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Mechanochemical transformation of α-Fe2O3 to Fe3−xO4–microstructural investigation
Journal of Alloys and Compounds 348 (2003) 278–284
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Mechanochemical transformation of a-Fe 2 O 3 to Fe 32x O 4 –microstructural investigation a, b ,1 b,c d M. Hofmann *, S.J. Campbell , W.A. Kaczmarek , S. Welzel a
b
Rutherford Appleton Laboratory, ISIS, Chilton, Didcot OX11 0 QX, UK School of Physics, UNSW@ ADFA, The University of New South Wales, Australian Defence Force Academy, Canberra, ACT 2600, Australia c Research School of Physical Sciences and Engineering, Australian National University, Canberra, ACT 0200, Australia d Hahn-Meitner-Institut, BENSC, Glienickerstr. 100, D-14109 Berlin, Germany Received 6 May 2002; accepted 17 May 2002
Abstract The effects of wet-milling a-Fe 2 O 3 in vacuum for up to 144 h have been investigated by neutron diffraction measurements at room temperature and in situ at |950 K. Rietveld refinements show that the main product is iron-deficient magnetite of approximate stoichiometry |Fe 2.8 O 4 . Comparison of the phases derived from the neutron data with results of the Fe 21 / Fe 31 oxidation states as determined by chemical analysis reveals that a significant fraction of the unreacted a-Fe 2 O 3 occurs in an amorphous-like or disordered state. The wet-milled products are also found to contain |8% g-Fe 2 O 3 . The transformation from the a-phase to the g-phase occurs as a result of the shearing during the low-energy milling, with the further collisions and impacts leading to defect magnetite on extended milling. While other contributing effects take place, the transformation process from a-Fe 2 O 3 to Fe 32x O 4 occurs mainly as a result of rupturing the oxide surface layers of a-Fe 2 O 3 and releasing oxygen with consequent reduction to Fe 32x O 4 . 2002 Elsevier Science B.V. All rights reserved. Keywords: Oxide materials; Mechanical alloying; Crystal structure; Neutron diffraction
1. Introduction The 16 iron oxides (including hydroxides and oxide hydroxides) have numerous applications and are of continuing importance for research and technology [1]. Applications range from their use as pigments for the paint and construction industry, raw materials for the iron and steel industries and production of ferrites such as BaFe 12 O 19 for use as permanent magnets. More sophisticated applications, including their use as ferrofluids for hermetic seals, exploit the unique magnetic properties of iron oxides, particularly magnetite, Fe 3 O 4 , haematite, a-Fe 2 O 3 and maghemite, g-Fe 2 O 3 . As an example, g-Fe 2 O 3 is the most widely used magnetic recording material with Co surfacemodified g-Fe 2 O 3 particularly important because of its
* Corresponding author. Technische Universitat ¨ Munchen, ¨ ZWE, FRMII, D-85747 Garching, Germany. E-mail address: [email protected] (M. Hofmann). 1 Former address: Johannes Gutenberg-University, D-55099 Mainz, Germany.
increased coercivity [1,2]. Haematite, a-Fe 2 O 3 , is equally important—the structural and magnetic properties of crystalline a-Fe 2 O 3 are well established [3] with the properties of haematite-like materials also attracting interest recently (e.g. Ref. [4]). Ball milling and related mechanical treatments continue to be developed as important methods for the preparation of new materials with enhanced physical properties (e.g. Refs. [5,6]). The effects of milling iron oxides have attracted particular attention, with an overall aim being to explore the structural effects and related changes in magnetic properties introduced on such mechanochemical treatments (e.g. Refs. [7–10]). For example, the influence of particle morphology in the magnetic properties of haematite has recently been investigated [8] as has the influence of interparticle interactions on the magnetic behaviour of nanoparticles of haematite [9] along with the effects of milling magnetite and maghemite [10]. The effects of milling a-Fe 2 O 3 under a variety of environments and milling conditions have been investigated by a number of groups (e.g. Refs. [7,11–19]) with the transformation of a-Fe 2 O 3 to iron deficient magnetite,
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S0925-8388( 02 )00808-3
M. Hofmann et al. / Journal of Alloys and Compounds 348 (2003) 278–284
Fe 32x O 4 , being of special interest. Recent investigations of magnetite prepared by ball-milling a-Fe 2 O 3 with Fe reveal enhancement in the magnetic properties of the milled magnetite product compared with those of single domain magnetite, albeit at the cost of reduced phase stability [18]. Here we investigate the possible role of the unstable maghemite phase, g-Fe 2 O 3 , in the transformation of aFe 2 O 3 to Fe 32x O 4 . However, given the similarity in lattice parameters of g-Fe 2 O 3 and Fe 3 O 4 it is difficult to separate them in mixtures of milled, iron oxide phases (g-Fe 2 O 3 can be described by the spinel structure of lattice parame˚ slightly reduced compared with that of ter a 0 58.3457 A, ˚ Likewise the similarity in Fe 3 O 4 for which a 0 58.399 A). the hyperfine interaction parameters for these phases means that their unequivocal identification by 57 Fe ¨ Mossbauer spectroscopy remains uncertain although applied magnetic fields can be used to assist in separating the spectral components [20]. This problem of overlap of parameters is exacerbated in the case of milled samples with a spread of structural and hyperfine values generally being obtained. Our approach has been to investigate the milled aFe 2 O 3 products by in situ neutron diffraction measurements up to |950 K. The extent to which Fe 2 O 3 phases remain in a disordered / amorphous form after milling has been determined by Rietveld refinements of the re-crystallised phases, with good agreement being obtained with the results of Fe 21 cation content from chemical analysis. The present work extends the preliminary findings reported elsewhere [21].
2. Experimental Full details of the milling treatments have been outlined previously [13,16,21]. In brief, the a-Fe 2 O 3 powder of purity 99.99% (Koch-Light Laboratories) was milled in vacuum with |5 ml water in a vertical stainless steel mill and stainless steel balls at a pressure of |10 2 –10 3 Pa. The materials were wet-milled for 72 h and 144 h in order to match the conditions and period over which the transformation from haematite to magnetite had been reported previously [7]. Neutron diffraction measurements were carried out in both time-of-flight and constant wavelength modes using LAD at room temperature (ISIS, Rutherford Appleton Laboratory, UK) and the E6 neutron diffractome˚ Hahnter at high temperatures (|300–950 K; l52.445 A; Meitner-Institut, Berlin). The diffraction patterns were analysed using Rietveld methods with nuclear and magnetic contributions being refined simultaneously. The SEM analyses were carried out using a Zeiss LEO 430 scanning electron microscope with the particle size distributions estimated from measurements of |660 crystallites for each powder sample.
279
Fig. 1. Room temperature time-of-flight neutron diffraction patterns and Rietveld refinements for a-Fe 2 O 3 samples wet-milled for 72 and 144 h. The phase markers are: a-Fe (top), a-Fe 2 O 3 (middle) and Fe 32x O 4 (bottom).
3. Results and discussion The room temperature time-of-flight neutron diffraction patterns of the samples wet-milled for 72 and 144 h are shown in Fig. 1 with the results of Rietveld refinements given in Table 1. As discussed below, the two main crystalline phases observed are off-stoichiometric f.c.c. spinel-type Fe 32x O 4 and unreacted rhombohedral aFe 2 O 3 . The refinements also reveal the presence of |2–3% b.c.c. a-Fe (resulting from the stainless steel balls and mill) in the milled products. The optimal refinements of the Fe 32x O 4 phases were Table 1 Results of the Rietveld refinements of the room temperature neutron diffraction patterns of a-Fe 2 O 3 wet-milled for 72 and 144 h (cf. Fig. 1)
Time (h) Frac. (%) ˚ a (A) a (8) m ( mB ) m( tet.) m(oct.) x kDl (nm) Da /a (%)
Fe 32x O 4
a-Fe 2 O 3
Fe 32x O 4
a-Fe 2 O 3
72 90(1) 8.3876(1) – – 4.19(4) 3.88(5) 0.15(1) 42(5) 0.30(3)
72 7(1) 5.429(2) 55.27(2) 3.9(1) – – – – 0.21(3)
144 96(1) 8.3845(2) – – 4.12(6) 3.50(5) 0.21(1) 47(5) 0.35(3)
144 3(1) 5.430(2) 55.25(2) 3.4(3) – – – 125(10) 0.11(3)
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M. Hofmann et al. / Journal of Alloys and Compounds 348 (2003) 278–284
obtained on the basis that the vacancies occur in the energetically more favourable octahedral Fe 21 / Fe 31 sites rather than the tetrahedral Fe 31 sites [10,16], leading to the off-stoichiometric Fe 32x O 4 values given in Table 1 with x|0.15–0.2. While refinements of the data with scope for occupancies of both the octahedral and tetrahedral sites were found to improve the reliability factors slightly, non-physical Fe contents in the Fe 32x O 4 phases were obtained with vacancy values of x.0.33. By comparison, our previous investigations of a-Fe 2 O 3 milled in the equivalent manner to the present samples revealed optimal refinements to the neutron diffraction pattern with vacancies occurring on tetrahedral sites as well as the octahedral sites [16,17]. In addition, weak superlattice lines corresponding to the onset of ordering of vacancies in fully ordered g-Fe 2 O 3 were observed in the neutron diffraction pattern of a-Fe 2 O 3 wet-milled for the extended period of 200 h [16]. These effects demonstrate the extent to which a wide range of distributions of vacancies and correspondingly Fe cations are obtained in the defect Fe 32x O 4 product that results from wet-milling a-Fe 2 O 3 . Novikov et al. [19] have recently carried out a detailed investigation ¨ by X-ray diffraction and Mossbauer spectroscopy of aFe 2 O 3 milled in water for periods of up to 30 h. They concluded that the Fe 32x O 4 product is chemically heterogeneous with a distribution of iron cations in the crystal lattice. For example, analysis of the relative intensities of XRD peak reflections of the defect Fe 32x O 4 produced on wet-milling a-Fe 2 O 3 for 30 h indicated that vacancies primarily occupy the octahedral sites but with a probability which differs from one crystallographic plane to another, depending on the type and degree of occupation ¨ with ions. Similarly, deconvolution of the Mossbauer spectrum of Fe 2.847 O 4 based on the probability distribution of magnetic hyperfine fields were consistent with local environments of the Fe 21 / Fe 31 octahedral sites with vacancy values ranging from x 0 50.1, x 1 |0.12, x 2 |0.18 to x 3 |0.26 [19]. The optimal analyses for the present samples indicate an increase in the apparent phase fraction of Fe 32x O 4 from |90% to |96% with increased milling time from 72 to 144 h. However, while this behaviour shows general agreement with an earlier X-ray diffraction investigation of milled a-Fe 2 O 3 [7], it should be noted that the refinements apply to the crystalline phases present in the milled products and, as discussed below, do not include the disordered phase fractions remaining in the milled products. The occurrence of disordered components in the diffraction patterns leads to uncertainties in the background level which in turn restricts detailed peak shape analyses and determination of the peak widths. This means that only lower level values for the crystallite sizes and the isotropic microstrains (root mean square of lattice strain Da /a) of the milled phases can be estimated from the peak broadening (Table 1). Within this limitation, the mean crystallite sizes for
Fe 32x O 4 in both milled samples (kDl | 45 nm) are found to agree generally with the particle size distributions determined by direct measurements (Fig. 2), with the SEM analyses for both milled samples exhibiting broad particle size distributions centred around |50 nm. The relatively large fractions of particles of average sizes . |100 nm indicated by the SEM data (Fig. 2a) can be linked with the presence of unreacted a-Fe 2 O 3 in the milled products. Indeed, the fraction of unreacted a-Fe 2 O 3 in the 72 h milled sample remained sufficiently large that the particle size cannot be estimated from the peak broadening (Table 1). As is evident from earlier studies of milled iron oxides, depending on the mill material and sample environment, a range of products can be observed (e.g. Refs. [22,23]). Indeed, adding to these complexities are the possible effects of contamination from the milling environment [24] and reversible transformations taking place during milling [15]. In the case of the transformation of a-Fe 2 O 3 to Fe 3 O 4 , Kosmac and Courtney [24] noted that the presence of Fe contamination from the milling treatment could reduce a-Fe 2 O 3 to Fe 3 O 4 . However, in considering the possible processes responsible for this reduction process, Linderoth et al. [15] concluded that, at least in their milling experiments (which are likely to be typical of most studies), the level of Fe contamination was insufficient to account for the amount of Fe 3 O 4 resulting from the milling. In the present milling experiments, for which the refinements reveal the occurrence of about |2–3% unreacted Fe (Fig. 1), it is clear that the process, 4a-Fe 2 O 3 1 Fe → 3Fe 3 O 4 could contribute in part to the transformation of a-Fe 2 O 3 to defect Fe 32x O 4 as observed in the present experiments. The extent to which this process contributes to this transformation could be checked using a ceramic, nonmetallic mill. In order to clarify further the phases and oxidation states present in the milled products, the Fe 21 cation contents of the milled samples were determined by chemical analysis; this provides insight to the oxide phases for comparison with the neutron results. Each sample was analysed using an acidified permanganate titration method in which the content of Fe 21 ions was measured relative to the whole Fe content. The chemical analysis revealed |59% of the Fe 32x O 4 phase in the 72 h milled sample and |61% in the 144 h milled sample (for simplicity the milled product was assumed to comprise stoichiometric Fe 3 O 4 ). These values are significantly lower than the fractional Fe 32x O 4 phase values of |90% and |96% determined from the refinements of the room temperature neutron diffraction patterns of the 72 h and the 144 h milled samples, respectively (Table 1). However, as mentioned above, these latter values are derived from the crystalline-like peaks present
M. Hofmann et al. / Journal of Alloys and Compounds 348 (2003) 278–284
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Fig. 2. (a) Particle size distributions of a-Fe 2 O 3 samples wet-milled for 72 and 144 h. (b) A typical SEM image as obtained for the sample milled for 144 h. The marker indicates a length scale of 200 nm.
in the diffraction patterns of the milled products rather than the pattern as a whole (Fig. 1). These differences could be linked with either the effects of an amorphous-like Fe 2 O 3 phase contributing to the diffuse background scattering or a residual, nanoscale a-Fe 2 O 3 haematite phase that remains undetected in the diffraction patterns due to peak broadening. Also, as noted above, given the similarity in the crystal structural parameters between the Fe 3 O 4 and
g-Fe 2 O 3 cells, a further factor is the possible presence of the g-Fe 2 O 3 maghemite phase. In order to clarify these points, additional in situ high temperature measurements were carried out at |950 K on the as-milled products. The milled samples were heated in high vacuum in order to prevent Fe 3 O 4 from being oxidised to Fe 2 O 3 . The temperature of |950 K was selected as it is well above both the transformation temperature from g-Fe 2 O 3 to a-Fe 2 O 3
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and the magnetic transition temperatures of the iron oxide phases. Fig. 3 shows a series of constant wavelength diffraction patterns for the 72 h sample—in the as-milled state at room temperature (Fig. 3a); at |950 K (Fig. 3b) and again at room temperature after cooling from |950 K (Fig. 3c). The equivalent diffraction patterns for the 144 h milled sample at room temperature and at |950 K are shown in Fig. 3d and e, respectively. Consistent with re-crystallisation of any amorphous / disordered components and grain growth as indicated by the narrow diffraction peaks, the background scattering is found to be reduced significantly for both samples at |950 K (Fig. 3b and d). The refinements of these diffraction patterns show that the relative fraction of a-Fe 2 O 3 has increased from its apparent value of |7% for the 72 h milled sample to |37% in its recrystallised state (Table 2). Similarly the a-Fe 2 O 3 fraction was found to increase from the apparent value of |4% for the 144 h milled sample to |29% in its
Table 2 Results of the Rietveld refinements of the high temperature neutron diffraction patterns of a-Fe 2 O 3 wet-milled for 72 h and 144 h (cf. Fig. 3)
Time (h) Frac. (%) ˚ a (A) a (8)
Fe 32x O 4
a-Fe 2 O 3
Fe 32x O 4
a-Fe 2 O 3
72 63(2) 8.4758(6) –
72 37(1) 5.462(8) 55.39(5)
144 71(4) 8.4725(6) –
144 29(2) 5.459(8) 55.42(5)
recrystallised state. These phase fractions for a-Fe 2 O 3 of |37% and |29% are in general agreement with the results of the chemical analyses (|41% and |39% in the Fe 2 O 3 form for the 72 h and 144 h milled samples, respectively). The agreement between these two different analytical methods is consistent with the conclusion that the majority of the Fe 2 O 3 phase (a or g) in the as-milled samples exists in an amorphous or nanostructured / disordered form, thereby contributing to the diffuse neutron scattering observed in the as-milled samples (Fig. 3a and d).
Fig. 3. Constant wavelength neutron diffraction patterns and Rietveld refinements for a-Fe 2 O 3 : wet-milled for 72 h (a, b, c) and 144 h (d, e): (a) room temperature (72 h sample); (b) in situ at |950 K (72 h sample); (c) at room temperature after exposure to |950 K (72 h sample); (d) room temperature (144 h sample); (e) in situ at |950 K (144 h sample). The phase markers are: Fe 32x O 4 (top) and a-Fe 2 O 3 (bottom).
M. Hofmann et al. / Journal of Alloys and Compounds 348 (2003) 278–284
The basic structure of g-Fe 2 O 3 corresponds to space group P4 3 32 [25] with vacancies on the octahedral sites. This results in additional reflections in the diffraction patterns that violate the reflection conditions for spineltype structures. Further it was also suggested that g-Fe 2 O 3 has a unique ordered vacancy distribution in the octahedral sites [26,27]. This vacancy ordering reduces the symmetry of the structure to tetragonal (space group P4 3 2 1 2), leading to additional satellite peaks. While no satellite peaks were observed in the present milled samples, indirect evidence for the presence of crystalline g-Fe 2 O 3 in the milled products has been obtained by comparing the integrated intensities of peaks that, because of the similarity in lattice parameters, may have contributions from both Fe 3 O 4 and g-Fe 2 O 3 . For example, comparison of the (220) peak intensities at 2u |48.98 for the 72 h milled sample before (Fig. 3a) and after exposure to |950 K (Fig. 3b) reveals a decrease in intensity of |8%. This decrease is likely to be due to the reversion of g-Fe 2 O 3 to the stable a-Fe 2 O 3 phase above the transition temperature of |300–400 8C. This result indicates that for the 72 h as-milled sample, g-Fe 2 O 3 would contribute a maximum of |20% (i.e. |8 / 41) of the Fe 2 O 3 (a or g) fraction. These findings demonstrate that a significant fraction of the Fe 2 O 3 phase in the milled samples exists in an amorphous / disordered phase; this behaviour is likely to be common to all milled iron oxides. The occurrence of g-Fe 2 O 3 as part of the transformation from a-Fe 2 O 3 to Fe 32x O 4 is consistent with the identification of |10–15% g-Fe 2 O 3 in a series of wet-milled a¨ Fe 2 O 3 samples by integral conversion electron Mossbauer spectroscopy measurements which are sensitive to the top layer (|10–100 nm) of materials [28]. Identification of the g-Fe 2 O 3 on the surfaces of particles milled for more than 24 h in this scattering experiment [28] agrees with the conclusions of Petrovsky et al. [18]. They reported that the most probable outcome of their milling of a-Fe 2 O 3 and Fe in an inert gas environment is a spinel magnetite core with a highly distorted surface. As is well known (e.g. Refs. [10,29]), defect, non-stoichiometric Fe 32x O 4 can be repre31 sented by [Fe 31 ] tetra [Fe 21 (123x) Fe ( 112x) h x ] octa O 4 where h x is the cationic vacancy. In the limit of non-stoichiometry (x51 / 3), g-Fe 2 O 3 is represented as [Fe 31 ] tetra [Fe 531/ 3 h 1 / 3 ] octa O 4 . Depending on the degree of distortion resulting from the range of defects introduced by the present wet-milling treatments in vacuum, the main products are therefore expected to include defect, nonstoichiometric Fe 32x O 4 and g-Fe 2 O 3 as observed in our neutron diffraction measurements. The present results are also consistent with the findings of Meillon et al. [14] who report a direct phase transformation of a-Fe 2 O 3 to g-Fe 2 O 3 on wet-grinding in ethanol. The mechanism for this latter transformation is based on the movement of the oxygen planes in a-Fe 2 O 3 due to the periodic shearing sequence used in their grinding experi-
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ments. In their scheme, the transformation is obtained by a shear sequence of planes followed by a rearrangement of the cations until domains of the g-Fe 2 O 3 phase occur with diminishing grain size. The g-Fe 2 O 3 product in their grinding experiment has a primitive cubic unit cell (space group P4 3 32 [25]) as evidenced by the 110, 210 and 112 reflections in their X-ray diffraction pattern (Fig. 1 of Ref. [14]). Given the low-energy mode used in the present mill [30] and experiments, the a-Fe 2 O 3 powder was worked predominantly by shearing with impact and collisions (due to powder trapped between balls and / or balls and vessel walls) also occurring. These mechanisms are therefore considered to contribute to shearing the a-Fe 2 O 3 powder with consequent transformation to g-Fe 2 O 3 and, with extended milling and continued exposure to impact as well as shearing, to the rupturing of the oxide surface layers and release of oxygen with consequent reduction to Fe 32x O 4 [7,13]. As noted above, these physical mechanisms would also account for the significant defect fractions measured for the off-stoichiometric Fe 32x O 4 with vacancy values of x|0.2 (Table 1).
4. Conclusions The effects of milling iron oxides are complex. Depending on the mill materials and sample environment a variety of phases—crystalline, nanostructured, as well as amorphous and disordered states can be obtained [5,22]. In the present investigation of wet-milling a-Fe 2 O 3 in vacuum for up to 144 h, neutron diffraction measurements have confirmed that the main crystalline products are nonstoichiometric Fe 32x O 4 (with octahedral Fe 21 / Fe 31 vacancy concentrations of x|0.2) and the unreacted aFe 2 O 3 . Comparisons of neutron diffraction patterns of the as-milled samples at room temperature and |950 K combined with chemical analyses indicate that a significant fraction of the milled products occurs in a disordered or amorphous form and that a maximum fraction of |8% g-Fe 2 O 3 is also present. Given that the predominant work on the a-Fe 2 O 3 powder for the low-energy mode used in the present mill takes place by shearing, this occurrence of g-Fe 2 O 3 is likely to be due to a shearing transformation involving movement of the oxygen planes [14]. Extended milling and continued exposure to impact as well as shearing is considered to lead to the rupturing of the oxide surface layers and release of oxygen with consequent reduction to Fe 32x O 4 [7]. In conclusion, while the occurrence of |2–3% Fe as a contaminant in the present low-energy milling experiments is likely to contribute via 4a-Fe 2 O 3 1Fe→3Fe 3 O 4 , the main transformation of a-Fe 2 O 3 to defect Fe 32x O 4 takes place by rupturing of the oxide surface layers of a-Fe 2 O 3 . This leads to release of oxygen and reduction to defect Fe 32x O 4 .
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Acknowledgements S.J.C. acknowledges renewal of an Alexander von Humboldt Research Fellowship while at the Johannes Gutenberg-University, Mainz. We acknowledge access to the LAD diffractometer at the ISIS Facility, Rutherford Appleton Laboratory, UK, and thank Dr. W.S. Howells for his assistance. We also thank Dr. A.V.J. Edge for assistance with the wet chemical analysis. This work was also supported in part by grants from the Australian Institute of Nuclear Science and Engineering and S.J.C. acknowledges support from the Access to Major Research Facilities Program, Australian Nuclear Science and Technology Organisation.
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Mechanochemical transformation of a-Fe 2 O 3 to Fe 32x O 4 –microstructural investigation a, b ,1 b,c d M. Hofmann *, S.J. Campbell , W.A. Kaczmarek , S. Welzel a
b
Rutherford Appleton Laboratory, ISIS, Chilton, Didcot OX11 0 QX, UK School of Physics, UNSW@ ADFA, The University of New South Wales, Australian Defence Force Academy, Canberra, ACT 2600, Australia c Research School of Physical Sciences and Engineering, Australian National University, Canberra, ACT 0200, Australia d Hahn-Meitner-Institut, BENSC, Glienickerstr. 100, D-14109 Berlin, Germany Received 6 May 2002; accepted 17 May 2002
Abstract The effects of wet-milling a-Fe 2 O 3 in vacuum for up to 144 h have been investigated by neutron diffraction measurements at room temperature and in situ at |950 K. Rietveld refinements show that the main product is iron-deficient magnetite of approximate stoichiometry |Fe 2.8 O 4 . Comparison of the phases derived from the neutron data with results of the Fe 21 / Fe 31 oxidation states as determined by chemical analysis reveals that a significant fraction of the unreacted a-Fe 2 O 3 occurs in an amorphous-like or disordered state. The wet-milled products are also found to contain |8% g-Fe 2 O 3 . The transformation from the a-phase to the g-phase occurs as a result of the shearing during the low-energy milling, with the further collisions and impacts leading to defect magnetite on extended milling. While other contributing effects take place, the transformation process from a-Fe 2 O 3 to Fe 32x O 4 occurs mainly as a result of rupturing the oxide surface layers of a-Fe 2 O 3 and releasing oxygen with consequent reduction to Fe 32x O 4 . 2002 Elsevier Science B.V. All rights reserved. Keywords: Oxide materials; Mechanical alloying; Crystal structure; Neutron diffraction
1. Introduction The 16 iron oxides (including hydroxides and oxide hydroxides) have numerous applications and are of continuing importance for research and technology [1]. Applications range from their use as pigments for the paint and construction industry, raw materials for the iron and steel industries and production of ferrites such as BaFe 12 O 19 for use as permanent magnets. More sophisticated applications, including their use as ferrofluids for hermetic seals, exploit the unique magnetic properties of iron oxides, particularly magnetite, Fe 3 O 4 , haematite, a-Fe 2 O 3 and maghemite, g-Fe 2 O 3 . As an example, g-Fe 2 O 3 is the most widely used magnetic recording material with Co surfacemodified g-Fe 2 O 3 particularly important because of its
* Corresponding author. Technische Universitat ¨ Munchen, ¨ ZWE, FRMII, D-85747 Garching, Germany. E-mail address: [email protected] (M. Hofmann). 1 Former address: Johannes Gutenberg-University, D-55099 Mainz, Germany.
increased coercivity [1,2]. Haematite, a-Fe 2 O 3 , is equally important—the structural and magnetic properties of crystalline a-Fe 2 O 3 are well established [3] with the properties of haematite-like materials also attracting interest recently (e.g. Ref. [4]). Ball milling and related mechanical treatments continue to be developed as important methods for the preparation of new materials with enhanced physical properties (e.g. Refs. [5,6]). The effects of milling iron oxides have attracted particular attention, with an overall aim being to explore the structural effects and related changes in magnetic properties introduced on such mechanochemical treatments (e.g. Refs. [7–10]). For example, the influence of particle morphology in the magnetic properties of haematite has recently been investigated [8] as has the influence of interparticle interactions on the magnetic behaviour of nanoparticles of haematite [9] along with the effects of milling magnetite and maghemite [10]. The effects of milling a-Fe 2 O 3 under a variety of environments and milling conditions have been investigated by a number of groups (e.g. Refs. [7,11–19]) with the transformation of a-Fe 2 O 3 to iron deficient magnetite,
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Fe 32x O 4 , being of special interest. Recent investigations of magnetite prepared by ball-milling a-Fe 2 O 3 with Fe reveal enhancement in the magnetic properties of the milled magnetite product compared with those of single domain magnetite, albeit at the cost of reduced phase stability [18]. Here we investigate the possible role of the unstable maghemite phase, g-Fe 2 O 3 , in the transformation of aFe 2 O 3 to Fe 32x O 4 . However, given the similarity in lattice parameters of g-Fe 2 O 3 and Fe 3 O 4 it is difficult to separate them in mixtures of milled, iron oxide phases (g-Fe 2 O 3 can be described by the spinel structure of lattice parame˚ slightly reduced compared with that of ter a 0 58.3457 A, ˚ Likewise the similarity in Fe 3 O 4 for which a 0 58.399 A). the hyperfine interaction parameters for these phases means that their unequivocal identification by 57 Fe ¨ Mossbauer spectroscopy remains uncertain although applied magnetic fields can be used to assist in separating the spectral components [20]. This problem of overlap of parameters is exacerbated in the case of milled samples with a spread of structural and hyperfine values generally being obtained. Our approach has been to investigate the milled aFe 2 O 3 products by in situ neutron diffraction measurements up to |950 K. The extent to which Fe 2 O 3 phases remain in a disordered / amorphous form after milling has been determined by Rietveld refinements of the re-crystallised phases, with good agreement being obtained with the results of Fe 21 cation content from chemical analysis. The present work extends the preliminary findings reported elsewhere [21].
2. Experimental Full details of the milling treatments have been outlined previously [13,16,21]. In brief, the a-Fe 2 O 3 powder of purity 99.99% (Koch-Light Laboratories) was milled in vacuum with |5 ml water in a vertical stainless steel mill and stainless steel balls at a pressure of |10 2 –10 3 Pa. The materials were wet-milled for 72 h and 144 h in order to match the conditions and period over which the transformation from haematite to magnetite had been reported previously [7]. Neutron diffraction measurements were carried out in both time-of-flight and constant wavelength modes using LAD at room temperature (ISIS, Rutherford Appleton Laboratory, UK) and the E6 neutron diffractome˚ Hahnter at high temperatures (|300–950 K; l52.445 A; Meitner-Institut, Berlin). The diffraction patterns were analysed using Rietveld methods with nuclear and magnetic contributions being refined simultaneously. The SEM analyses were carried out using a Zeiss LEO 430 scanning electron microscope with the particle size distributions estimated from measurements of |660 crystallites for each powder sample.
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Fig. 1. Room temperature time-of-flight neutron diffraction patterns and Rietveld refinements for a-Fe 2 O 3 samples wet-milled for 72 and 144 h. The phase markers are: a-Fe (top), a-Fe 2 O 3 (middle) and Fe 32x O 4 (bottom).
3. Results and discussion The room temperature time-of-flight neutron diffraction patterns of the samples wet-milled for 72 and 144 h are shown in Fig. 1 with the results of Rietveld refinements given in Table 1. As discussed below, the two main crystalline phases observed are off-stoichiometric f.c.c. spinel-type Fe 32x O 4 and unreacted rhombohedral aFe 2 O 3 . The refinements also reveal the presence of |2–3% b.c.c. a-Fe (resulting from the stainless steel balls and mill) in the milled products. The optimal refinements of the Fe 32x O 4 phases were Table 1 Results of the Rietveld refinements of the room temperature neutron diffraction patterns of a-Fe 2 O 3 wet-milled for 72 and 144 h (cf. Fig. 1)
Time (h) Frac. (%) ˚ a (A) a (8) m ( mB ) m( tet.) m(oct.) x kDl (nm) Da /a (%)
Fe 32x O 4
a-Fe 2 O 3
Fe 32x O 4
a-Fe 2 O 3
72 90(1) 8.3876(1) – – 4.19(4) 3.88(5) 0.15(1) 42(5) 0.30(3)
72 7(1) 5.429(2) 55.27(2) 3.9(1) – – – – 0.21(3)
144 96(1) 8.3845(2) – – 4.12(6) 3.50(5) 0.21(1) 47(5) 0.35(3)
144 3(1) 5.430(2) 55.25(2) 3.4(3) – – – 125(10) 0.11(3)
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obtained on the basis that the vacancies occur in the energetically more favourable octahedral Fe 21 / Fe 31 sites rather than the tetrahedral Fe 31 sites [10,16], leading to the off-stoichiometric Fe 32x O 4 values given in Table 1 with x|0.15–0.2. While refinements of the data with scope for occupancies of both the octahedral and tetrahedral sites were found to improve the reliability factors slightly, non-physical Fe contents in the Fe 32x O 4 phases were obtained with vacancy values of x.0.33. By comparison, our previous investigations of a-Fe 2 O 3 milled in the equivalent manner to the present samples revealed optimal refinements to the neutron diffraction pattern with vacancies occurring on tetrahedral sites as well as the octahedral sites [16,17]. In addition, weak superlattice lines corresponding to the onset of ordering of vacancies in fully ordered g-Fe 2 O 3 were observed in the neutron diffraction pattern of a-Fe 2 O 3 wet-milled for the extended period of 200 h [16]. These effects demonstrate the extent to which a wide range of distributions of vacancies and correspondingly Fe cations are obtained in the defect Fe 32x O 4 product that results from wet-milling a-Fe 2 O 3 . Novikov et al. [19] have recently carried out a detailed investigation ¨ by X-ray diffraction and Mossbauer spectroscopy of aFe 2 O 3 milled in water for periods of up to 30 h. They concluded that the Fe 32x O 4 product is chemically heterogeneous with a distribution of iron cations in the crystal lattice. For example, analysis of the relative intensities of XRD peak reflections of the defect Fe 32x O 4 produced on wet-milling a-Fe 2 O 3 for 30 h indicated that vacancies primarily occupy the octahedral sites but with a probability which differs from one crystallographic plane to another, depending on the type and degree of occupation ¨ with ions. Similarly, deconvolution of the Mossbauer spectrum of Fe 2.847 O 4 based on the probability distribution of magnetic hyperfine fields were consistent with local environments of the Fe 21 / Fe 31 octahedral sites with vacancy values ranging from x 0 50.1, x 1 |0.12, x 2 |0.18 to x 3 |0.26 [19]. The optimal analyses for the present samples indicate an increase in the apparent phase fraction of Fe 32x O 4 from |90% to |96% with increased milling time from 72 to 144 h. However, while this behaviour shows general agreement with an earlier X-ray diffraction investigation of milled a-Fe 2 O 3 [7], it should be noted that the refinements apply to the crystalline phases present in the milled products and, as discussed below, do not include the disordered phase fractions remaining in the milled products. The occurrence of disordered components in the diffraction patterns leads to uncertainties in the background level which in turn restricts detailed peak shape analyses and determination of the peak widths. This means that only lower level values for the crystallite sizes and the isotropic microstrains (root mean square of lattice strain Da /a) of the milled phases can be estimated from the peak broadening (Table 1). Within this limitation, the mean crystallite sizes for
Fe 32x O 4 in both milled samples (kDl | 45 nm) are found to agree generally with the particle size distributions determined by direct measurements (Fig. 2), with the SEM analyses for both milled samples exhibiting broad particle size distributions centred around |50 nm. The relatively large fractions of particles of average sizes . |100 nm indicated by the SEM data (Fig. 2a) can be linked with the presence of unreacted a-Fe 2 O 3 in the milled products. Indeed, the fraction of unreacted a-Fe 2 O 3 in the 72 h milled sample remained sufficiently large that the particle size cannot be estimated from the peak broadening (Table 1). As is evident from earlier studies of milled iron oxides, depending on the mill material and sample environment, a range of products can be observed (e.g. Refs. [22,23]). Indeed, adding to these complexities are the possible effects of contamination from the milling environment [24] and reversible transformations taking place during milling [15]. In the case of the transformation of a-Fe 2 O 3 to Fe 3 O 4 , Kosmac and Courtney [24] noted that the presence of Fe contamination from the milling treatment could reduce a-Fe 2 O 3 to Fe 3 O 4 . However, in considering the possible processes responsible for this reduction process, Linderoth et al. [15] concluded that, at least in their milling experiments (which are likely to be typical of most studies), the level of Fe contamination was insufficient to account for the amount of Fe 3 O 4 resulting from the milling. In the present milling experiments, for which the refinements reveal the occurrence of about |2–3% unreacted Fe (Fig. 1), it is clear that the process, 4a-Fe 2 O 3 1 Fe → 3Fe 3 O 4 could contribute in part to the transformation of a-Fe 2 O 3 to defect Fe 32x O 4 as observed in the present experiments. The extent to which this process contributes to this transformation could be checked using a ceramic, nonmetallic mill. In order to clarify further the phases and oxidation states present in the milled products, the Fe 21 cation contents of the milled samples were determined by chemical analysis; this provides insight to the oxide phases for comparison with the neutron results. Each sample was analysed using an acidified permanganate titration method in which the content of Fe 21 ions was measured relative to the whole Fe content. The chemical analysis revealed |59% of the Fe 32x O 4 phase in the 72 h milled sample and |61% in the 144 h milled sample (for simplicity the milled product was assumed to comprise stoichiometric Fe 3 O 4 ). These values are significantly lower than the fractional Fe 32x O 4 phase values of |90% and |96% determined from the refinements of the room temperature neutron diffraction patterns of the 72 h and the 144 h milled samples, respectively (Table 1). However, as mentioned above, these latter values are derived from the crystalline-like peaks present
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Fig. 2. (a) Particle size distributions of a-Fe 2 O 3 samples wet-milled for 72 and 144 h. (b) A typical SEM image as obtained for the sample milled for 144 h. The marker indicates a length scale of 200 nm.
in the diffraction patterns of the milled products rather than the pattern as a whole (Fig. 1). These differences could be linked with either the effects of an amorphous-like Fe 2 O 3 phase contributing to the diffuse background scattering or a residual, nanoscale a-Fe 2 O 3 haematite phase that remains undetected in the diffraction patterns due to peak broadening. Also, as noted above, given the similarity in the crystal structural parameters between the Fe 3 O 4 and
g-Fe 2 O 3 cells, a further factor is the possible presence of the g-Fe 2 O 3 maghemite phase. In order to clarify these points, additional in situ high temperature measurements were carried out at |950 K on the as-milled products. The milled samples were heated in high vacuum in order to prevent Fe 3 O 4 from being oxidised to Fe 2 O 3 . The temperature of |950 K was selected as it is well above both the transformation temperature from g-Fe 2 O 3 to a-Fe 2 O 3
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and the magnetic transition temperatures of the iron oxide phases. Fig. 3 shows a series of constant wavelength diffraction patterns for the 72 h sample—in the as-milled state at room temperature (Fig. 3a); at |950 K (Fig. 3b) and again at room temperature after cooling from |950 K (Fig. 3c). The equivalent diffraction patterns for the 144 h milled sample at room temperature and at |950 K are shown in Fig. 3d and e, respectively. Consistent with re-crystallisation of any amorphous / disordered components and grain growth as indicated by the narrow diffraction peaks, the background scattering is found to be reduced significantly for both samples at |950 K (Fig. 3b and d). The refinements of these diffraction patterns show that the relative fraction of a-Fe 2 O 3 has increased from its apparent value of |7% for the 72 h milled sample to |37% in its recrystallised state (Table 2). Similarly the a-Fe 2 O 3 fraction was found to increase from the apparent value of |4% for the 144 h milled sample to |29% in its
Table 2 Results of the Rietveld refinements of the high temperature neutron diffraction patterns of a-Fe 2 O 3 wet-milled for 72 h and 144 h (cf. Fig. 3)
Time (h) Frac. (%) ˚ a (A) a (8)
Fe 32x O 4
a-Fe 2 O 3
Fe 32x O 4
a-Fe 2 O 3
72 63(2) 8.4758(6) –
72 37(1) 5.462(8) 55.39(5)
144 71(4) 8.4725(6) –
144 29(2) 5.459(8) 55.42(5)
recrystallised state. These phase fractions for a-Fe 2 O 3 of |37% and |29% are in general agreement with the results of the chemical analyses (|41% and |39% in the Fe 2 O 3 form for the 72 h and 144 h milled samples, respectively). The agreement between these two different analytical methods is consistent with the conclusion that the majority of the Fe 2 O 3 phase (a or g) in the as-milled samples exists in an amorphous or nanostructured / disordered form, thereby contributing to the diffuse neutron scattering observed in the as-milled samples (Fig. 3a and d).
Fig. 3. Constant wavelength neutron diffraction patterns and Rietveld refinements for a-Fe 2 O 3 : wet-milled for 72 h (a, b, c) and 144 h (d, e): (a) room temperature (72 h sample); (b) in situ at |950 K (72 h sample); (c) at room temperature after exposure to |950 K (72 h sample); (d) room temperature (144 h sample); (e) in situ at |950 K (144 h sample). The phase markers are: Fe 32x O 4 (top) and a-Fe 2 O 3 (bottom).
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The basic structure of g-Fe 2 O 3 corresponds to space group P4 3 32 [25] with vacancies on the octahedral sites. This results in additional reflections in the diffraction patterns that violate the reflection conditions for spineltype structures. Further it was also suggested that g-Fe 2 O 3 has a unique ordered vacancy distribution in the octahedral sites [26,27]. This vacancy ordering reduces the symmetry of the structure to tetragonal (space group P4 3 2 1 2), leading to additional satellite peaks. While no satellite peaks were observed in the present milled samples, indirect evidence for the presence of crystalline g-Fe 2 O 3 in the milled products has been obtained by comparing the integrated intensities of peaks that, because of the similarity in lattice parameters, may have contributions from both Fe 3 O 4 and g-Fe 2 O 3 . For example, comparison of the (220) peak intensities at 2u |48.98 for the 72 h milled sample before (Fig. 3a) and after exposure to |950 K (Fig. 3b) reveals a decrease in intensity of |8%. This decrease is likely to be due to the reversion of g-Fe 2 O 3 to the stable a-Fe 2 O 3 phase above the transition temperature of |300–400 8C. This result indicates that for the 72 h as-milled sample, g-Fe 2 O 3 would contribute a maximum of |20% (i.e. |8 / 41) of the Fe 2 O 3 (a or g) fraction. These findings demonstrate that a significant fraction of the Fe 2 O 3 phase in the milled samples exists in an amorphous / disordered phase; this behaviour is likely to be common to all milled iron oxides. The occurrence of g-Fe 2 O 3 as part of the transformation from a-Fe 2 O 3 to Fe 32x O 4 is consistent with the identification of |10–15% g-Fe 2 O 3 in a series of wet-milled a¨ Fe 2 O 3 samples by integral conversion electron Mossbauer spectroscopy measurements which are sensitive to the top layer (|10–100 nm) of materials [28]. Identification of the g-Fe 2 O 3 on the surfaces of particles milled for more than 24 h in this scattering experiment [28] agrees with the conclusions of Petrovsky et al. [18]. They reported that the most probable outcome of their milling of a-Fe 2 O 3 and Fe in an inert gas environment is a spinel magnetite core with a highly distorted surface. As is well known (e.g. Refs. [10,29]), defect, non-stoichiometric Fe 32x O 4 can be repre31 sented by [Fe 31 ] tetra [Fe 21 (123x) Fe ( 112x) h x ] octa O 4 where h x is the cationic vacancy. In the limit of non-stoichiometry (x51 / 3), g-Fe 2 O 3 is represented as [Fe 31 ] tetra [Fe 531/ 3 h 1 / 3 ] octa O 4 . Depending on the degree of distortion resulting from the range of defects introduced by the present wet-milling treatments in vacuum, the main products are therefore expected to include defect, nonstoichiometric Fe 32x O 4 and g-Fe 2 O 3 as observed in our neutron diffraction measurements. The present results are also consistent with the findings of Meillon et al. [14] who report a direct phase transformation of a-Fe 2 O 3 to g-Fe 2 O 3 on wet-grinding in ethanol. The mechanism for this latter transformation is based on the movement of the oxygen planes in a-Fe 2 O 3 due to the periodic shearing sequence used in their grinding experi-
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ments. In their scheme, the transformation is obtained by a shear sequence of planes followed by a rearrangement of the cations until domains of the g-Fe 2 O 3 phase occur with diminishing grain size. The g-Fe 2 O 3 product in their grinding experiment has a primitive cubic unit cell (space group P4 3 32 [25]) as evidenced by the 110, 210 and 112 reflections in their X-ray diffraction pattern (Fig. 1 of Ref. [14]). Given the low-energy mode used in the present mill [30] and experiments, the a-Fe 2 O 3 powder was worked predominantly by shearing with impact and collisions (due to powder trapped between balls and / or balls and vessel walls) also occurring. These mechanisms are therefore considered to contribute to shearing the a-Fe 2 O 3 powder with consequent transformation to g-Fe 2 O 3 and, with extended milling and continued exposure to impact as well as shearing, to the rupturing of the oxide surface layers and release of oxygen with consequent reduction to Fe 32x O 4 [7,13]. As noted above, these physical mechanisms would also account for the significant defect fractions measured for the off-stoichiometric Fe 32x O 4 with vacancy values of x|0.2 (Table 1).
4. Conclusions The effects of milling iron oxides are complex. Depending on the mill materials and sample environment a variety of phases—crystalline, nanostructured, as well as amorphous and disordered states can be obtained [5,22]. In the present investigation of wet-milling a-Fe 2 O 3 in vacuum for up to 144 h, neutron diffraction measurements have confirmed that the main crystalline products are nonstoichiometric Fe 32x O 4 (with octahedral Fe 21 / Fe 31 vacancy concentrations of x|0.2) and the unreacted aFe 2 O 3 . Comparisons of neutron diffraction patterns of the as-milled samples at room temperature and |950 K combined with chemical analyses indicate that a significant fraction of the milled products occurs in a disordered or amorphous form and that a maximum fraction of |8% g-Fe 2 O 3 is also present. Given that the predominant work on the a-Fe 2 O 3 powder for the low-energy mode used in the present mill takes place by shearing, this occurrence of g-Fe 2 O 3 is likely to be due to a shearing transformation involving movement of the oxygen planes [14]. Extended milling and continued exposure to impact as well as shearing is considered to lead to the rupturing of the oxide surface layers and release of oxygen with consequent reduction to Fe 32x O 4 [7]. In conclusion, while the occurrence of |2–3% Fe as a contaminant in the present low-energy milling experiments is likely to contribute via 4a-Fe 2 O 3 1Fe→3Fe 3 O 4 , the main transformation of a-Fe 2 O 3 to defect Fe 32x O 4 takes place by rupturing of the oxide surface layers of a-Fe 2 O 3 . This leads to release of oxygen and reduction to defect Fe 32x O 4 .
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Acknowledgements S.J.C. acknowledges renewal of an Alexander von Humboldt Research Fellowship while at the Johannes Gutenberg-University, Mainz. We acknowledge access to the LAD diffractometer at the ISIS Facility, Rutherford Appleton Laboratory, UK, and thank Dr. W.S. Howells for his assistance. We also thank Dr. A.V.J. Edge for assistance with the wet chemical analysis. This work was also supported in part by grants from the Australian Institute of Nuclear Science and Engineering and S.J.C. acknowledges support from the Access to Major Research Facilities Program, Australian Nuclear Science and Technology Organisation.
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