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Mechanochemical synthesis of yttrium manganite
Journal of Alloys and Compounds 552 (2013) 451–456
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Mechanochemical synthesis of yttrium manganite M. Pocˇucˇa-Nešic´ a, Z. Marinkovic´ Stanojevic´ a,⇑, Z. Brankovic´ a, P. Coticˇ b, S. Bernik c, M. Sousa Góes d, B.A. Marinkovic´ e, J.A. Varela d, G. Brankovic´ a a
Institute for Multidisciplinary Research, University of Belgrade, Kneza Višeslava 1a, 11030 Belgrade, Serbia Institute of Mathematics, Physics and Mechanics, Jadranska 19, 1000 Ljubljana, Slovenia c Jozˇef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia d Instituto de Química, Depto. Físico-Química, Universidade Estadual Paulista-UNESP, R. Prof. Francisco Degni, n. 55, 14800-900 Araraquara SP, Brazil e Departamento de Engenharia de Materiais, Pontifícia Universidade Católica de Rio de Janeiro – PUC-Rio, Rua Marquês de São Vicente 225, Gávea, Rio de Janeiro, RJ, Brazil b
a r t i c l e
i n f o
Article history: Received 3 July 2012 Received in revised form 6 November 2012 Accepted 6 November 2012 Available online 16 November 2012 Keywords: Mechanochemistry Multiferroic materials Yttrium manganite Microstructure Magnetic properties
a b s t r a c t Yttrium manganite (YMnO3) is a multiferroic material, which means that it exhibits both ferromagnetic and ferroelectric properties, so making it interesting for a variety of technological applications. In this work, single-phase YMnO3 was prepared for the first time by mechanochemical synthesis in a planetary ball mill. The YMnO3 was formed directly from the highly activated constituent oxides, Y2O3 and Mn2O3, after 60 min of milling time. During prolonged milling, the growth of the particles occurred. The cumulative energy introduced into the system during milling for 60 min was 86 kJ/g. The X-ray powder-diffraction analysis indicates that the as-prepared samples crystallize with an orthorhombic (Pnma) YMnO3 structure. The morphology and chemical composition of the powder were investigated by SEM and FESEM. The magnetic properties of the obtained YMnO3 powders were found to change as a function of the milling time in a manner consistent with the variation in the nanocomposite microstructure. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Multiferroic materials that exhibit any type of long-range magnetic ordering, spontaneous electric polarization, and/or ferroelasticity in one phase have been the subject of numerous scientific studies in recent years. Their interesting properties offer possibilities for many different technological applications in non-volatile ferroelectric memories, sensors, transducers and microwave devices [1]. As possible multiferroic materials, rare-earth manganites, RMnO3 (R being rare earth or transition metal ion), have attracted a lot of attention recently. These manganese (III) oxides have an interesting structure that depends on the size of the R cation. Those with larger R cations crystallize in the perovskite orthorhombic structure (space group Pnma), while those with a smaller R cation radii crystallize in the hexagonal structure (space group P63cm). Among these materials YMnO3 (YMO) stands out as a proven multiferroic material and because of its intermediate size the Y3+ ion allows the formation of both structures: the hexagonal (h-YMO) structure, stable under atmospheric conditions, and orthorhombic (o-YMO), the metastable one. The literature data shows that the Néel temperatures (TN) for h-YMO and o-YMO are 80 and 42 K,
⇑ Corresponding author. Tel.: +381 11 2085842. E-mail address: [email protected] (Z. Marinkovic´ Stanojevic´). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.11.031
respectively, while the ferroelectric Curie temperatures (TC) are 914 K (h-YMO) and 31 K (o-YMO) [2–5]. Being the stable one, h-YMO can be prepared by various techniques. Nanosized powders are obtained by a solid-state reaction from Mn(III) and Y(III) oxides [6,7], a hydrothermal process [8] or chemical methods employing different chelating agents, like citric acid or EDTA [9,10]. Polycrystalline and epitaxial YMO thin films with a hexagonal structure can be obtained by means of molecular-beam epitaxy, laser ablation, sputtering methods, pulsed-laser deposition or the sol–gel process [11,12]. Physical methods can also be used for the preparation of orthorhombic YMO thin films. But in this case, an appropriate choice of perovskite substrates and/or the use of a seeding layer, could be the crucial factor for the stabilization of the metastable phase [11,13–15]. On the other hand, the characterization of o-YMO powders is very difficult, since the preparation and stabilization of this phase are very complex. Orthorhombic single crystals of YMO can be prepared by the transformation of the hexagonal phase at high pressures and high temperatures [2,16]. Apart from this treatment, recommended synthetic routes for powder preparation are soft chemical methods, with a thermal treatment at lower temperatures [17,18]. These methods can be very complex, since the precursor structure has to fulfill certain requirements. The Mn3+ ion in the orthorhombic YMO is six-coordinated and these surroundings should be preserved throughout the whole synthesis and thermal treatment, without transforming to the five-coordinated Mn3+
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ion in the hexagonal phase. Even if this can be accomplished, the resulting powder can be a mixture of both YMO phases [19,20]. To the best of our knowledge, no literature data has been reported on the mechanochemical synthesis of YMO powders. This synthetic method is based on solid-state reactions taking part while the reactants are being milled in a high-energy ball mill. During these reactions the particle size of the reactants is decreasing, the contact area between the particles is increasing and the nucleation starts at lower temperatures, which results in the formation of products with a smaller particle size. Intensive milling also provides high temperatures and high pressures for the start of the reaction and for the formation of the product [21]. In this paper mechanochemical synthesis was proposed as a possible method for the preparation of orthorhombic YMnO3, having in mind that by using this method some perovskite manganitebased materials have already been prepared without post-milling high-pressure or high-temperature treatments [21]. The milling conditions were investigated, and the characterization of the asprepared powders was performed. The main result of our work is the preparation of a pure orthorhombic YMnO3 phase after only 240 min of milling. 2. Experimental procedure For the mechanochemical synthesis of YMnO3, Y2O3 (Alfa Aeser, p.a. 99.9%) and Mn2O3 (Aldrich, p.a. 99%) were used as the precursors. The appropriate amounts of oxides, with an equimolar Y3+ and Mn3+ ratio, according to Eq. (1), were weighed, mixed and then dry milled in a planetary ball mill (Fritsch Pulverisette 5). The milling conditions were as follows: tungsten carbide cylindrical vials (V = 250 cm3, Dv = 75 mm) and balls (35 balls with d = 10 mm, q = 14.95 g/cm3); a ball-to-powder weight ratio of 30:1; an air atmosphere; a basic disc rotational speed of 325 rpm; a rotation speed of the discs with vials equal to 400 rpm; and a milling time of 60– 720 min.
Y2 O3 þ Mn2 O3 ! 2YMnO3
ð1Þ
Based on the milling times the samples were marked as YMO60, YMO120, YMO240, YMO360 and YMO720. The ball-impact energy was calculated (DEb⁄ 60 mJ/hit) according to the model derived by Burgio et al. [22] The experimentally calculated weight-normalized cumulative energies (Ecum) required for starting (after 60 min) and the complete (after 240 min) formation of the YMnO3 phase were 86 kJ/g and 344 kJ/g, respectively. The X-ray diffraction (XRD) data were obtained using a RigakuÒ RINT2000 diffractometer (42 kV 120 mA) with CuKa radiation (kka1 = 1.5405 Å, kka2 = 1.5443 Å, Ika1/Ika2 = 2) in the 2h range between 10° and 80°, step 0.02° (2h) with 10.0 s per point. To obtain the microstructural data and the quantitative phase compositions of the as-milled powders a Rietveld refinement [23] was performed using the General Structure Analysis System (GSAS) program [24] suite with the EXPGUI interface [25]. The peak-profile function was modeled using a convolution of the Thompson–Cox–Hastings pseudo-Voigt function (pV-TCH) [26], employing the asymmetry function described by Finger et al. [27], which accounts for the
Fig. 1. (a) XRD patterns for Y2O3 and Mn2O3 mixture milled between 60 and 720 min and (b) observed, calculated and difference profiles for the sample YMO360.
asymmetry resulting from the axial divergence. The bi-dimensional model for crystallite size described by Larson & Von Dreele [24] was used to account for the anisotropy in the half width of the reflections. Differential scanning calorimetry (DSC) and thermogravimetric analyses (TGA) of the as-prepared powders were performed in the temperature interval from 20 to 1000 °C, with a heating rate of 10 °C/min (model SDT Q-600). The powders were
Table 1 Results of structural analysis (unit-cell parameters, cell volume, density and crystallite size (hti)), mass percentages and agreement indexes at the end of the refinements for the powders obtained after milling for 60–720 min with Rietveld refinement. YMO60 a (nm) b (nm) c (nm) V (nm3) q (g/cm3) hti \ (nm) hti// (nm) s = hti \ /hti// wt.% YMnO3 wt.% Y2O3 wt.% Mn2O3 Rwp(%) RF2(%) S
0.5744(2) 0.7433(2) 0.5263(1) 0.22478(6) 5.670
44.7(7) 9.3(4) 45.9(8) 3.14 0.78 1.41
Rwp is the weighted index, S the goodness of fit ð¼
YMO120 0.57493(8) 0.74216(9) 0.52592(5) 0.22441(6) 5.677
99.81(3) 0.18(6) – 3.36 0.87 1.49 pffiffiffiffiffi x2 Þ, RF is the structure factor index [23].
YMO240
YMO360
YMO720
0.57507(7) 0.74190(8) 0.52610(5) 0.22446(5) 5.677 40 21 1.90 100 – – 3.13 0.60 1.40
0.57584(7) 0.74151(8) 0.52655(5) 0.22484(5) 5.667 16 20 0.80 100 – – 3.53 0.78 1.58
0.57630(6) 0.74182(8) 0.52653(5) 0.22510(5) 5.662 27 62 0.44 100 – – 3.22 0.58 1.44
M. Pocˇucˇa-Nešic´ et al. / Journal of Alloys and Compounds 552 (2013) 451–456 also characterized by SEM (SEM TESCAN Vega TS 5130MM), FE-SEM (JEOL, Model JSM-6701F) and particle size analyses (HORIBA LA-920 laser scattering particle size distribution analyzer). The magnetic measurements of the YMO samples were carried out with a SQUID MPMS-XL-5 magnetometer from Quantum Design. The zero-field-cooled (ZFC) and field-cooled (FC) magnetization vs. temperature curves were studied in the temperature range 2–300 K, while isothermal magnetization measurements were recorded between 50 and 50 kOe at two temperatures, 5 and 300 K.
If we take into account only single-phase YMO samples, the lattice parameter a and the cell volume (V) of the mechano-synthesized YMnO3 perovskite phase are significantly higher in comparison to the reported ICSD values (166217). The Rietveld refinement indicates that prolonging the milling time from 240 to 720 min results in an increase of the cell parameter a and the
3. Results and discussion The XRD patterns of the powders obtained after milling for different times are shown in Fig. 1a. It is clear that already after 60 min of milling time, the distinctive peaks of the orthorhombic phase of YMO are present: (2h = 25.73, 33.41, 48.46, and 60.50). The YMO60 XRD pattern shows that, apart from the orthorhombic YMO phase, there is also some quantity of starting oxides and an amount of amorphous phase (detectable through the pronounced background halo around the 2h = 33.41 peak). The XRD pattern for the YMO120 shows almost no traces of the starting oxides and a significantly reduced quantity of amorphous material. The patterns for the powders obtained after milling for 240, 360 and 720 min are almost identical. They have peaks that can only be attributed to the orthorhombic YMO phase, while the amorphous halo disappeared. These conclusions coming from a qualitative data analysis were additionally verified using the Rietveld-refinement method. The corresponding results for the sample YMO360 are illustrated in Fig. 1b. Quantitative phase and microstructural analyses results and the Rietveld agreement indexes after the YMO refinements are presented in Table 1. The quantitative phase analysis for the YMO60 powder sample shows that the crystallization of the YMO started after milling times as short as 60 min, although there are also some quantities of starting Y2O3 (9.3 wt.%) and Mn2O3 (45.9 wt.%) present in the powder. The quantities of these oxides are insignificant in the YMO120 sample, i.e., only 0.18 wt.% of Y2O3 and no traces of Mn2O3. It is clear that the milling times of 240, 360 and 720 min resulted in the preparation of pure YMnO3 with the orthorhombic structure, o-YMO, space group Pnma. The starting crystal-structure parameters used in the Rietveld quantitative data analysis were in accordance with the structures presented in the 166217 (YMnO3), 86815 (Y2O3) and 9091 (Mn2O3) data cards, available from the Inorganic Crystal Structure Database (ICSD). The Rietveld-analysis-based results proved that the pure, orthorhombic, metastable phase of YMnO3 can be prepared through mechanochemical synthesis.
Fig. 2. DSC–TGA curve for YMO240 sample.
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Fig. 3. Results of particle size distribution for YMO powders.
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Fig. 4. SEM micrographs for powder samples: (a) YMO120, (b) YMO360, (c) and (d) YMO720.
cell volume. The variation in the reported lattice parameter might result from the presence of non-stoichiometry as a consequence of the formation of point and lattice defects during the mechanochemical synthesis. The values of the crystallite sizes are relatively small (<65 nm), while the time of milling promotes the presence of a crystallite-shape (T) anisotropy. This is in accordance with some previous studies [28,29]. Fig. 2. shows the results of the thermal analysis for sample YMO240. It is clear that the weight loss during the thermal treatment is insignificant and it may be attributed to the loss of the always-present hygroscopic water and carbonates, or some small quantities of impurities. This result is in accordance with the idea that after 240 min of milling, pure o-YMO is formed and since it does not have any volatile components, only structural changes are expected to occur during heating. However, the absence of any distinctive peaks up to 1000 °C indicates the complete crystallization of o-YMO, while the phase transformation to stable h-YMO is postponed to higher temperatures. This is in accordance with the literature data, which suggest that the pure hexagonal phase can be prepared at 1100 °C or higher temperatures [18,30]. The results of the particle size analysis for the as-prepared powders of pure o-YMO are given in Fig 3. They reveal a bimodal particle size distribution for all the samples, except YMO120 and YMO240. It can be seen that the smallest particle sizes are present in the powder milled for 120 min, which is the sample with the narrowest particle size distribution. A characteristic of all the powders is the prevalence of a median particle size in the range between 1 and 3 lm. By prolonging the milling time, the particle size increases and the particle size distribution is broadened, suggesting the possibility of further particle
agglomeration, which is a consequence of the long milling times during the mechanochemical synthesis. The results of the microstructural analysis of the as-prepared powders are in accordance with the particle size analysis. From Fig. 4a, which shows an SEM micrograph of the YMO120 sample, it is clear that the particles have a similar, uniformly distributed size around 1 lm. However, prolongation of the milling time resulted in the formation of powders with larger particle sizes and with large agglomerates, even larger than 10 lm. SEM micrographs for both single particles and agglomerates, for the sample YMO720, are presented in Fig. 4c and d. An additional microstructural analysis of the sample YMO360 (Fig. 5), using the FE-SEM technique, offered a new insight into the inner microstructure of the particle. From this micrograph (Fig. 5) we can see that the as-prepared particles are dense and pore-free. Also, it is clear that the primary particles (crystallites) that are composing the secondary particles previously identified with the SEM and the particle size analyses are much smaller than 100 nm and that the as-prepared o-YMO powders obtained from the mechanochemical synthesis could be considered as nanopowders, being in accordance with the XRD results (Table 1). Fig. 6 shows the magnetization of the single-phase YMO samples as a function of magnetic field, M(H), under ZFC at a temperature of 5 K and YMO240 samples at 300 K. The absence of any hysteresis in the magnetization curves at room temperature is evidence for paramagnetic behavior, which confirms that the ferromagnetic component disappears above TN. The M vs. H curve at low temperatures (5 K) deviates from the linear relation, but the magnetization does not show saturation in a field up to 50 kOe. The minor magnetic hysteresis at 5 K (see inset of Fig. 6) and the
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Table 2 Different parameters deduced from the magnetization curves of mechanochemically synthesized YMO powder samples. Sample
TN (K)
Hc 103 (Oe)
hCW (K)
leff (lB)
YMO240 YMO360 YMO720
42 37 37
0.742 0.503 0.501
44 42 37
4.57 4.45 4.41
TN, Neel temperature (antiferro-paramagnetic transition); Hc, coercive field at 5 K; hCW, the paramagnetic Curie–Weiss temperature; and leff, effective magnetic moment.
Fig. 5. FE-SEM micrograph of a secondary particle of the sample YMO360, indicating nanosized dimensions of the primary particles (crystallites).
very low coercive fields (see Table 2) indicate that the samples are basically anti-ferromagnetic with weak ferromagnetism. The thermal evolution of the dc magnetic susceptibility of the single-phase YMO samples (YMO240, YMO360 and YMO720) is exhibited in Fig. 7. It is clear that initially the magnetization increases gradually over a broad region as the temperature decreases. The diffuseness of the magnetic transition probably indicates that an inhomogeneous magnetic state prevails in the transition region. The broadening of the magnetic transition is consistent with the observed broadening of the X-ray reflections and with the model of short-range anti-ferromagnetism and a spinglass state [31,32]. A typical cusp exists in the ZFC curve and the low-temperature ZFC and FC curves deviate substantially for all the samples. The phenomenon of the discrepancy between the ZFC and the FC magnetization curves is usually ascribed to the appearance of the spin-glass or cluster-glass state induced by the competition between the ferromagnetic and anti-ferromagnetic exchange interactions. In addition, the M(H) curve displays a small hysteresis at 5 K (see inset in Fig. 6), consistent with a conventional spin-glass system [8,17,33]. Moreover, in Fig. 7 we can see that the onset of magnetization occurs at about 50–70 K, which is above the Néel temperature (40 K) reported for bulk orthorhombic YMnO3 [6,17,31]. This shift could also reflect the strain-induced unbalance of the anti-ferromagnetic competing interactions in the samples [14].
Fig. 6. Isothermal magnetization curves of YMO samples measured at 5 and 300 K. The inset shows the enlarged version of M (H).
Fig. 7. Magnetization as a function of temperature, ZFC and FC cycles, measured under an applied field of 100 Oe for YMO240, YMO360 and YMO720 samples. The solid symbols denote ZFC, whereas the open symbols denote FC magnetization. Inset shows the thermal evolution of the inverse susceptibility for all the YMO samples.
The inset of Fig. 7 exhibits the reciprocal susceptibility vs. temperature of mechanochemically synthesized YMO samples. By fitting with the Curie–Weiss law, v = C/(T hCW), in the temperature range above T = 120 K, the values of the effective paramagnetic moments (leff) and the paramagnetic Curie–Weiss temperatures (hCW) could be obtained. Table 2. collects these parameters together with the TN of the single-phase YMO samples. TN was defined by the peak in the ZFC d(vT)/dT vs. T curve (not presented here). The negative value of the paramagnetic Curie–Weiss temperature indicates the dominance of the anti-ferromagnetic interaction in all the samples. The sets of Curie–Weiss parameters are in good agreement with the previously reported values for the orthorhombic YMnO3 [6,20,31]. The effective magnetic moment obtained from the reciprocal susceptibility is below the expected value for the spin-only Mn3+ ions in the ground state (lMn3+ = 4.90lB) (see Table 2). This could be a result of the covalence effects and could also be understood by considering that a small content of Mn4+ (lMn4+ = 3.87lB) with a smaller expected effective magnetic moment is present in the samples YMO240, YMO360 and YMO720. Probably, the presence of some Mn4+ ions results from the slight oxygen excess localized in the interstitial sites. In the stoichiometric YMnO3, the magnetic interaction of the two neighboring Mn3+ is mediated by the oxygen ions between them and is ascribed to the super-exchange mechanism with the resulting anti-ferromagnetic coupling of the moments. The non-stoichiometry-related defects could lead to the formation of Mn4+ states, which introduce some disorder in the crystal lattice and the subsequent magnetic disorder. These Mn4+ ions do not participate in the super-exchange interaction but open the double-exchange interaction channel whenever they are next
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to a Mn3+ ion. In that case, double-exchange Mn3+/4+–Mn3+/4+ ferromagnetic interactions can be introduced, thus leading to competing ferromagnetic-anti-ferromagnetic interactions that may behave as a spin-glass-like material [17,34]. In addition, we can speculate that the slight difference in the crystal structure due to the microstrain effect (induced by nonstoichiometry-related defects) might have an influence on the decrease of the magnetic transition temperature through the changes in the bond lengths and the angles between the magnetic and oxygen atoms. Because of that, the small depression in transition temperature (TN) in YMO360 and YMO720 can be attributed to the non-stoichiometric composition of this sample. 4. Conclusions Pure orthorhombic YMO powders were successfully prepared by mechanochemical synthesis in a planetary ball mill, after 240 min of milling, starting from Y2O3 and Mn2O3 precursors. According to the XRD analysis and the Rietveld refinement, the YMnO3 obtained after milling for 240 min had a orthorhombic perovskite structure (Space group Pnma) and the following lattice parameters: a = 0.57507 nm, b = 0.74190 nm, c = 0.52610nm. The mean particle size for the YMO powders milled for 240 min is 1.25 lm. The SEM analysis proved the agglomeration of powders with prolonged milling times; however, the FE-SEM analysis evidence that the primary particles (crystallites) are smaller than 100 nm. The irreversibility in the magnetic response and the small hysteresis at low temperature indicate that the samples obtained by mechanochemical synthesis are basically anti-ferromagnetic, with weak ferromagnetism. We have proposed that this behavior could be associated with a modification of the magnetic interactions in the samples due to the presence of non-stoichiometry-related defects. Acknowledgement This work was supported by Ministry of Education and Science of Republic of Serbia through project III45007 and COST Action MP0904. B. A. Marinkovic´ is grateful to CNPq for a Research Productivity Grant. References [1] M. Fiebig, J. Phys. D: Appl. Phys. 38 (2005) R123–R152.
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Mechanochemical synthesis of yttrium manganite M. Pocˇucˇa-Nešic´ a, Z. Marinkovic´ Stanojevic´ a,⇑, Z. Brankovic´ a, P. Coticˇ b, S. Bernik c, M. Sousa Góes d, B.A. Marinkovic´ e, J.A. Varela d, G. Brankovic´ a a
Institute for Multidisciplinary Research, University of Belgrade, Kneza Višeslava 1a, 11030 Belgrade, Serbia Institute of Mathematics, Physics and Mechanics, Jadranska 19, 1000 Ljubljana, Slovenia c Jozˇef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia d Instituto de Química, Depto. Físico-Química, Universidade Estadual Paulista-UNESP, R. Prof. Francisco Degni, n. 55, 14800-900 Araraquara SP, Brazil e Departamento de Engenharia de Materiais, Pontifícia Universidade Católica de Rio de Janeiro – PUC-Rio, Rua Marquês de São Vicente 225, Gávea, Rio de Janeiro, RJ, Brazil b
a r t i c l e
i n f o
Article history: Received 3 July 2012 Received in revised form 6 November 2012 Accepted 6 November 2012 Available online 16 November 2012 Keywords: Mechanochemistry Multiferroic materials Yttrium manganite Microstructure Magnetic properties
a b s t r a c t Yttrium manganite (YMnO3) is a multiferroic material, which means that it exhibits both ferromagnetic and ferroelectric properties, so making it interesting for a variety of technological applications. In this work, single-phase YMnO3 was prepared for the first time by mechanochemical synthesis in a planetary ball mill. The YMnO3 was formed directly from the highly activated constituent oxides, Y2O3 and Mn2O3, after 60 min of milling time. During prolonged milling, the growth of the particles occurred. The cumulative energy introduced into the system during milling for 60 min was 86 kJ/g. The X-ray powder-diffraction analysis indicates that the as-prepared samples crystallize with an orthorhombic (Pnma) YMnO3 structure. The morphology and chemical composition of the powder were investigated by SEM and FESEM. The magnetic properties of the obtained YMnO3 powders were found to change as a function of the milling time in a manner consistent with the variation in the nanocomposite microstructure. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Multiferroic materials that exhibit any type of long-range magnetic ordering, spontaneous electric polarization, and/or ferroelasticity in one phase have been the subject of numerous scientific studies in recent years. Their interesting properties offer possibilities for many different technological applications in non-volatile ferroelectric memories, sensors, transducers and microwave devices [1]. As possible multiferroic materials, rare-earth manganites, RMnO3 (R being rare earth or transition metal ion), have attracted a lot of attention recently. These manganese (III) oxides have an interesting structure that depends on the size of the R cation. Those with larger R cations crystallize in the perovskite orthorhombic structure (space group Pnma), while those with a smaller R cation radii crystallize in the hexagonal structure (space group P63cm). Among these materials YMnO3 (YMO) stands out as a proven multiferroic material and because of its intermediate size the Y3+ ion allows the formation of both structures: the hexagonal (h-YMO) structure, stable under atmospheric conditions, and orthorhombic (o-YMO), the metastable one. The literature data shows that the Néel temperatures (TN) for h-YMO and o-YMO are 80 and 42 K,
⇑ Corresponding author. Tel.: +381 11 2085842. E-mail address: [email protected] (Z. Marinkovic´ Stanojevic´). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.11.031
respectively, while the ferroelectric Curie temperatures (TC) are 914 K (h-YMO) and 31 K (o-YMO) [2–5]. Being the stable one, h-YMO can be prepared by various techniques. Nanosized powders are obtained by a solid-state reaction from Mn(III) and Y(III) oxides [6,7], a hydrothermal process [8] or chemical methods employing different chelating agents, like citric acid or EDTA [9,10]. Polycrystalline and epitaxial YMO thin films with a hexagonal structure can be obtained by means of molecular-beam epitaxy, laser ablation, sputtering methods, pulsed-laser deposition or the sol–gel process [11,12]. Physical methods can also be used for the preparation of orthorhombic YMO thin films. But in this case, an appropriate choice of perovskite substrates and/or the use of a seeding layer, could be the crucial factor for the stabilization of the metastable phase [11,13–15]. On the other hand, the characterization of o-YMO powders is very difficult, since the preparation and stabilization of this phase are very complex. Orthorhombic single crystals of YMO can be prepared by the transformation of the hexagonal phase at high pressures and high temperatures [2,16]. Apart from this treatment, recommended synthetic routes for powder preparation are soft chemical methods, with a thermal treatment at lower temperatures [17,18]. These methods can be very complex, since the precursor structure has to fulfill certain requirements. The Mn3+ ion in the orthorhombic YMO is six-coordinated and these surroundings should be preserved throughout the whole synthesis and thermal treatment, without transforming to the five-coordinated Mn3+
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ion in the hexagonal phase. Even if this can be accomplished, the resulting powder can be a mixture of both YMO phases [19,20]. To the best of our knowledge, no literature data has been reported on the mechanochemical synthesis of YMO powders. This synthetic method is based on solid-state reactions taking part while the reactants are being milled in a high-energy ball mill. During these reactions the particle size of the reactants is decreasing, the contact area between the particles is increasing and the nucleation starts at lower temperatures, which results in the formation of products with a smaller particle size. Intensive milling also provides high temperatures and high pressures for the start of the reaction and for the formation of the product [21]. In this paper mechanochemical synthesis was proposed as a possible method for the preparation of orthorhombic YMnO3, having in mind that by using this method some perovskite manganitebased materials have already been prepared without post-milling high-pressure or high-temperature treatments [21]. The milling conditions were investigated, and the characterization of the asprepared powders was performed. The main result of our work is the preparation of a pure orthorhombic YMnO3 phase after only 240 min of milling. 2. Experimental procedure For the mechanochemical synthesis of YMnO3, Y2O3 (Alfa Aeser, p.a. 99.9%) and Mn2O3 (Aldrich, p.a. 99%) were used as the precursors. The appropriate amounts of oxides, with an equimolar Y3+ and Mn3+ ratio, according to Eq. (1), were weighed, mixed and then dry milled in a planetary ball mill (Fritsch Pulverisette 5). The milling conditions were as follows: tungsten carbide cylindrical vials (V = 250 cm3, Dv = 75 mm) and balls (35 balls with d = 10 mm, q = 14.95 g/cm3); a ball-to-powder weight ratio of 30:1; an air atmosphere; a basic disc rotational speed of 325 rpm; a rotation speed of the discs with vials equal to 400 rpm; and a milling time of 60– 720 min.
Y2 O3 þ Mn2 O3 ! 2YMnO3
ð1Þ
Based on the milling times the samples were marked as YMO60, YMO120, YMO240, YMO360 and YMO720. The ball-impact energy was calculated (DEb⁄ 60 mJ/hit) according to the model derived by Burgio et al. [22] The experimentally calculated weight-normalized cumulative energies (Ecum) required for starting (after 60 min) and the complete (after 240 min) formation of the YMnO3 phase were 86 kJ/g and 344 kJ/g, respectively. The X-ray diffraction (XRD) data were obtained using a RigakuÒ RINT2000 diffractometer (42 kV 120 mA) with CuKa radiation (kka1 = 1.5405 Å, kka2 = 1.5443 Å, Ika1/Ika2 = 2) in the 2h range between 10° and 80°, step 0.02° (2h) with 10.0 s per point. To obtain the microstructural data and the quantitative phase compositions of the as-milled powders a Rietveld refinement [23] was performed using the General Structure Analysis System (GSAS) program [24] suite with the EXPGUI interface [25]. The peak-profile function was modeled using a convolution of the Thompson–Cox–Hastings pseudo-Voigt function (pV-TCH) [26], employing the asymmetry function described by Finger et al. [27], which accounts for the
Fig. 1. (a) XRD patterns for Y2O3 and Mn2O3 mixture milled between 60 and 720 min and (b) observed, calculated and difference profiles for the sample YMO360.
asymmetry resulting from the axial divergence. The bi-dimensional model for crystallite size described by Larson & Von Dreele [24] was used to account for the anisotropy in the half width of the reflections. Differential scanning calorimetry (DSC) and thermogravimetric analyses (TGA) of the as-prepared powders were performed in the temperature interval from 20 to 1000 °C, with a heating rate of 10 °C/min (model SDT Q-600). The powders were
Table 1 Results of structural analysis (unit-cell parameters, cell volume, density and crystallite size (hti)), mass percentages and agreement indexes at the end of the refinements for the powders obtained after milling for 60–720 min with Rietveld refinement. YMO60 a (nm) b (nm) c (nm) V (nm3) q (g/cm3) hti \ (nm) hti// (nm) s = hti \ /hti// wt.% YMnO3 wt.% Y2O3 wt.% Mn2O3 Rwp(%) RF2(%) S
0.5744(2) 0.7433(2) 0.5263(1) 0.22478(6) 5.670
44.7(7) 9.3(4) 45.9(8) 3.14 0.78 1.41
Rwp is the weighted index, S the goodness of fit ð¼
YMO120 0.57493(8) 0.74216(9) 0.52592(5) 0.22441(6) 5.677
99.81(3) 0.18(6) – 3.36 0.87 1.49 pffiffiffiffiffi x2 Þ, RF is the structure factor index [23].
YMO240
YMO360
YMO720
0.57507(7) 0.74190(8) 0.52610(5) 0.22446(5) 5.677 40 21 1.90 100 – – 3.13 0.60 1.40
0.57584(7) 0.74151(8) 0.52655(5) 0.22484(5) 5.667 16 20 0.80 100 – – 3.53 0.78 1.58
0.57630(6) 0.74182(8) 0.52653(5) 0.22510(5) 5.662 27 62 0.44 100 – – 3.22 0.58 1.44
M. Pocˇucˇa-Nešic´ et al. / Journal of Alloys and Compounds 552 (2013) 451–456 also characterized by SEM (SEM TESCAN Vega TS 5130MM), FE-SEM (JEOL, Model JSM-6701F) and particle size analyses (HORIBA LA-920 laser scattering particle size distribution analyzer). The magnetic measurements of the YMO samples were carried out with a SQUID MPMS-XL-5 magnetometer from Quantum Design. The zero-field-cooled (ZFC) and field-cooled (FC) magnetization vs. temperature curves were studied in the temperature range 2–300 K, while isothermal magnetization measurements were recorded between 50 and 50 kOe at two temperatures, 5 and 300 K.
If we take into account only single-phase YMO samples, the lattice parameter a and the cell volume (V) of the mechano-synthesized YMnO3 perovskite phase are significantly higher in comparison to the reported ICSD values (166217). The Rietveld refinement indicates that prolonging the milling time from 240 to 720 min results in an increase of the cell parameter a and the
3. Results and discussion The XRD patterns of the powders obtained after milling for different times are shown in Fig. 1a. It is clear that already after 60 min of milling time, the distinctive peaks of the orthorhombic phase of YMO are present: (2h = 25.73, 33.41, 48.46, and 60.50). The YMO60 XRD pattern shows that, apart from the orthorhombic YMO phase, there is also some quantity of starting oxides and an amount of amorphous phase (detectable through the pronounced background halo around the 2h = 33.41 peak). The XRD pattern for the YMO120 shows almost no traces of the starting oxides and a significantly reduced quantity of amorphous material. The patterns for the powders obtained after milling for 240, 360 and 720 min are almost identical. They have peaks that can only be attributed to the orthorhombic YMO phase, while the amorphous halo disappeared. These conclusions coming from a qualitative data analysis were additionally verified using the Rietveld-refinement method. The corresponding results for the sample YMO360 are illustrated in Fig. 1b. Quantitative phase and microstructural analyses results and the Rietveld agreement indexes after the YMO refinements are presented in Table 1. The quantitative phase analysis for the YMO60 powder sample shows that the crystallization of the YMO started after milling times as short as 60 min, although there are also some quantities of starting Y2O3 (9.3 wt.%) and Mn2O3 (45.9 wt.%) present in the powder. The quantities of these oxides are insignificant in the YMO120 sample, i.e., only 0.18 wt.% of Y2O3 and no traces of Mn2O3. It is clear that the milling times of 240, 360 and 720 min resulted in the preparation of pure YMnO3 with the orthorhombic structure, o-YMO, space group Pnma. The starting crystal-structure parameters used in the Rietveld quantitative data analysis were in accordance with the structures presented in the 166217 (YMnO3), 86815 (Y2O3) and 9091 (Mn2O3) data cards, available from the Inorganic Crystal Structure Database (ICSD). The Rietveld-analysis-based results proved that the pure, orthorhombic, metastable phase of YMnO3 can be prepared through mechanochemical synthesis.
Fig. 2. DSC–TGA curve for YMO240 sample.
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Fig. 3. Results of particle size distribution for YMO powders.
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Fig. 4. SEM micrographs for powder samples: (a) YMO120, (b) YMO360, (c) and (d) YMO720.
cell volume. The variation in the reported lattice parameter might result from the presence of non-stoichiometry as a consequence of the formation of point and lattice defects during the mechanochemical synthesis. The values of the crystallite sizes are relatively small (<65 nm), while the time of milling promotes the presence of a crystallite-shape (T) anisotropy. This is in accordance with some previous studies [28,29]. Fig. 2. shows the results of the thermal analysis for sample YMO240. It is clear that the weight loss during the thermal treatment is insignificant and it may be attributed to the loss of the always-present hygroscopic water and carbonates, or some small quantities of impurities. This result is in accordance with the idea that after 240 min of milling, pure o-YMO is formed and since it does not have any volatile components, only structural changes are expected to occur during heating. However, the absence of any distinctive peaks up to 1000 °C indicates the complete crystallization of o-YMO, while the phase transformation to stable h-YMO is postponed to higher temperatures. This is in accordance with the literature data, which suggest that the pure hexagonal phase can be prepared at 1100 °C or higher temperatures [18,30]. The results of the particle size analysis for the as-prepared powders of pure o-YMO are given in Fig 3. They reveal a bimodal particle size distribution for all the samples, except YMO120 and YMO240. It can be seen that the smallest particle sizes are present in the powder milled for 120 min, which is the sample with the narrowest particle size distribution. A characteristic of all the powders is the prevalence of a median particle size in the range between 1 and 3 lm. By prolonging the milling time, the particle size increases and the particle size distribution is broadened, suggesting the possibility of further particle
agglomeration, which is a consequence of the long milling times during the mechanochemical synthesis. The results of the microstructural analysis of the as-prepared powders are in accordance with the particle size analysis. From Fig. 4a, which shows an SEM micrograph of the YMO120 sample, it is clear that the particles have a similar, uniformly distributed size around 1 lm. However, prolongation of the milling time resulted in the formation of powders with larger particle sizes and with large agglomerates, even larger than 10 lm. SEM micrographs for both single particles and agglomerates, for the sample YMO720, are presented in Fig. 4c and d. An additional microstructural analysis of the sample YMO360 (Fig. 5), using the FE-SEM technique, offered a new insight into the inner microstructure of the particle. From this micrograph (Fig. 5) we can see that the as-prepared particles are dense and pore-free. Also, it is clear that the primary particles (crystallites) that are composing the secondary particles previously identified with the SEM and the particle size analyses are much smaller than 100 nm and that the as-prepared o-YMO powders obtained from the mechanochemical synthesis could be considered as nanopowders, being in accordance with the XRD results (Table 1). Fig. 6 shows the magnetization of the single-phase YMO samples as a function of magnetic field, M(H), under ZFC at a temperature of 5 K and YMO240 samples at 300 K. The absence of any hysteresis in the magnetization curves at room temperature is evidence for paramagnetic behavior, which confirms that the ferromagnetic component disappears above TN. The M vs. H curve at low temperatures (5 K) deviates from the linear relation, but the magnetization does not show saturation in a field up to 50 kOe. The minor magnetic hysteresis at 5 K (see inset of Fig. 6) and the
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Table 2 Different parameters deduced from the magnetization curves of mechanochemically synthesized YMO powder samples. Sample
TN (K)
Hc 103 (Oe)
hCW (K)
leff (lB)
YMO240 YMO360 YMO720
42 37 37
0.742 0.503 0.501
44 42 37
4.57 4.45 4.41
TN, Neel temperature (antiferro-paramagnetic transition); Hc, coercive field at 5 K; hCW, the paramagnetic Curie–Weiss temperature; and leff, effective magnetic moment.
Fig. 5. FE-SEM micrograph of a secondary particle of the sample YMO360, indicating nanosized dimensions of the primary particles (crystallites).
very low coercive fields (see Table 2) indicate that the samples are basically anti-ferromagnetic with weak ferromagnetism. The thermal evolution of the dc magnetic susceptibility of the single-phase YMO samples (YMO240, YMO360 and YMO720) is exhibited in Fig. 7. It is clear that initially the magnetization increases gradually over a broad region as the temperature decreases. The diffuseness of the magnetic transition probably indicates that an inhomogeneous magnetic state prevails in the transition region. The broadening of the magnetic transition is consistent with the observed broadening of the X-ray reflections and with the model of short-range anti-ferromagnetism and a spinglass state [31,32]. A typical cusp exists in the ZFC curve and the low-temperature ZFC and FC curves deviate substantially for all the samples. The phenomenon of the discrepancy between the ZFC and the FC magnetization curves is usually ascribed to the appearance of the spin-glass or cluster-glass state induced by the competition between the ferromagnetic and anti-ferromagnetic exchange interactions. In addition, the M(H) curve displays a small hysteresis at 5 K (see inset in Fig. 6), consistent with a conventional spin-glass system [8,17,33]. Moreover, in Fig. 7 we can see that the onset of magnetization occurs at about 50–70 K, which is above the Néel temperature (40 K) reported for bulk orthorhombic YMnO3 [6,17,31]. This shift could also reflect the strain-induced unbalance of the anti-ferromagnetic competing interactions in the samples [14].
Fig. 6. Isothermal magnetization curves of YMO samples measured at 5 and 300 K. The inset shows the enlarged version of M (H).
Fig. 7. Magnetization as a function of temperature, ZFC and FC cycles, measured under an applied field of 100 Oe for YMO240, YMO360 and YMO720 samples. The solid symbols denote ZFC, whereas the open symbols denote FC magnetization. Inset shows the thermal evolution of the inverse susceptibility for all the YMO samples.
The inset of Fig. 7 exhibits the reciprocal susceptibility vs. temperature of mechanochemically synthesized YMO samples. By fitting with the Curie–Weiss law, v = C/(T hCW), in the temperature range above T = 120 K, the values of the effective paramagnetic moments (leff) and the paramagnetic Curie–Weiss temperatures (hCW) could be obtained. Table 2. collects these parameters together with the TN of the single-phase YMO samples. TN was defined by the peak in the ZFC d(vT)/dT vs. T curve (not presented here). The negative value of the paramagnetic Curie–Weiss temperature indicates the dominance of the anti-ferromagnetic interaction in all the samples. The sets of Curie–Weiss parameters are in good agreement with the previously reported values for the orthorhombic YMnO3 [6,20,31]. The effective magnetic moment obtained from the reciprocal susceptibility is below the expected value for the spin-only Mn3+ ions in the ground state (lMn3+ = 4.90lB) (see Table 2). This could be a result of the covalence effects and could also be understood by considering that a small content of Mn4+ (lMn4+ = 3.87lB) with a smaller expected effective magnetic moment is present in the samples YMO240, YMO360 and YMO720. Probably, the presence of some Mn4+ ions results from the slight oxygen excess localized in the interstitial sites. In the stoichiometric YMnO3, the magnetic interaction of the two neighboring Mn3+ is mediated by the oxygen ions between them and is ascribed to the super-exchange mechanism with the resulting anti-ferromagnetic coupling of the moments. The non-stoichiometry-related defects could lead to the formation of Mn4+ states, which introduce some disorder in the crystal lattice and the subsequent magnetic disorder. These Mn4+ ions do not participate in the super-exchange interaction but open the double-exchange interaction channel whenever they are next
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to a Mn3+ ion. In that case, double-exchange Mn3+/4+–Mn3+/4+ ferromagnetic interactions can be introduced, thus leading to competing ferromagnetic-anti-ferromagnetic interactions that may behave as a spin-glass-like material [17,34]. In addition, we can speculate that the slight difference in the crystal structure due to the microstrain effect (induced by nonstoichiometry-related defects) might have an influence on the decrease of the magnetic transition temperature through the changes in the bond lengths and the angles between the magnetic and oxygen atoms. Because of that, the small depression in transition temperature (TN) in YMO360 and YMO720 can be attributed to the non-stoichiometric composition of this sample. 4. Conclusions Pure orthorhombic YMO powders were successfully prepared by mechanochemical synthesis in a planetary ball mill, after 240 min of milling, starting from Y2O3 and Mn2O3 precursors. According to the XRD analysis and the Rietveld refinement, the YMnO3 obtained after milling for 240 min had a orthorhombic perovskite structure (Space group Pnma) and the following lattice parameters: a = 0.57507 nm, b = 0.74190 nm, c = 0.52610nm. The mean particle size for the YMO powders milled for 240 min is 1.25 lm. The SEM analysis proved the agglomeration of powders with prolonged milling times; however, the FE-SEM analysis evidence that the primary particles (crystallites) are smaller than 100 nm. The irreversibility in the magnetic response and the small hysteresis at low temperature indicate that the samples obtained by mechanochemical synthesis are basically anti-ferromagnetic, with weak ferromagnetism. We have proposed that this behavior could be associated with a modification of the magnetic interactions in the samples due to the presence of non-stoichiometry-related defects. Acknowledgement This work was supported by Ministry of Education and Science of Republic of Serbia through project III45007 and COST Action MP0904. B. A. Marinkovic´ is grateful to CNPq for a Research Productivity Grant. References [1] M. Fiebig, J. Phys. D: Appl. Phys. 38 (2005) R123–R152.
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