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Structure and magnetic properties in CoCu granular alloys
Nuclear Instruments and Methods in Physics Research B 200 (2003) 215–219 www.elsevier.com/locate/nimb
Structure and magnetic properties in CoCu granular alloys C. Meneghini a
a,b,*
, S. Mobilio
a,c
, A. Garcıa-Prieto d, M.L.F. Fdez-Gubieda
d
Dip. di Fisica E. Amaldi, Universit a di Roma Tre, Via della Vasca Navale 84, I-00146 Roma, Italy b INFM-GILDA c/o ESRF Grenoble, France c Laboratori Nazionali di Frascati dellÕINFN, via E. Fermi 40, I-00044 Frascati, Italy d Departamento de Electricidad y Electr onica, UPV-EHU, Apartado 644, Bilbao 48080, Spain
Abstract Granular alloys, composed of magnetic clusters embedded in non-magnetic metallic matrices, can develop giant magnetoresistance effect after suitable preparation and thermal treatments. The structural effect of annealing on the structure of Co10 Cu90 samples has been directly probed by in situ time resolved X-ray diffraction (TR-XRD) during thermal treatment. TR-XRD definitively proves the occurrence of an anomalous behaviour in the thermally activated segregation process that is related to the evolution of magnetotransport properties in these materials. Ó 2002 Elsevier Science B.V. All rights reserved. PACS: 61.72.Cc; 61.10.Nz; 75.47.De Keywords: Giant magnetoresistance; Granular alloys; Time resolved-XRD
1. Introduction The magnetoresistance (MR) effect consists on change of electrical resistance under an applied magnetic field. Granular alloys, which are composed of magnetic clusters embedded in a nonmagnetic matrix, stimulated a large research interest since the discovery [1,2] that they can develop a giant magnetoresistance (GMR) effect. The current understanding of GMR phenomena in granular alloys is that they mainly arise from the spin-dependent scattering that takes place at the
*
Corresponding author. Tel.: +39-06-55177054; fax: +39-065579303. E-mail address: meneghini@fis.uniroma3.it (C. Meneghini).
interfaces between the magnetic granules and the matrix. The magnetotransport properties are also influenced by the mean free path in the magnetic and non-magnetic regions, thus by the composition, shape and distribution of magnetic particles within the matrix [3–5]. The knowledge of the sample microstructure is crucial for deeply understanding the peculiar properties of these materials. Co–Cu alloys are among the systems presenting GMR. Co and Cu are immiscible elements at room temperature but metastable solid solutions can be prepared by special methods. After preparation, suitable thermal treatments that cause spontaneous partial segregation of the components favour the evolution of the system in the form of a granular alloy and can enhance the MR. Although granular Co–Cu alloys have been extensively
0168-583X/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 2 ) 0 1 7 2 2 - 6
216
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
studied in the last years [6–8], there are no conclusive results about their real microstructure, mostly due to the experimental difficulty of resolving small structures out of a matrix when the elements involved have very close atomic numbers, as for Co and Cu. The relevant information gathered about these systems mainly comes from macroscopic evidence, like magnetic and transport properties from which the microscopic information available is deduced using more or less accurate models [6,8,9]. An interesting aspect of granular Co–Cu systems is that the maximum of MR is reached after annealing the sample around 450 °C, while the MR rapidly falls for higher annealing temperatures [10,11]. This behaviour is largely independent from the sample composition, preparation method (melt spun, laser ablation or electro-deposition) and the form of the sample form (bulk or films). In previous works we combined X-ray absorption fine structure (EXAFS), high resolution X-ray diffraction (XRD) and standard magnetic and magnetotransport characterization [10–12], in order to relate the evolution of the microstructure in Cox Cu1x alloys to changes of magnetotransport properties as a function of composition and annealing temperature (Ta ). The net result coming out from these works is that the MR firstly increases as a function of Ta (i.e. Ta 6 450 °C) due to the segregation of small (2–3 nm) superparamagnetic Co clusters into the Cu-rich matrix. On the contrary raising Ta above 450–500 °C, worsens the MR properties of the material. EXAFS and XRD results demonstrated that the deterioration of magnetotransport is triggered by an anomaly in the segregation process that provokes the partial remixing of the elements causing the roughening of Co–Cu interfaces. The purpose of this work is to get deeper insight on such anomalous effects probing the structural evolution of granular CoCu samples as a function of temperature in situ, during the annealing process. This has been done exploiting the time resolved XRD (TR-XRD) technique on Co10 Cu90 samples prepared by melt spun and heated in situ between 300 and 700 °C. Different heating speed were used in order to reveal the effect of heating kinetics.
2. Experiment and data analysis Co10 Cu90 samples in form of long ribbon about 30 mm thick have been prepared by rapid quenching in He controlled atmosphere. Different pieces of the same ribbon, annealed ex situ (Ta ¼ 400–650 °C) have been already carefully investigated by XRD, EXAFS as well as magnetic and magnetotransport techniques [10–12]. TR-XRD experiments have been performed on the GILDA beam-line at the ESRF (Grenoble, France) using an angle dispersed set-up based on a translatingimaging plate (TIP) camera [13]. The monochro-
Fig. 1. Panel (a): typical TR-XRD image collected heating the Co10 Cu90 sample between 300 and 700 °C (VT ¼ 2 °C/min). The inset details the diffraction lines corresponding to the more intense Cu-rich phase (a) and to the weaker Co-rich phase (b). Panels (b) and (c) display an experimental diffraction pattern (þ) obtained integrating a 2 mm wide vertical stripe (t) of the digitised image. The Rietveld refined pattern (line) and its residual (shifted for clarity) are also shown. Markers indicate the refined Bragg position of Cu- and Co-rich diffraction lines.
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
. Pieces of matic beam wavelength was 0.68795 A the same original ribbon have been enclosed in cylindrical quartz capillaries (inner diameter 0.3 mm). Capillaries have been mounted horizontally on a goniometer head, aligned on the beam and kept rotating during acquisition in order to improve the grain statistics. Samples were heated using a gas heater. Instrumental parameters (sample to TIP distance, temperature and so on) were calibrated by refining diffractograms collected of reference samples (c-Si and a pure Cu foil) in the same experimental conditions. Time resolved diffractograms were collected by heating the sample between 300 and 700 °C at different heating ratios: VT ¼ 6:7, 4, 2.7 and 2 °C/min. Diffraction patterns, stored in the IP latent image, have been digitised using a BAS2500 reader with a resolution of 100 100 lm2 /pixel and 16 bit of dynamical range (Fig. 1). Digitised images were processed with the program ScanZero [14] to extract 90 diffraction patterns as a function of temperature with temperature step of 4.4 °C. The extracted diffraction patterns have been quantitatively analysed within the Rietveld structural refinement approach as implemented in the GSAS package [15] (Fig. 1(b) and (c)) using a pseudoVoigth profile function.
3. Results and discussion Before we discuss the in situ results let us resume the knowledge on these compounds on the basis of previous investigations [10–12]. XRD experiments revealed the presence of two fcc phases also in the as-quenched ribbon: the one richer in Cu, the other richer in Co. Thermal treatment provokes the progressive expulsion of minority elements from metastable phases: Co atoms are expelled from the Cu-rich phase; Cu atoms are expelled out from the Co-rich one. This, according to the VegardÕs law, results in the progressive increase (decrease) of Cu-rich (Co-rich) lattice parameter toward the value expected in pure Cu, that (Co-fcc, aCo-fcc ¼ 3:544 A ). is aCu-fcc ¼ 3:615 A However Cu-rich and Co-rich phases evolves differently as a function of Ta : the segregation process in Cu-rich phase is activated only heating the
217
samples at Ta > 550–600 °C, in Co-rich phase it starts earlier (Ta > 400 °C) but it stops for Ta ¼ 500 °C that is the same temperature at which the MR start to decrease. Combined EXAFS and magnetic measurements allowed demonstrating that this anomalous behaviour signals a partial remixing of the elements causing the roughening of Co–Cu interfaces and worsening the MR. This phenomenon has been observed in all the Cox Cu1x samples investigated, despite the composition (x ¼ 5, 10, 15, 20), suggesting that it must derive from peculiar features in the phase diagram and/or in the thermodynamics of Co and Cu solution (as supposed in [16]) that must favour the a partial re-mixing of the immiscible elements. In situ experiments directly probe the effect of temperature during heating allowing to get deeper insights on this phenomenon. TR-XRD patterns, according to previous high-resolution XRD experiments [10–12], show two fcc phases (Fig. 1(a)). The refined structures definitively show the phase separation in a Cu-rich phase, corresponding to the most intense diffraction lines, and a Co-rich one, having weaker peaks (Fig. 1(b)). Raising the annealing temperature provokes the lattice expansion in both the phases (Fig. 2(a)) but the effect is definitively different in Co- and Cu-rich phases. The lattice parameter of the Co-rich phase (aCo ) increases linearly heating the sample between 300 and 700 °C while the expansion of Cu-rich phase (aCu ) depart from the linear trend heating above 600–650 °C. Linear expansion coefficients for Corich (kCo ) and Cu-rich (kCu ) phases have been estimated by linear fitting ai ðT Þ curves. The kCu ¼ 1:69 105 K1 , obtained from the aCu ðT Þ data (T < 600 °C), is the same of pure Cu-fcc (aCo ¼ 1:7 105 K1 [17]) and independent from heating speed (VT ). Above 600 °C the slope of aCu ðT Þ suddenly increases pointing out the activated segregation process in the Cu-rich phase that raises the lattice parameter toward the value expected in pure Cu. This effect is larger for the slower heating ramps. This finding comes in agreement with ex situ measurements showing that the Cu-rich phase does not change appreciably until annealing above 550–600 °C [10–12]. The Co-rich phase thermal expansion follows a linear behaviour in the whole temperature range
218
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
3.66
(a)
a (Å)
a
Cu
3.65 -5
-1
k = 1.69e K 3.64
a + 0.05 Co
-5
k = 1.36e K
3.63
∆a (Å) 0.07
-1
(b)
2
slower heating provokes the larger departure of Da from a linear trend signalling the larger segregation effect. Additional information on the structural evolution under annealing can be derived looking at the diffraction line profile and diffraction intensity (Fig. 3). The gaussian contribution (G) to the profile function reflects the distribution of lattice parameters resulting from stress/strain effects as well as from composition spread. The lorentzian (L) contribution, on the contrary, is mainly related to the crystallite size.
2.7
(x 10 deg)
4
500
600
700
o
T ( C)
2.7 32
4
28
6.7
(x 10 deg)
(b)
2
-2
400
2.7
G
Co
300 4 200 6.7
100
(c)
2
1.8
2.7 Co
with kCo ¼ 1:36 105 K1 , lower than that for pure Co-fcc (aCo ¼ 1:4 105 K1 [18]) and almost independent on the heating ratios. This effect is due to the competition of two effects: the thermal expansion, raising the lattice parameter, and the expulsion of Cu foreign atoms that according to the VegardÕs law reduces the aCo ðT Þ toward the values expected in pure Co-fcc. The evolution of aCo as a function of annealing temperature does not show the anomalies reported in ex situ experiments. This is probably related to the different annealing kinetics and deserves further investigation. The differences DaðT Þ ¼ aCu aCo (Fig. 2(b)) have all the same slope in the T < 600 °C region, suggesting that the heating kinetics affects only weakly the low temperature segregation process. On the contrary in the high temperature region the
36
500
H
Fig. 2. The lattice parameters for Cu-rich (aCu ) and Co-rich (aCo ) phases (VT ¼ 2 °C/min) as a function of annealing temperature are shown shifted for clarity (panel (a)). Full lines are linear fit to the data (see text). Panel (b) reports the differences Da ¼ aCu aCo as a function of heating speed (VT ). The numbers represent the VT for each data set. DaðT Þ depict a definitive departure from a linear trend heating above 600 °C, such an effect is more pronounced for the slower heating curves. Straight lines are guide for the eyes.
(a) 2
Cu
400
-2
6.7
40
L
0.06
1.4 4 1 6.7 400
500
600
700
o
T( C) Fig. 3. Results of diffraction profile refinement: the lorentzian contribution to the Cu-rich line shape (panel (a)) and the gaussian contribution to the Co-rich diffraction profile (panel (b)) are shown as a function of temperature for different VT (VT are reported for each data set). Panel (c) reports the histogram (HCo ) scale factor for Co-rich phase (see text).
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
Sharper diffraction lines characterize the Curich contribution to the diffractograms. The gaussian term in the Cu-rich phase peaks profile is about 0.15–0.2°, roughly constant as a function of T , and largely independent from VT . On the contrary the lorentzian contribution (LCu , Fig. 2(a)) depicts a clear temperature dependence: LCu ðT Þ decreases as a function of T in all the four kinetics, pointing out the progressive growth of crystallites. However, all the LCu ðT Þ curves depict an evident change of slope on heating the sample across 450– 500 °C. This effect appears independent on VT . Broader and weaker diffraction lines characterize the contribution of the Co-rich phase which profile appears dominated by the gaussian contribution (GCo ), while the lorentzian term is negligible. GCo ðT Þ is more than one order of magnitude larger than in Cu-rich phase and, interestingly, depicts an evident maximum as a function of T around 500 °C, which amplitude and position is largely independent on VT . This finding suggests a larger spread of Co-rich phase lattice parameters heating the sample between 450–500 °C signifying the activation of a remixing process in the Co-rich phase. The histogram scale factor (Hi ) of GSAS accounts for the relative amount of each phase within the sample. The relative weight of Cu-rich phase, HCu , is roughly independent from T . On the contrary the Co-rich scale factor (HCo ) depicts evident temperature dependence (Fig. 3(c)) with a clear minimum around 500 °C that seem unaffected by changing VT . This finding comes in agreement with the previous hypothesis of a Co– Cu remixing process activated on heating the sample around 500 °C that would cause the partial dissolution of Co-rich phase. TR-XRD results definitively prove that Co- and Cu-rich phases evolve differently as a function of annealing temperature. The thermally activated segregation process in the Co-rich phase depicts an anomalous behaviour that is largely independent on the heating kinetics.
219
Acknowledgements We acknowledge the excellent technical support of F. Campolungo, V. Sciarra, V. Tullio (INFNLNF) and F. DÕAnca (INFM-OGG Grenoble). Gilda Beamline is financed by the Italian institutions CNR, INFM and INFN. This work has been partially supported by the Spanish CICyT under project no. MAT99-0667.
References [1] A.E. Berkowitz, J.R. Mitchell, M.J. Carey, A.P. Young, S. Zhang, F.E. Spada, F.T. Parjer, A. Hutten, G. Thomas, Phys. Rev. Lett. 68 (1992) 3745. [2] J.Q. Xiao, J.S. Jiang, C.L. Chien, Phys. Rev. Lett. 68 (1992) 3749. [3] S. Zhang, P.M. Levy, J. Appl. Phys. 73 (1993) 5315. [4] M. Rubinstein, Phys. Rev. B 50 (1994) 3830. [5] R.Y. Gu et al., Phys. Rev. B 53 (1996) 11685. [6] P. Allia, M. Knobel, P. Tiberto, F. Vinai, Phys. Rev. B 52 (1995) 15398. [7] B.J. Hickey et al., Phys. Rev. B 51 (1995) 667. [8] F.C.S. da Silva, E.F. Ferrari, M. Knobel, J. Appl. Phys. 86 (1999) 7170. [9] A.D.C. Viegas et al., J. Appl. Phys. 82 (1997) 3047. [10] A. Garcıa-Prieto, M.L. Fdez-Gubieda, A. Garcıa Arribas, J.M. Barandiaran, C. Meneghini, S. Mobilio, J. Magn. Magn. Mater. 221 (2000) 80. [11] M.L. Fdez-Gubieda, A. Garcıa-Prieto, A. Garcıa-Arribas, C. Meneghini, S. Mobilio, ESRF Highlights (2000) 73. [12] M.L. Fdez-Gubieda, A. Garcıa-Prieto, A. Garcıa-Arribas, C. Meneghini, S. Mobilio, Europhys. Lett. 59 (2002) 855. [13] C. Meneghini, G. Artioli, A. Balerna, A.F. Gualtieri, P. Norby, S. Mobilio, J. Synchr. Rad. 8 (2001) 1162. [14] C. Meneghini, C. Dalconi, G. Cruciani, unpublished. [15] A.C. Larson, R.B. Von Dreele, LANSCE, MS-H805, Los Alamos National Laboratory. [16] Z.F. Li, Q. Zhang, D.P. Yu, C. Lin, B.X. Liu, Phys. Rev. B 64 (2001) 14102. [17] J. Rodrıguez-Carvajal, computer code FULLPROF, Rietveld Pattern Matching Analysis of Powder Patterns, ILL, Grenoble, 1990. [18] R.M. Bozorth, Ferromagnetism, D. Van Nostram Company, 1964.
Structure and magnetic properties in CoCu granular alloys C. Meneghini a
a,b,*
, S. Mobilio
a,c
, A. Garcıa-Prieto d, M.L.F. Fdez-Gubieda
d
Dip. di Fisica E. Amaldi, Universit a di Roma Tre, Via della Vasca Navale 84, I-00146 Roma, Italy b INFM-GILDA c/o ESRF Grenoble, France c Laboratori Nazionali di Frascati dellÕINFN, via E. Fermi 40, I-00044 Frascati, Italy d Departamento de Electricidad y Electr onica, UPV-EHU, Apartado 644, Bilbao 48080, Spain
Abstract Granular alloys, composed of magnetic clusters embedded in non-magnetic metallic matrices, can develop giant magnetoresistance effect after suitable preparation and thermal treatments. The structural effect of annealing on the structure of Co10 Cu90 samples has been directly probed by in situ time resolved X-ray diffraction (TR-XRD) during thermal treatment. TR-XRD definitively proves the occurrence of an anomalous behaviour in the thermally activated segregation process that is related to the evolution of magnetotransport properties in these materials. Ó 2002 Elsevier Science B.V. All rights reserved. PACS: 61.72.Cc; 61.10.Nz; 75.47.De Keywords: Giant magnetoresistance; Granular alloys; Time resolved-XRD
1. Introduction The magnetoresistance (MR) effect consists on change of electrical resistance under an applied magnetic field. Granular alloys, which are composed of magnetic clusters embedded in a nonmagnetic matrix, stimulated a large research interest since the discovery [1,2] that they can develop a giant magnetoresistance (GMR) effect. The current understanding of GMR phenomena in granular alloys is that they mainly arise from the spin-dependent scattering that takes place at the
*
Corresponding author. Tel.: +39-06-55177054; fax: +39-065579303. E-mail address: meneghini@fis.uniroma3.it (C. Meneghini).
interfaces between the magnetic granules and the matrix. The magnetotransport properties are also influenced by the mean free path in the magnetic and non-magnetic regions, thus by the composition, shape and distribution of magnetic particles within the matrix [3–5]. The knowledge of the sample microstructure is crucial for deeply understanding the peculiar properties of these materials. Co–Cu alloys are among the systems presenting GMR. Co and Cu are immiscible elements at room temperature but metastable solid solutions can be prepared by special methods. After preparation, suitable thermal treatments that cause spontaneous partial segregation of the components favour the evolution of the system in the form of a granular alloy and can enhance the MR. Although granular Co–Cu alloys have been extensively
0168-583X/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 2 ) 0 1 7 2 2 - 6
216
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
studied in the last years [6–8], there are no conclusive results about their real microstructure, mostly due to the experimental difficulty of resolving small structures out of a matrix when the elements involved have very close atomic numbers, as for Co and Cu. The relevant information gathered about these systems mainly comes from macroscopic evidence, like magnetic and transport properties from which the microscopic information available is deduced using more or less accurate models [6,8,9]. An interesting aspect of granular Co–Cu systems is that the maximum of MR is reached after annealing the sample around 450 °C, while the MR rapidly falls for higher annealing temperatures [10,11]. This behaviour is largely independent from the sample composition, preparation method (melt spun, laser ablation or electro-deposition) and the form of the sample form (bulk or films). In previous works we combined X-ray absorption fine structure (EXAFS), high resolution X-ray diffraction (XRD) and standard magnetic and magnetotransport characterization [10–12], in order to relate the evolution of the microstructure in Cox Cu1x alloys to changes of magnetotransport properties as a function of composition and annealing temperature (Ta ). The net result coming out from these works is that the MR firstly increases as a function of Ta (i.e. Ta 6 450 °C) due to the segregation of small (2–3 nm) superparamagnetic Co clusters into the Cu-rich matrix. On the contrary raising Ta above 450–500 °C, worsens the MR properties of the material. EXAFS and XRD results demonstrated that the deterioration of magnetotransport is triggered by an anomaly in the segregation process that provokes the partial remixing of the elements causing the roughening of Co–Cu interfaces. The purpose of this work is to get deeper insight on such anomalous effects probing the structural evolution of granular CoCu samples as a function of temperature in situ, during the annealing process. This has been done exploiting the time resolved XRD (TR-XRD) technique on Co10 Cu90 samples prepared by melt spun and heated in situ between 300 and 700 °C. Different heating speed were used in order to reveal the effect of heating kinetics.
2. Experiment and data analysis Co10 Cu90 samples in form of long ribbon about 30 mm thick have been prepared by rapid quenching in He controlled atmosphere. Different pieces of the same ribbon, annealed ex situ (Ta ¼ 400–650 °C) have been already carefully investigated by XRD, EXAFS as well as magnetic and magnetotransport techniques [10–12]. TR-XRD experiments have been performed on the GILDA beam-line at the ESRF (Grenoble, France) using an angle dispersed set-up based on a translatingimaging plate (TIP) camera [13]. The monochro-
Fig. 1. Panel (a): typical TR-XRD image collected heating the Co10 Cu90 sample between 300 and 700 °C (VT ¼ 2 °C/min). The inset details the diffraction lines corresponding to the more intense Cu-rich phase (a) and to the weaker Co-rich phase (b). Panels (b) and (c) display an experimental diffraction pattern (þ) obtained integrating a 2 mm wide vertical stripe (t) of the digitised image. The Rietveld refined pattern (line) and its residual (shifted for clarity) are also shown. Markers indicate the refined Bragg position of Cu- and Co-rich diffraction lines.
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
. Pieces of matic beam wavelength was 0.68795 A the same original ribbon have been enclosed in cylindrical quartz capillaries (inner diameter 0.3 mm). Capillaries have been mounted horizontally on a goniometer head, aligned on the beam and kept rotating during acquisition in order to improve the grain statistics. Samples were heated using a gas heater. Instrumental parameters (sample to TIP distance, temperature and so on) were calibrated by refining diffractograms collected of reference samples (c-Si and a pure Cu foil) in the same experimental conditions. Time resolved diffractograms were collected by heating the sample between 300 and 700 °C at different heating ratios: VT ¼ 6:7, 4, 2.7 and 2 °C/min. Diffraction patterns, stored in the IP latent image, have been digitised using a BAS2500 reader with a resolution of 100 100 lm2 /pixel and 16 bit of dynamical range (Fig. 1). Digitised images were processed with the program ScanZero [14] to extract 90 diffraction patterns as a function of temperature with temperature step of 4.4 °C. The extracted diffraction patterns have been quantitatively analysed within the Rietveld structural refinement approach as implemented in the GSAS package [15] (Fig. 1(b) and (c)) using a pseudoVoigth profile function.
3. Results and discussion Before we discuss the in situ results let us resume the knowledge on these compounds on the basis of previous investigations [10–12]. XRD experiments revealed the presence of two fcc phases also in the as-quenched ribbon: the one richer in Cu, the other richer in Co. Thermal treatment provokes the progressive expulsion of minority elements from metastable phases: Co atoms are expelled from the Cu-rich phase; Cu atoms are expelled out from the Co-rich one. This, according to the VegardÕs law, results in the progressive increase (decrease) of Cu-rich (Co-rich) lattice parameter toward the value expected in pure Cu, that (Co-fcc, aCo-fcc ¼ 3:544 A ). is aCu-fcc ¼ 3:615 A However Cu-rich and Co-rich phases evolves differently as a function of Ta : the segregation process in Cu-rich phase is activated only heating the
217
samples at Ta > 550–600 °C, in Co-rich phase it starts earlier (Ta > 400 °C) but it stops for Ta ¼ 500 °C that is the same temperature at which the MR start to decrease. Combined EXAFS and magnetic measurements allowed demonstrating that this anomalous behaviour signals a partial remixing of the elements causing the roughening of Co–Cu interfaces and worsening the MR. This phenomenon has been observed in all the Cox Cu1x samples investigated, despite the composition (x ¼ 5, 10, 15, 20), suggesting that it must derive from peculiar features in the phase diagram and/or in the thermodynamics of Co and Cu solution (as supposed in [16]) that must favour the a partial re-mixing of the immiscible elements. In situ experiments directly probe the effect of temperature during heating allowing to get deeper insights on this phenomenon. TR-XRD patterns, according to previous high-resolution XRD experiments [10–12], show two fcc phases (Fig. 1(a)). The refined structures definitively show the phase separation in a Cu-rich phase, corresponding to the most intense diffraction lines, and a Co-rich one, having weaker peaks (Fig. 1(b)). Raising the annealing temperature provokes the lattice expansion in both the phases (Fig. 2(a)) but the effect is definitively different in Co- and Cu-rich phases. The lattice parameter of the Co-rich phase (aCo ) increases linearly heating the sample between 300 and 700 °C while the expansion of Cu-rich phase (aCu ) depart from the linear trend heating above 600–650 °C. Linear expansion coefficients for Corich (kCo ) and Cu-rich (kCu ) phases have been estimated by linear fitting ai ðT Þ curves. The kCu ¼ 1:69 105 K1 , obtained from the aCu ðT Þ data (T < 600 °C), is the same of pure Cu-fcc (aCo ¼ 1:7 105 K1 [17]) and independent from heating speed (VT ). Above 600 °C the slope of aCu ðT Þ suddenly increases pointing out the activated segregation process in the Cu-rich phase that raises the lattice parameter toward the value expected in pure Cu. This effect is larger for the slower heating ramps. This finding comes in agreement with ex situ measurements showing that the Cu-rich phase does not change appreciably until annealing above 550–600 °C [10–12]. The Co-rich phase thermal expansion follows a linear behaviour in the whole temperature range
218
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
3.66
(a)
a (Å)
a
Cu
3.65 -5
-1
k = 1.69e K 3.64
a + 0.05 Co
-5
k = 1.36e K
3.63
∆a (Å) 0.07
-1
(b)
2
slower heating provokes the larger departure of Da from a linear trend signalling the larger segregation effect. Additional information on the structural evolution under annealing can be derived looking at the diffraction line profile and diffraction intensity (Fig. 3). The gaussian contribution (G) to the profile function reflects the distribution of lattice parameters resulting from stress/strain effects as well as from composition spread. The lorentzian (L) contribution, on the contrary, is mainly related to the crystallite size.
2.7
(x 10 deg)
4
500
600
700
o
T ( C)
2.7 32
4
28
6.7
(x 10 deg)
(b)
2
-2
400
2.7
G
Co
300 4 200 6.7
100
(c)
2
1.8
2.7 Co
with kCo ¼ 1:36 105 K1 , lower than that for pure Co-fcc (aCo ¼ 1:4 105 K1 [18]) and almost independent on the heating ratios. This effect is due to the competition of two effects: the thermal expansion, raising the lattice parameter, and the expulsion of Cu foreign atoms that according to the VegardÕs law reduces the aCo ðT Þ toward the values expected in pure Co-fcc. The evolution of aCo as a function of annealing temperature does not show the anomalies reported in ex situ experiments. This is probably related to the different annealing kinetics and deserves further investigation. The differences DaðT Þ ¼ aCu aCo (Fig. 2(b)) have all the same slope in the T < 600 °C region, suggesting that the heating kinetics affects only weakly the low temperature segregation process. On the contrary in the high temperature region the
36
500
H
Fig. 2. The lattice parameters for Cu-rich (aCu ) and Co-rich (aCo ) phases (VT ¼ 2 °C/min) as a function of annealing temperature are shown shifted for clarity (panel (a)). Full lines are linear fit to the data (see text). Panel (b) reports the differences Da ¼ aCu aCo as a function of heating speed (VT ). The numbers represent the VT for each data set. DaðT Þ depict a definitive departure from a linear trend heating above 600 °C, such an effect is more pronounced for the slower heating curves. Straight lines are guide for the eyes.
(a) 2
Cu
400
-2
6.7
40
L
0.06
1.4 4 1 6.7 400
500
600
700
o
T( C) Fig. 3. Results of diffraction profile refinement: the lorentzian contribution to the Cu-rich line shape (panel (a)) and the gaussian contribution to the Co-rich diffraction profile (panel (b)) are shown as a function of temperature for different VT (VT are reported for each data set). Panel (c) reports the histogram (HCo ) scale factor for Co-rich phase (see text).
C. Meneghini et al. / Nucl. Instr. and Meth. in Phys. Res. B 200 (2003) 215–219
Sharper diffraction lines characterize the Curich contribution to the diffractograms. The gaussian term in the Cu-rich phase peaks profile is about 0.15–0.2°, roughly constant as a function of T , and largely independent from VT . On the contrary the lorentzian contribution (LCu , Fig. 2(a)) depicts a clear temperature dependence: LCu ðT Þ decreases as a function of T in all the four kinetics, pointing out the progressive growth of crystallites. However, all the LCu ðT Þ curves depict an evident change of slope on heating the sample across 450– 500 °C. This effect appears independent on VT . Broader and weaker diffraction lines characterize the contribution of the Co-rich phase which profile appears dominated by the gaussian contribution (GCo ), while the lorentzian term is negligible. GCo ðT Þ is more than one order of magnitude larger than in Cu-rich phase and, interestingly, depicts an evident maximum as a function of T around 500 °C, which amplitude and position is largely independent on VT . This finding suggests a larger spread of Co-rich phase lattice parameters heating the sample between 450–500 °C signifying the activation of a remixing process in the Co-rich phase. The histogram scale factor (Hi ) of GSAS accounts for the relative amount of each phase within the sample. The relative weight of Cu-rich phase, HCu , is roughly independent from T . On the contrary the Co-rich scale factor (HCo ) depicts evident temperature dependence (Fig. 3(c)) with a clear minimum around 500 °C that seem unaffected by changing VT . This finding comes in agreement with the previous hypothesis of a Co– Cu remixing process activated on heating the sample around 500 °C that would cause the partial dissolution of Co-rich phase. TR-XRD results definitively prove that Co- and Cu-rich phases evolve differently as a function of annealing temperature. The thermally activated segregation process in the Co-rich phase depicts an anomalous behaviour that is largely independent on the heating kinetics.
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Acknowledgements We acknowledge the excellent technical support of F. Campolungo, V. Sciarra, V. Tullio (INFNLNF) and F. DÕAnca (INFM-OGG Grenoble). Gilda Beamline is financed by the Italian institutions CNR, INFM and INFN. This work has been partially supported by the Spanish CICyT under project no. MAT99-0667.
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