hysteresis behavior (M–H) of sintered bulk ZnO ceramics (processed at 1300 1C for 24 h) at 10 and 300 K is presented in Fig. 7 after correcting necessary background diamagnetic
ARTICLE IN PRESS J. Das et al. / Physica B 405 (2010) 2492–2497
subtraction. A clear M–H loop (inset) with saturation magnetization Ms ¼1.88 10 3 emu/g, coercivity Hc ¼34.44 Oe and remanent magnetization ¼8.07 10 5 emu/g is observed at 300 K (Fig. 7), indicating that pure ZnO bulk ceramics processed at 1300 1C by our SSSS method is ferromagnetic at RT, but becomes antiferromagnetic at low temperature (i.e., at 10 K) as clearly observed in the figure. The antiferromagnetic behavior is only due to the isolation of defects at low temperature whereas at room temperature there is a substantial enhancement in the ferromagnetic contribution in the sample, which is possibly due to the formation of clusters of defects in the sample when sintered at an elevated temperature of 1300 1C. As we have not incorporated any kind of magnetic impurities, this unconventional ferromagnetism observed here can definitely be assigned to the defects generated during material processing of this monoelemental (Zn) oxide. As known, contribution of defectinduced magnetization in the ZnO system is limited to oxygen vacancy (VO), oxygen interstitial (IO), Zn vacancy (VZn) and Zn interstitials (IZn), which are all non-magnetic, it is argued that only oxygen vacancy or zinc vacancy is responsible for RTFM. However, it has been explained by Banerjee et al. [10] that, in case of single O vacancy, at each of the next four Zn ions one sp3 hybrid of 4s and 4p orbital are found, which points towards the vacancy. As per their calculation a neutral O vacancy has two electrons in one state out of the four (one a1 and three degenerate t2), which is a singlet with zero spin, because of the splitting between a1 and t2 levels greater than Hund’s coupling (1 eV). Hence a single neutral oxygen vacancy cannot have a net magnetic moment associated with it. On the other hand, for a cluster of more than three O vacancies, the splitting is lower than Hund’s coupling, giving rise to a considerable net moment for the cluster. It is speculated by the researchers that heating at elevated temperature, may cause the O vacancies to migrate within the material and form clusters, favoring the decrease in the strain field energy ([10] and the Refs therein). Oxygen deficiency is manifested in the form of oxygen vacancy (deep donors) and zinc interstitials (shallow donors). It is suggested that the unpaired electron trapped in oxygen vacancies (F centre) are polarized to give room temperature ferromagnetism. Upon heating at elevated temperature, the ferromagnetic property was enhanced. It is also shown that for oxygen nonstoichiometry of 0.1% the magnetic moment contribution is considerably significant. It is understood that native defects play an essential role in understanding the behavior as well as for the successful application of any semiconductor and as such, in ZnO, specific defects have long been believed to play a more important role as its n-type conductivity has often been explained by the native point defects like oxygen vacancies and zinc interstitials [32]. We emphasize that during the processing, unreacted zinc outgases and annihilates the oxygen octahedron and tetrahedron networks and finally produces the above-mentioned defects and cluster of defect molecules, which are responsible to induce ferromagnetism in the slow step sintered specimen. Our microRaman as well as XPS study also clearly supported the presence of such defect structures in the sintered ceramics. However, the exact explanation of the ferromagnetic contribution by vacancy clusters needs to be studied more extensively both experimentally and theoretically.
4. Conclusion In conclusion, we emphasize that our experimental observation presents an alternate material preparation process for achieving ferromagnetic character in pure ZnO semiconductor
2497
without any magnetic cation substitution. Observation of defectinduced ferromagnetism with only small oxygen vacancies in bulk ZnO ceramics processed by SSSS at high temperature by us is definitely a challenge to explore and find the real explanation, which would be propitious for the development of multifunctional devices using ZnO. However, we also strongly urge that more research work in processing pure ZnO bulk ceramics (especially for bulk target for thin film materials) is essential along with satisfactory theoretical interpretation in this wellstudied material, which would open up a different avenue for magneto-optical applications.
Acknowledgement Authors are highly grateful to the ex-Director IMS, Prof. S.N. Behera for his constant guidance and encouragement to carry out research work at IMS, BBSR. We are also highly thankful to Mr. Santosh Choudhury to carry out the XPS study of our sample.
References [1] T. Dietl, H. Ohno, F. Matsukura, Science 287 (2000) 1019. [2] X.M. Cheng, C.L. Chien, J. Appl. Phys. 93 (2003) 7876. [3] S.J. Han, T.H. Jang, Y.B. Kim, B.G. Park, J.H. Park, Y.H. Jeonga, Appl. Phys. Lett. 83 (2003) 920. [4] T. Fukumura, Z. Jin, M. Kawasaki, T. Shono, T. Hasegawa, S. Koshihara, H. Koinuma, Appl. Phys. Lett. 78 (2001) 958. [5] A. Tiwari, C. Jin, A. Kvit Kumar, J.F. Muth, J. Narayana, Solid State Commun. 121 (2002) 371. [6] S.W. Jung, S.J. An, G.C. Yi, C.U. Jung, S. Lee, S. Cho, Appl. Phys. Lett. 80 (2002) 4561. [7] P. Sharma, A. Gupta, K.V. Rao, F.J. Owens, R. Sharma, R. Ahuja, J.M. Osorio Gullien, B. Johansson, G.A. Gehring, Nat. Mater. 2 (2003) 673. [8] J. Luo, J.K. Liang, Q.I. Liu, F.S. Liu, Y. Zhang, B.J. Sun, G.H. Rao, J. Appl. Phys. 97 (2005) 086106. [9] W. Chen, L.F. Zhao, T.Q. Wang, J.H. Miao, S. Liu, Z.C. Xia, S.L. Yuan, Solid State Commun. 134 (2005) 827. [10] S. Banerjee, M. Mandal, N. Gayathri, M. Sardar, Appl. Phys. Lett. 91 (2007) 182501. [11] K. Potzer, Z. Shengqiang, J. Grenzer, M. Helm, Fassbender, Appl. Phys. Lett. 92 (2008) 182504. [12] Q. Xu, H. Schmidt, S. Zhou, K. Potzger, M. Helm, H. Hochmuth, M. Lorenz, A. Setzer, P. Esquinazi, C. Meinecke, M. Grandmann, Appl. Phys. Lett. 92 (2008) 082508. [13] X. Zuo, S.D. Yoon, A. Yang, W.H. Duan, C. Vittoria, V.G. Harris, J. Appl. Phys. 105 (2009) 07C508. [14] H. Ohno, Science 281 (1998) 951. [15] T. Jungwirth, J. Sinova, J. Masek, J. Kucera, A.H. MacDonald, Rev. Mod. Phys. 78 (2006) 809. [16] J.M.D. Coey, Solid State Sci. 7 (2005) 660. ¨ [17] For a review, see G. Heiland, E. Mollwo, and F.Stockmann, in Solid State Physics, edited by F. Seitz and D. Turnbull, Academic, New York, vol 8, 1959. p. 191. [18] D.A. Schwartz, D.R. Gamelin, Adv. Mater. 16 (2004) 2115. [19] R. Radovanovic, D.R. Gamelan, Phys. Rev. Lett. 91 (2003) 157202. [20] S.A. Wolf, D.D. Awschalom, R.A. Buhraman, Science 294 (2001) 1488. [21] A. Sundaresan, R. Bhargavi, N. Rangarajan, U. Siddesh, C.N.R. Rao, Phys. Rev. B 74 (2006) 161306. [22] C.D. Wagner, W.M. Riggs, L.E. Davis, J.F. Moulder, G.E. Muilenberg, in: Handbook of X-ray Photoelectron Spectroscopy, Perkin Elmer, Eden Prairie, 1979. [23] M. Chen, X. Wang, Y.H. Yu, Z.L. Pei, X.D. Bai, C. Sun, R.F. Huang, L.S. Wen, Appl. Surf. Sci. 158 (2000) 134. [24] J.C.C. Fan, J.B. Goodenough, J. Appl. Phys. 48 (1977) 3524. [25] N. Ashkelon, et al., J. Appl. Phys. 93 (2003) 126. [26] J.F. Scott, Phys. Rev. B 2 (1970) 1209. [27] K.A. Alim, V.A. Fonoberov, A.A. Balandin, Appl. Phys. Lett. 86 (2005) 053103. [28] J. Serrano, F.J. Manjon, A.H. Romero, F. Widulle, R. Lauck, M. Cardona, Phys. Rev. Lett. 90 (2003) 055510. [29] R. Cusco, E.A. Llado, et al., Phys. Rev. B 75 (2007) 165202. [30] S.H. Jeong, J.K. Kim, B.T. Lee, J. Phys. D: Appl. Phys. 36 (2003) 2017. [31] Z.Q. Chen, et al., J. Appl. Phys. 97 (2005) 013528. [32] Q. Wang, Q. Sun, G. Chen, Y. KawaZoe, Puru Jena, Phys. Rev. B 77 (2008) 205411.