Weak ferromagnetism in textured Zn1−x(TM)xO thin films

Superlattices and Microstructures 39 (2006) 334–339 www.elsevier.com/locate/superlattices

Weak ferromagnetism in textured Zn1−x(TM)x O thin films H. Schmidta,∗, M. Diaconua, H. Hochmutha, M. Lorenza, A. Setzera, P. Esquinazia, A. Pöppla, D. Spemanna, K.-W. Nielsenb, R. Grossb, G. Wagnerc, M. Grundmanna a Institut für Experimentelle Physik II, Universität Leipzig, Linnéstrasse 5, 04103 Leipzig, Germany b Walther-Meißner-Institut für Tieftemperaturforschung, Bayerische Akademie der Wissenschaften,

Walther-Meißner-Str. 8, 85748 Garching, Germany c Universität Leipzig, Institut für Mineralogie, Kristallographie und Materialwissenschaft,

Scharnhorststrasse 20, 04275 Leipzig, Germany Available online 8 September 2005

Abstract Room-temperature ferromagnetic Zn1−x (TM)x O semiconductor thin films have been grown by pulsed laser deposition (PLD) on c-plane sapphire substrates and the effect of the substitution of Zn2+ ions by transition metal ions (Mn2+ or Ti2+ ) has been investigated. In contrast to ZnTiO thin films, the grain size of the ZnMnO thin films can be controlled in a wide range by the PLD growth conditions. The magnetic properties of the Zn1−x (TM)x O films have been investigated by a superconducting quantum interference device (SQUID). For semi-insulating ZnMnO films with an optimized grain size, we observed a weak ferromagnetism for temperatures up to 400 K [M. Diaconu, H. Schmidt, H. Hochmuth, M. Lorenz, G. Benndorf, J. Lenzner, D. Spemann, A. Setzer, K.-W. Nielsen, P. Esquinazi, M. Grundmann, Thin Solid Films 486 (2005) 117–121], while the ZnTiO films were only paramagnetic or superparamagnetic. This is expected with respect to the micromagnetic model. © 2005 Elsevier Ltd. All rights reserved. PACS: 61.72.V; 75.50.P; 61.16.C Keywords: Dilute magnetic semiconductors; ZnO; Mosaicity; Micromagnetism

∗ Corresponding author. Tel.: +49 (0) 341/97 32666; fax: +49 (0) 341/97 32668.

E-mail address: [email protected] (H. Schmidt). 0749-6036/$ - see front matter © 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.spmi.2005.08.059

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1. Introduction Dilute magnetic semiconductors (DMS) in which magnetic ions substitute cations of the host semiconducting material are assumed to be ideal systems for spintronics. There are two major criteria for selecting the most prospective materials for semiconductor spintronics: (i) Ferromagnetism should be retained above room temperature. (ii) The electrical and optical properties of ferromagnetic semiconductors should allow for spin manipulation. Based on theoretical studies for p-conducting II–VI DMS [2], attention was focussed on wide band gap semiconductors as prospective DMS for achieving magnetic ordering at ambient temperatures. In particular, the possibility of producing ZnO-based DMS has been explored. The intrinsically n-conducting, wide band gap II–VI semiconductor ZnO is already widely used, for example in surface acoustic wave (SAW) devices, gas sensors, microactuators, and conducting transparent layers. Since metallic Mn or Ti is paramagnetic, any ferromagnetism detected in ZnTMO films with TM = Mn or Ti cannot be due to Mn or Ti clusters possibly formed during the growth process. Ferromagnetism near or above room temperature has already been reported for Mn-alloyed ZnO pellets and thin films [1,3,4] whereas in other studies only paramagnetic behavior was reported [5–9]. It was suggested that the weak ferromagnetism observed in semi-insulating, Mn-alloyed ZnO films is due to antiferromagnetically coupled, but nonparallel (canted) magnetic moments [10]. Furthermore, Venkatesan et al. [11] also reported on ferromagnetism in Ti-alloyed ZnO thin films. Several investigations across the 3d transition metal series in ZnO underline the different electrical and magnetic behavior of Mn and Ti in ZnO. The contribution of 3d transition metal ions to the electronic conduction in ZnO has been elucidated by first-principle pseudopotential calculations [12] which predict that Mn generates isolated energy levels within the energy gap of ZnO, whereas Ti generates energy levels in the conduction band region of ZnO. Ab initio studies [13] explain the ferromagnetism in ZnMnO by a hole-mediated double exchange mechanism and predict a zero total magnetic moment of Ti2+ , thus precluding magnetic ordering in p-conducting ZnTiO. On the other hand, the phenomenological spin–split donor impurity-band model [11] predicts a nearly zero-valued magnetic moment of Mn ions and a large total magnetic moment of Ti ions in n-conducting ZnO. In our experimental study we observed grain-size dependent magnetic properties, namely weak ferromagnetism in ZnMnO with larger grains and superparamagnetism in ZnTiO with smaller grains. Existing theories on DMS do not include such effects. Thus, the clarification of the relation between mosaicity and magnetism in ZnO-based DMS is of fundamental interest. The solubility of Mn2+ into the ZnO matrix is relatively high and amounts to x < 0.35 [14]. Because of the similar ionic radii of Mn2+ and Zn2+ , the crystalline quality of ZnMnO is comparable to that of pure ZnO. On the other hand, the analysis of the (0002) peaks in the θ –2θ XRDscans of ZnTiO thin films provides evidence for a deteriorated crystalline quality [15] and hints toward a small grain size in thin ZnTiO films. This observation is expected because the volume of TiOn polyhedra decreases with the coordination number n and the valence of Ti. Furthermore, the preparation of semi-insulating, magnetic ZnTMO films is described in Section 2 and the experimental investigations on magnetic properties are discussed in Section 3. Finally, we give a phenomenological explanation for the coercivity mechanism

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and emphasize the importance of film mosaicity. One possible future application of magnetic, semi-insulating ZnTMO films is magnetic piezoelectrics [16]. 2. Sample preparation The ZnTMO films have been grown under oxygen atmosphere on 10 × 10 mm2 c-plane sapphire substrates by pulsed laser deposition (PLD) using a KrF excimer laser [17]. The distance between PLD target and substrate was 10 cm; 30 000 laser pulses with an energy density of 2 J cm−2 have been used to grow ca. 1 µm thick films. The ZnMnO-PLD targets with a nominal Mn content of 3 and 10 at.% were prepared by mixing and pressing appropriate amounts of ZnO (5N5) and MnO2 (5N) powders. The ZnTiO–PLD targets with a Ti content of 4, 5, and 10 at.% were prepared from ZnO (5N5) and TiO2 . Afterwards, the ZnMnO and ZnTiO target material has been sintered for 16 h in air at 500 or 700 ◦C or at 1100 ◦C, respectively. For the PLD growth optimization, we varied the oxygen partial pressure between 0.1 and 30 Pa and the growth temperature between 200 and 800 ◦C. The resistivity of the ZnTMO films increases exponentially with pressure and growth temperature and only semi-insulating ZnTMO films revealed magnetic domain formation. Generally, ZnO films on c-plane sapphire substrates grow with the c-axis perpendicular to the substrate surface and exhibit a columnar structure. The size and misorientation of the columns varies as a function of the PLD growth conditions. The mosaicity of magnetic ZnMnO and ZnTiO films can be recognized from the surface structure, where the formation of columnar structures with an average diameter of 170 nm (not shown here) and 25 nm (Fig. 1(a)), respectively, is clearly visible. In contrast to ZnTiO thin films, the grain size in ZnMnO thin films can be controlled in a wide range from 50 nm up to 320 nm by the PLD growth conditions (Fig. 2). The grain size in ZnTiO films on c-sapphire is smaller than 50 nm. The TM content in the magnetic Znx TM1−x O films has been investigated by combined Rutherford backscattering spectrometry (RBS) and particle-induced X-ray emission (PIXE) using 1.2 MeV protons. The presence of secondary phases is excluded on the basis of X-ray diffraction (XRD) measurements. On average, the chemical composition of the 3 and 10 at.% ZnMnO and of the 4 at.% ZnTiO target is transferred into the corresponding thin films. For the 5 at.% ZnTiO target the average transfer factor amounts to 2. 3. Experimental Because the magnetic signal detected by SQUID measurements on magnetic semiinsulating ZnTMO films is very weak, in addition to the bulk-sensitive SQUID technique we also used the surface-sensitive magnetic force microscopy (MFM) technique to observe magnetic domain formation in ZnMnO films [1] and in ZnTiO films (Fig. 1(b)). Revealing a periodicity of about 1.7 µm, the experimentally observed magnetic domains are much larger than the grains with a diameter of 25 nm. In a previous paper we showed that the coercivity in ferromagnetic ZnMnO films passes through a maximum for an optimal grain size of 140 nm [1] and we assumed a coercivity mechanism like in granular metals, where the coercivity passes through a maximum at a grain size comparable to the domain-wall width [18]. The temperature dependent magnetization measured on two magnetic films

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Fig. 1. 20 × 20 µm2 (a) AFM image and (b) MFM image of a magnetic ZnTiO film with a Ti content of 9.9 at.% revealing small grains and magnetic domains, respectively.

Fig. 2. Grain diameter as a function of growth temperature measured by AFM on different ZnMnO and ZnTiO thin films. The ZnTMO films discussed in Figs. 1 and 3 are labelled by ZnMnO and ZnTiO.

with a comparable TM concentration of 9.1 at.% Mn and 9.9 at.% Ti is shown in Fig. 3(a) and (b), respectively. The shape of field cooled (FC) and zero-field cooled (ZFC) curves

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Fig. 3. Temperature dependent 500 Oe FC and ZFC magnetization measured on a (a) ZnMnO thin film with a Mn content of 9.1 at.% and on a (b) ZnTiO thin film with a Ti content of 9.9 at.%.

and the temperature dependence of their nonzero difference can be used for deciding which type of magnetization, namely ferromagnetic, spin-glass, superparamagnetism, or paramagnetic has been observed. For superparamagnets [19] or spin-glasses [20], the ZFC curves show a peak below the blocking temperature or freezing point, respectively. We observed a nonzero difference between the FC and ZFC for magnetic ZnMnO (Fig. 3(a)) and ZnTiO (Fig. 3(b)) films at an applied magnetic field of 500 Oe. From the temperature dependent shape of the FC and ZFC magnetization curves we conclude that the ZnMnO films show ferromagnetic ordering (Fig. 3(a)) and that the ZnTiO thin films show superparamagnetic ordering (Fig. 3(b)). It should be noted that the blocking temperature increases with increasing volume of superparamagnetic particles. Therefore, the broad ZFC peak in Fig. 3(b) indicates a nonuniform size distribution of the superparamagnetic particles in the ZnTiO film.

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4. Discussion and conclusions We used the PLD technique to grow ZnTMO thin films and we succeeded in obtaining, ferromagnetic or superparamagnetic, semi-insulating thin films at room temperature. One possible application of semi-insulating, magnetic ZnTMO films is magnetic piezoelectrics [16]. Magnetic force microscopy measurements on the magnetic ZnTMO films revealed homogeneous magnetic domain formation and point toward the intrinsic nature of the weak ferromagnetism or superparamagnetism observed. We detected a clear correlation between the columnar structure in ferromagnetic ZnMnO films measured by AFM and the coercive field measured by SQUID. Temperature dependent magnetization measurements revealed ferromagnetic ordering in ZnMnO and superparamagnetic ordering in ZnTiO. This can be explained in the micromagnetic model because, in contrast to ZnMnO thin films, the mosaicity of ZnTiO thin films cannot be controlled over a wide range. It would be instructive to strengthen the exchange interaction via itinerant carriers and to increase the conductivity by additional doping of the ZnTMO films. Acknowledgement The funding from BMBF (code: 03N8708) is gratefully acknowledged by Schmidt, Diaconu and Hochmuth. References [1] M. Diaconu, H. Schmidt, H. Hochmuth, M. Lorenz, G. Benndorf, J. Lenzner, D. Spemann, A. Setzer, K.-W. Nielsen, P. Esquinazi, M. Grundmann, Thin Solid Films 486 (2005) 117–121. [2] T. Dietl, H. Ohno, F. Matsukura, J. Cibert, D. Ferrand, Science 287 (2000) 1019–1022. [3] P. Sharma, A. Gupta, K.V. Rao, F.J. Owens, R. Sharma, R. Ahuja, J.M. Osorio Guillen, B. Johansson, G.A. Gehring, Nature Mater. 2 (2003) 673–677. [4] D.P. Norton, S.J. Pearton, A.F. Hebard, N. Theodoropoulou, L.A. Boatner, R.G. Wilson, Appl. Phys. Lett. 82 (2003) 239–241. [5] K. Ueda, H. Tabata, T. Kawai, Appl. Phys. Lett. 79 (2001) 988–990. [6] T. Fukumura, Z. Jin, M. Kawasaki, T. Shono, T. Hasegawa, S. Koshihara, H. Koinuma, Appl. Phys. Lett. 78 (2001) 958–960. [7] A. Tiwari, C. Jin, A. Kvit, D. Kumar, J.F. Muth, J. Narayan, Solid State Commun. 121 (2002) 371–374. [8] S.S. Kim, J.H. Moon, B.T. Lee, O.S. Song, J.H. Je, J. Appl. Phys. 95 (2004) 454–459. [9] X.M. Cheng, C.L. Chien, J. Appl. Phys. 93 (2003) 7876–7878. [10] G. Lawes, A.S. Risbud, A.P. Ramirez, R. Seshadri, Phys. Rev. B 71 (2005) 045201-1–045201-5. [11] M. Venkatesan, C.B. Fitzgerald, J.G. Lunney, J.M.D. Coey, Phys. Rev. Lett. 93 (2004) 177206-1–177206-4. [12] Y. Imai, A. Watanabe, J. Mater. Sci., Mater. Electron. 15 (2004) 743–749. [13] K. Sato, H. Katayama-Yoshida, Japan. J. Appl. Phys. 39 (2000) L555–L558. [14] T. Fukumura, Z. Jin, A. Ohtomo, H. Koinuma, M. Kawasaki, Appl. Phys. Lett. 75 (1999) 3366–3368. [15] Y.R. Park, K.J. Kim, Solid State Commun. 123 (2002) 147–150. [16] N.A. Spaldin, Phys. Rev. B 69 (2004) 125201-1–125201-7. [17] M. Lorenz, E.M. Kaidashev, H. von Wenckstern, V. Riede, C. Bundesmann, D. Spemann, G. Benndorf, H. Hochmuth, A. Rahm, H.C. Semmelhack, M. Grundmann, Solid-State Electron. 47 (2003) 2205–2209. [18] J.F. Löffler, W. Wagner, G. Kostorz, J. Appl. Cryst. 33 (2000) 451–455. [19] J.G. Moore, E.J. Lochner, C. Ramsey, N.S. Dalal, A.E. Stiegman, Angew. Chem., Int. Edn 42 (2003) 2741–2743. [20] Young-II Jang, F.C. Chou, Y.-M. Chiang, Appl. Phys. Lett. 74 (1999) 2504–2506.