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Martensite–austenite transformation in Fe80Ni20 ball-milled powder
ARTICLE IN PRESS
Journal of Magnetism and Magnetic Materials 316 (2007) 328–331 www.elsevier.com/locate/jmmm
Martensite–austenite transformation in Fe80Ni20 ball-milled powder D. Martı´ nez-Blancoa, P. Gorriaa,, M.J. Pe´reza, J.A. Blancoa, R.I. Smithb a
Department of Physics, University of Oviedo, Calvo Sotelo s/n, Oviedo 33007, Spain b ISIS Facility, RAL, Chilton, Didcot, Oxon OX11 0QX, UK Available online 4 March 2007
Abstract The martensite–austenite (MA) transformation in Fe80Ni20 ball-milled powder has been studied through neutron thermo-diffraction experiments and magnetisation vs. temperature measurements between 300 and 1100 K. For the as-milled powder, the estimated size of mean coherent diffraction domains and lattice microstrains at room temperature (RT) are 11(2) nm and 0.8(2)%, respectively. The evolution of both parameters with temperature is presented and discussed. On heating the as-milled powder, the MA transformation begins above 630 K (200 K below the expected temperature for as-cast coarse-grained Fe80Ni20 conventional alloys) and finishes at 890 K. On cooling down from high temperature, the reverse transformation occurs below 400 K. These changes in the MA transformation temperatures are attributed to the initial nanostructure of the Fe80Ni20 material. r 2007 Published by Elsevier B.V. PACS: 75.50.Bb; 81.30.Kf; 61.12.Ld Keywords: Fe–Ni alloy; Mechanical alloying; Martensite transformation; Rietveld method
Modern fabrication techniques are used to design and develop new materials for technological applications. In particular, mechanical alloying or ball milling (BM) allows us to prepare metastable solid solutions with intermetallic elements, where structural and magnetic phase diagrams are still far from being completely determined [1]. One of such systems is FexNi100x in which BM procedure gives rise to disordered solid solutions in the whole compositional range, with average grain sizes in the nanometre scale and a large amount of microstrains that can reach 1% [2]. The crystal structure of these alloys depends on composition, that is, face-centred cubic (FCC) crystal structure is obtained for x435, while body-centred cubic (BCC) one does for xo20, with a mixture of both phases for intermediate compositions [2–4]. Moreover, it is well known that BCC-FeNi alloys prepared by conventional annealing suffer the so-called martensite–austenite (MA) transition (BCC to FCC, a–g) on heating above 650 K [5,6]. The temperatures at which the transition starts, As, and finishes, Af, depend on composition, same as, Corresponding author. Tel.: +34 985102899; fax: +34 985103324.
E-mail address: [email protected] (P. Gorria). 0304-8853/$ - see front matter r 2007 Published by Elsevier B.V. doi:10.1016/j.jmmm.2007.03.004
DTMA ¼ AsAf, which does not exceed 50 K [5]. Besides that, on cooling from high temperatures the reverse transformation, austenite–martensite (FCC to BCC, g–a0 ), takes place. However, the starting temperature, Ms, and the finishing temperature, Mf, are lower than As and Af, with differences that can reach 400 K [5] as the Ni content increases. Previous works show that this MA transformation also takes place in ball-milled BCC-FeNi alloys [2]. However, up to now most of the structural information on this subject comes from X-ray diffraction and Mo¨ssbauer spectroscopy measurements at room temperature (RT), performed on as-milled and annealed samples [7–9]. Hence, variations on relative percentages of austenite and martensite phases could appear depending on the annealing temperature and cooling rate. On the other hand, magnetisation vs. temperature curves, M(T), show an anomalous decrease above 700 K, which is a signature of an advanced MA transformation [2]. Hence, this M(T) behaviour suggests the need to carry out a detailed in situ structural study during the MA transformation process. Nowadays, powder neutron thermo-diffraction is considered as a very useful tool for studying the temperature
ARTICLE IN PRESS D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331
dependence of nuclear and/or magnetic structure in metastable systems [10]. We report on the structural evolution of BCC-Fe80Ni20 powders obtained from in situ neutron thermo-diffraction in the temperature range between 300 and 1100 K with the aim to better understand the observed magnetic behaviour of these metastable alloys. Pure (99.99%) Fe and Ni powders were milled in a highenergy ball mill (Retsch PM/400) during 30 h under controlled Ar atmosphere. The powder neutron thermodiffraction experiment was carried out at the POLARIS time-of-flight diffractometer (ISIS facility, UK) in the temperature range from 300 to 1100 K, with a heating rate of 2 K/min. Diffraction patterns were collected every 5 min. Rietveld refinement of neutron diffraction patterns has been realized using the Fullprof Suite Package [11]. Magnetisation vs. temperature, M(T), measurements between 300 and 1073 K and back again to 300 K were performed using a Faraday susceptometer under 1 kOe applied magnetic field. At RT, powder diffraction patterns show that the as-milled sample is nearly a BCC singlephase (a ¼ 2.873070.0005 A˚) with up to 5% of an FCC phase (a ¼ 3.58370.001 A˚). The particle size distribution was obtained from TEM images (180 keV JEOL-2000 EX-II model) by counting a large number of particles. In Fig. 1, M(T) curves for the as-milled sample (top), on heating and cooling procedures, together with the quantitative phase analysis obtained from neutron diffraction patterns (bottom) are shown. During the heating of the as-milled sample, a decrease in the magnetisation down to zero values is observed between 700 and 940 K whereas on cooling the sample, the magnetisation assumes a low and constant value down to 400 K, temperature at which M increases drastically. This behaviour differ from that expected for a conventional ferromagnetic system near the Curie temperature, TC. This could indicate that a gradual structural transformation from the ferromagnetic BCC–FeNi to the paramagnetic FCC–FeNi phase is taken place along this temperature interval. Powder neutron thermo-diffraction results confirm this point, the MA transformation begins at around 640 K, and then g-FeNi phase increases as the temperature is raised. The transformation is completed at T ¼ 920 K. Following the conventional criterion to assign the start (10% of austenite phase) and the end (90% of austenite phase) temperature for the MA transformation [5], it can be seen in Fig. 1 (bottom) that AsE630 K and AfE890 K. These temperatures do not exactly coincide with that of M(T) curve, probably due to the fact that the furnace and temperature controllers in both experiments are different. However, both measurements suggest that As is at least 150 K lower to that measured in as-cast coarse-grained conventional alloy and also that the temperature interval for the MA transition DTMA exceeds 200 K. These findings point out the important role played by non-equilibrium microstructure in its kind of solid-state transformations [2]. On the other hand, M(T) curve on cooling reveals that the
329
Fig. 1. Top: temperature dependence of the relative magnetisation M/ MRT (MRT corresponds to the magnetisation value measured at RT) for as-milled Fe80Ni20 sample. Bottom: temperature dependence of the percentage of the a-Fe80Ni20 and g-Fe80Ni20 phases on heating.
reverse MA transition, g-a0 , does not occur until the sample is cooled below 400 K, when magnetisation is drastically recovered (see Fig. 1). In order to better illustrate the MA transition, we show three neutron diffraction patterns (see Fig. 2) at selected temperatures during the heating process. These temperatures T1, T2 and T3 are marked in the quantitative phase analysis plot (Fig. 1). In Fig. 2, the diffraction pattern at T1 ¼ 600 K (upper panel) was taken before MA transition starts, where a-Fe80Ni20 phase is predominant. In the middle panel, the pattern at T1 ¼ 845 K is shown, where strong Bragg reflections of martensite phase are still present, more than 100 K after the MA transition begins. The third pattern (bottom panel) exhibits only sharp peaks belonging to g-Fe80Ni20 phase (T3 ¼ 915 K). On cooling down to RT, less than 5% of FCC phase is obtained from the fit of the diffraction pattern at RT [12], indicating that the reverse MA transformation is almost complete. Furthermore, the temperature dependence of the microstructure has been studied through the variation of the peak width of the Bragg reflections corresponding to the a-Fe80Ni20 phase [13]. Line-broadening analysis was done using a linear combination of spherical harmonics for modelling the ‘size’ broadening [14], while Stephens’
ARTICLE IN PRESS 330
D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331
Fig. 3. Temperature dependence of the microstructure (microstrain, e and apparent size, Dv) for the a-Fe80Ni20 phase, being derived from profile fitting of neutron diffraction pattern. The inset shows a histogram of grain-size distribution for the as-milled Fe80Ni20 sample obtained by TEM images at RT, and the fit to a log-normal distribution (see text).
Fig. 2. Fitted powder neutron diffraction patterns of a-Fe80Ni20 and gFe80Ni20 phases for three selected temperatures: T1 before beginning, T2 during and T3 after finishing the MA transformation. Observed (+); calculated (solid line); positions of the Bragg reflections are represented by vertical bars. The difference (observed–calculated) pattern is displayed at the bottom of each panel. Bragg reflections corresponding to traces of Fe oxides (below 4%) are also presented (Fe3O4 third row of vertical bars at T1 and T2, and FeO fourth and second rows of vertical bars at T2 and T3, respectively).
formulation [15] was used in order to determine the ‘microstrain’ broadening contribution. The apparent size, Dv, and microstrain, e, of the diffraction domains were deduced from the obtained parameters after fitting the profile of (h k l) reflections at all reciprocal directions. The mean values for Dv and e are represented in Fig. 3. The calculated initial values for the as-milled sample at RT are Dv ¼ 11(2) nm and e ¼ 0.8(2)%, respectively (where the number in the parenthesis is a measure of the degree of anisotropy, not of the estimated error). It is worth noting that for ToAs (630 K) the average grain size slightly increases its value (up to 15 nm), while the microstrain quickly decreases (down to 0.3%) due to thermal relaxation. The growth of the grain size is faster above 750 K, once the a–g transformation have started, thus
suggesting that both processes are interconnected. On the other hand, the microstrain diminishes at slower rate. Besides that, we compare the average grain size values obtained from diffraction analysis with those from dark field TEM images for the as-milled sample. Inset in Fig. 3 shows the histogram corresponding to the grain size distribution after counting a large number of grains in these images. The mean grain size diameter is obtained from a fit to a log-normal distribution with the characteristic parameters: tc ¼ 14.6(9) nm and s ¼ 0.38(2). This value is in good agreement with that estimated from the diffraction peak width of 11(2) nm, and it means that the microstructure has a quite homogenous distribution with a small anisotropy in this length scale. The experimental results of magnetisation vs. temperature measurements and powder neutron thermo-diffraction on Fe80Ni20 clearly indicate that the MA transformation spreads out for more than 150 K, exceeding the temperature range observed in as-cast coarse-grained Fe80Ni20 conventional alloys. This feature seems to be attributed to small size of diffraction domains and important microstrain effects that are still presented inside a-Fe80Ni20 grains. Acknowledgements The authors thank Spanish MEC for financial support under projects MAT2005-06806-C04-01 and MAT200306492. We also thank ISIS facility for neutron beam time, the SCT (University of Oviedo) for XRD facility and Carlos Alvarez Villa (TEM experiments). D.M.-B. thanks MEC for Ph.D. grant. References [1] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1. [2] C. Kuhrt, L. Schultz, J. Appl. Phys. 73 (1993) 1975.
ARTICLE IN PRESS D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331 [3] L. Hong, B. Fultz, J. Appl. Phys. 79 (1996) 3946. [4] S.D. Kaloshkin, V.V. Tcherdyntsev, I.A. Tomilin, Yu.V. Baldokhin, E.V. Shelekhov, Physica B 299 (2001) 236. [5] H.P.J. Wijn (Ed.), Magnetic Properties of Metals, d-Elements, Alloys and Compounds, Springer, Berlin, 1991. [6] R.M. Bozorth, Ferromagnetism, Wiley-IEEE Press, New York, 1993. [7] Yu.V. Baldohkin, V.V. Tcherdyntsev, S.D. Kaloshkin, G.A. Kochetov, Yu.A. Pustov, J. Magn. Magn. Mater. 203 (1999) 313. [8] M. Pekala. D. Oleszak, E. Jartych, J.K. Zurawicz, Nanostruct. Mater. 11 (1999) 789.
331
[9] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, J. Non-Cryst. Solids 287 (2001) 239. [10] P. Gorria, D. Martı´ nez-Blanco, J.A. Blanco, A. Hernando, J.S. Garitaonandia, L. Ferna´ndez Barquı´ n, J. Campo, R.I. Smith, Phys. Rev. B 69 (2004) 214421. [11] J. Rodrı´ guez Carvajal, Physica B 192 (1993) 55. [12] P. Gorria, D. Martı´ nez-Blanco, M.J. Pe´rez, J.A. Blanco, R.I. Smith, J. Magn. Magn. Mater. 294 (2005) 159. [13] R.A. Young (Ed.), The Rietveld Method, IUCRr Monographs on Crystallography 5, Oxford University Press, Oxford, 1993. [14] M. Ja¨rvinen, J. Appl. Cryst. 26 (1993) 525. [15] P.W. Stephens, J. Appl. Cryst. 32 (1999) 281.
Journal of Magnetism and Magnetic Materials 316 (2007) 328–331 www.elsevier.com/locate/jmmm
Martensite–austenite transformation in Fe80Ni20 ball-milled powder D. Martı´ nez-Blancoa, P. Gorriaa,, M.J. Pe´reza, J.A. Blancoa, R.I. Smithb a
Department of Physics, University of Oviedo, Calvo Sotelo s/n, Oviedo 33007, Spain b ISIS Facility, RAL, Chilton, Didcot, Oxon OX11 0QX, UK Available online 4 March 2007
Abstract The martensite–austenite (MA) transformation in Fe80Ni20 ball-milled powder has been studied through neutron thermo-diffraction experiments and magnetisation vs. temperature measurements between 300 and 1100 K. For the as-milled powder, the estimated size of mean coherent diffraction domains and lattice microstrains at room temperature (RT) are 11(2) nm and 0.8(2)%, respectively. The evolution of both parameters with temperature is presented and discussed. On heating the as-milled powder, the MA transformation begins above 630 K (200 K below the expected temperature for as-cast coarse-grained Fe80Ni20 conventional alloys) and finishes at 890 K. On cooling down from high temperature, the reverse transformation occurs below 400 K. These changes in the MA transformation temperatures are attributed to the initial nanostructure of the Fe80Ni20 material. r 2007 Published by Elsevier B.V. PACS: 75.50.Bb; 81.30.Kf; 61.12.Ld Keywords: Fe–Ni alloy; Mechanical alloying; Martensite transformation; Rietveld method
Modern fabrication techniques are used to design and develop new materials for technological applications. In particular, mechanical alloying or ball milling (BM) allows us to prepare metastable solid solutions with intermetallic elements, where structural and magnetic phase diagrams are still far from being completely determined [1]. One of such systems is FexNi100x in which BM procedure gives rise to disordered solid solutions in the whole compositional range, with average grain sizes in the nanometre scale and a large amount of microstrains that can reach 1% [2]. The crystal structure of these alloys depends on composition, that is, face-centred cubic (FCC) crystal structure is obtained for x435, while body-centred cubic (BCC) one does for xo20, with a mixture of both phases for intermediate compositions [2–4]. Moreover, it is well known that BCC-FeNi alloys prepared by conventional annealing suffer the so-called martensite–austenite (MA) transition (BCC to FCC, a–g) on heating above 650 K [5,6]. The temperatures at which the transition starts, As, and finishes, Af, depend on composition, same as, Corresponding author. Tel.: +34 985102899; fax: +34 985103324.
E-mail address: [email protected] (P. Gorria). 0304-8853/$ - see front matter r 2007 Published by Elsevier B.V. doi:10.1016/j.jmmm.2007.03.004
DTMA ¼ AsAf, which does not exceed 50 K [5]. Besides that, on cooling from high temperatures the reverse transformation, austenite–martensite (FCC to BCC, g–a0 ), takes place. However, the starting temperature, Ms, and the finishing temperature, Mf, are lower than As and Af, with differences that can reach 400 K [5] as the Ni content increases. Previous works show that this MA transformation also takes place in ball-milled BCC-FeNi alloys [2]. However, up to now most of the structural information on this subject comes from X-ray diffraction and Mo¨ssbauer spectroscopy measurements at room temperature (RT), performed on as-milled and annealed samples [7–9]. Hence, variations on relative percentages of austenite and martensite phases could appear depending on the annealing temperature and cooling rate. On the other hand, magnetisation vs. temperature curves, M(T), show an anomalous decrease above 700 K, which is a signature of an advanced MA transformation [2]. Hence, this M(T) behaviour suggests the need to carry out a detailed in situ structural study during the MA transformation process. Nowadays, powder neutron thermo-diffraction is considered as a very useful tool for studying the temperature
ARTICLE IN PRESS D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331
dependence of nuclear and/or magnetic structure in metastable systems [10]. We report on the structural evolution of BCC-Fe80Ni20 powders obtained from in situ neutron thermo-diffraction in the temperature range between 300 and 1100 K with the aim to better understand the observed magnetic behaviour of these metastable alloys. Pure (99.99%) Fe and Ni powders were milled in a highenergy ball mill (Retsch PM/400) during 30 h under controlled Ar atmosphere. The powder neutron thermodiffraction experiment was carried out at the POLARIS time-of-flight diffractometer (ISIS facility, UK) in the temperature range from 300 to 1100 K, with a heating rate of 2 K/min. Diffraction patterns were collected every 5 min. Rietveld refinement of neutron diffraction patterns has been realized using the Fullprof Suite Package [11]. Magnetisation vs. temperature, M(T), measurements between 300 and 1073 K and back again to 300 K were performed using a Faraday susceptometer under 1 kOe applied magnetic field. At RT, powder diffraction patterns show that the as-milled sample is nearly a BCC singlephase (a ¼ 2.873070.0005 A˚) with up to 5% of an FCC phase (a ¼ 3.58370.001 A˚). The particle size distribution was obtained from TEM images (180 keV JEOL-2000 EX-II model) by counting a large number of particles. In Fig. 1, M(T) curves for the as-milled sample (top), on heating and cooling procedures, together with the quantitative phase analysis obtained from neutron diffraction patterns (bottom) are shown. During the heating of the as-milled sample, a decrease in the magnetisation down to zero values is observed between 700 and 940 K whereas on cooling the sample, the magnetisation assumes a low and constant value down to 400 K, temperature at which M increases drastically. This behaviour differ from that expected for a conventional ferromagnetic system near the Curie temperature, TC. This could indicate that a gradual structural transformation from the ferromagnetic BCC–FeNi to the paramagnetic FCC–FeNi phase is taken place along this temperature interval. Powder neutron thermo-diffraction results confirm this point, the MA transformation begins at around 640 K, and then g-FeNi phase increases as the temperature is raised. The transformation is completed at T ¼ 920 K. Following the conventional criterion to assign the start (10% of austenite phase) and the end (90% of austenite phase) temperature for the MA transformation [5], it can be seen in Fig. 1 (bottom) that AsE630 K and AfE890 K. These temperatures do not exactly coincide with that of M(T) curve, probably due to the fact that the furnace and temperature controllers in both experiments are different. However, both measurements suggest that As is at least 150 K lower to that measured in as-cast coarse-grained conventional alloy and also that the temperature interval for the MA transition DTMA exceeds 200 K. These findings point out the important role played by non-equilibrium microstructure in its kind of solid-state transformations [2]. On the other hand, M(T) curve on cooling reveals that the
329
Fig. 1. Top: temperature dependence of the relative magnetisation M/ MRT (MRT corresponds to the magnetisation value measured at RT) for as-milled Fe80Ni20 sample. Bottom: temperature dependence of the percentage of the a-Fe80Ni20 and g-Fe80Ni20 phases on heating.
reverse MA transition, g-a0 , does not occur until the sample is cooled below 400 K, when magnetisation is drastically recovered (see Fig. 1). In order to better illustrate the MA transition, we show three neutron diffraction patterns (see Fig. 2) at selected temperatures during the heating process. These temperatures T1, T2 and T3 are marked in the quantitative phase analysis plot (Fig. 1). In Fig. 2, the diffraction pattern at T1 ¼ 600 K (upper panel) was taken before MA transition starts, where a-Fe80Ni20 phase is predominant. In the middle panel, the pattern at T1 ¼ 845 K is shown, where strong Bragg reflections of martensite phase are still present, more than 100 K after the MA transition begins. The third pattern (bottom panel) exhibits only sharp peaks belonging to g-Fe80Ni20 phase (T3 ¼ 915 K). On cooling down to RT, less than 5% of FCC phase is obtained from the fit of the diffraction pattern at RT [12], indicating that the reverse MA transformation is almost complete. Furthermore, the temperature dependence of the microstructure has been studied through the variation of the peak width of the Bragg reflections corresponding to the a-Fe80Ni20 phase [13]. Line-broadening analysis was done using a linear combination of spherical harmonics for modelling the ‘size’ broadening [14], while Stephens’
ARTICLE IN PRESS 330
D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331
Fig. 3. Temperature dependence of the microstructure (microstrain, e and apparent size, Dv) for the a-Fe80Ni20 phase, being derived from profile fitting of neutron diffraction pattern. The inset shows a histogram of grain-size distribution for the as-milled Fe80Ni20 sample obtained by TEM images at RT, and the fit to a log-normal distribution (see text).
Fig. 2. Fitted powder neutron diffraction patterns of a-Fe80Ni20 and gFe80Ni20 phases for three selected temperatures: T1 before beginning, T2 during and T3 after finishing the MA transformation. Observed (+); calculated (solid line); positions of the Bragg reflections are represented by vertical bars. The difference (observed–calculated) pattern is displayed at the bottom of each panel. Bragg reflections corresponding to traces of Fe oxides (below 4%) are also presented (Fe3O4 third row of vertical bars at T1 and T2, and FeO fourth and second rows of vertical bars at T2 and T3, respectively).
formulation [15] was used in order to determine the ‘microstrain’ broadening contribution. The apparent size, Dv, and microstrain, e, of the diffraction domains were deduced from the obtained parameters after fitting the profile of (h k l) reflections at all reciprocal directions. The mean values for Dv and e are represented in Fig. 3. The calculated initial values for the as-milled sample at RT are Dv ¼ 11(2) nm and e ¼ 0.8(2)%, respectively (where the number in the parenthesis is a measure of the degree of anisotropy, not of the estimated error). It is worth noting that for ToAs (630 K) the average grain size slightly increases its value (up to 15 nm), while the microstrain quickly decreases (down to 0.3%) due to thermal relaxation. The growth of the grain size is faster above 750 K, once the a–g transformation have started, thus
suggesting that both processes are interconnected. On the other hand, the microstrain diminishes at slower rate. Besides that, we compare the average grain size values obtained from diffraction analysis with those from dark field TEM images for the as-milled sample. Inset in Fig. 3 shows the histogram corresponding to the grain size distribution after counting a large number of grains in these images. The mean grain size diameter is obtained from a fit to a log-normal distribution with the characteristic parameters: tc ¼ 14.6(9) nm and s ¼ 0.38(2). This value is in good agreement with that estimated from the diffraction peak width of 11(2) nm, and it means that the microstructure has a quite homogenous distribution with a small anisotropy in this length scale. The experimental results of magnetisation vs. temperature measurements and powder neutron thermo-diffraction on Fe80Ni20 clearly indicate that the MA transformation spreads out for more than 150 K, exceeding the temperature range observed in as-cast coarse-grained Fe80Ni20 conventional alloys. This feature seems to be attributed to small size of diffraction domains and important microstrain effects that are still presented inside a-Fe80Ni20 grains. Acknowledgements The authors thank Spanish MEC for financial support under projects MAT2005-06806-C04-01 and MAT200306492. We also thank ISIS facility for neutron beam time, the SCT (University of Oviedo) for XRD facility and Carlos Alvarez Villa (TEM experiments). D.M.-B. thanks MEC for Ph.D. grant. References [1] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1. [2] C. Kuhrt, L. Schultz, J. Appl. Phys. 73 (1993) 1975.
ARTICLE IN PRESS D. Martı´nez-Blanco et al. / Journal of Magnetism and Magnetic Materials 316 (2007) 328–331 [3] L. Hong, B. Fultz, J. Appl. Phys. 79 (1996) 3946. [4] S.D. Kaloshkin, V.V. Tcherdyntsev, I.A. Tomilin, Yu.V. Baldokhin, E.V. Shelekhov, Physica B 299 (2001) 236. [5] H.P.J. Wijn (Ed.), Magnetic Properties of Metals, d-Elements, Alloys and Compounds, Springer, Berlin, 1991. [6] R.M. Bozorth, Ferromagnetism, Wiley-IEEE Press, New York, 1993. [7] Yu.V. Baldohkin, V.V. Tcherdyntsev, S.D. Kaloshkin, G.A. Kochetov, Yu.A. Pustov, J. Magn. Magn. Mater. 203 (1999) 313. [8] M. Pekala. D. Oleszak, E. Jartych, J.K. Zurawicz, Nanostruct. Mater. 11 (1999) 789.
331
[9] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, J. Non-Cryst. Solids 287 (2001) 239. [10] P. Gorria, D. Martı´ nez-Blanco, J.A. Blanco, A. Hernando, J.S. Garitaonandia, L. Ferna´ndez Barquı´ n, J. Campo, R.I. Smith, Phys. Rev. B 69 (2004) 214421. [11] J. Rodrı´ guez Carvajal, Physica B 192 (1993) 55. [12] P. Gorria, D. Martı´ nez-Blanco, M.J. Pe´rez, J.A. Blanco, R.I. Smith, J. Magn. Magn. Mater. 294 (2005) 159. [13] R.A. Young (Ed.), The Rietveld Method, IUCRr Monographs on Crystallography 5, Oxford University Press, Oxford, 1993. [14] M. Ja¨rvinen, J. Appl. Cryst. 26 (1993) 525. [15] P.W. Stephens, J. Appl. Cryst. 32 (1999) 281.