- Email: [email protected]
Development and characterization of Fe70Pd30 ferromagnetic shape memory splats
Intermetallics 18 (2010) 877–882
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Development and characterization of Fe70Pd30 ferromagnetic shape memory splats Iris Kock a, *, Sven Hamann b, **, Hayo Brunken b, Tobias Edler a, S.G. Mayr a, c, Alfred Ludwig b a
¨t Go ¨ttingen, Friedrich-Hund-Platz 1, 37077 Go ¨ttingen, Germany I. Physikalisches Institut, Georg-August-Universita Institute for Materials, Ruhr-University Bochum, 44780 Bochum, Germany c ¨chenmodifizierung e.V., Translationszentrum fu ¨t fu ¨t Leipzig, 04318 ¨ r Oberfla ¨ r regenerative Medizin und Fakulta ¨ r Physik und Geowissenschaften der Universita Leibniz-Institut fu Leipzig, Germany b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 18 March 2009 Received in revised form 23 October 2009 Accepted 12 December 2009 Available online 12 January 2010
Freestanding Fe70Pd30 foils with a thickness of about 60 mm were fabricated using the splat-quenching technique. A shift of the martensitic transformation temperatures as a function of different annealing treatments (600 C, 700 C, 800 C, 900 C, 1000 C for 15 min) was observed by temperature-dependent X-ray diffraction (XRD), resistance and magnetization measurements. The sample annealed at 800 C showed the highest degree of crystallinity for the (200) fcc austenite peak and no secondary phases. Samples annealed below 800 C kept austenite remainders even at 25 C. The transformation temperatures, determined by all three-measurement methods, showed an increase with increasing annealing temperature. Ó 2009 Elsevier Ltd. All rights reserved.
Keywords: B. Martensitic transformation C. Rapid solidification processing C. Heat treatment B. Magnetic properties B. Electrical resistance
1. Introduction Ferromagnetic shape memory alloys (FSMAs) have attracted tremendous scientific interest during the past years – primarily due to their high potential for actuators in engineering and applied sciences. The magnetic shape memory effect was first observed in Ni2MnGa bulk samples [1] and subsequently other compounds including Fe70Pd30. In Fe–Pd the actuation mechanism is based on the growth of martensite variants aligned by an external magnetic field in expense of other variants, which leads to a magnetic field induced strain (MFIS) up to 3% [2]. Application in the form of functional thin foils offers a promising route for microscale actuation. In contrast to conventional SMAs the speed of actuation is controlled by the switching of the magnetic field and the rearrangement of variants and is not limited by heat conduction. This leads to higher switching frequencies and thus to an expanded field of application. Generally there are some advantages of the Fe–Pd system in comparison to Ni–Mn–Ga. Fe70Pd30 has a higher saturation magnetization, a higher Curie temperature (490 C), better ductility and higher blocking stress [3]. Therefore Fe–Pd is much more applicable for sensor and actuator systems. From the Fe–Pd
phase diagram is known that the transforming fcc (Fe70Pd30) austenite phase occurs at 800 C or higher [4]. Therefore a rapid cooling after annealing is necessary to keep this metastable fcc austenite phase, needed for the MFIS effect, at room temperature. Slow cooling from the annealing temperature leads to decomposition into the stable, ordered L10 phase (Fe50Pd50) and the stable, iron rich, bcc Fe(Pd) solid solution of the equillibrium state [3]. FSMA application requires freestanding materials with thicknesses ranging from a few mm to several 100 mm [5]. Films produced by sputtering [6–8], thermal evaporation [9] and pulsed laser deposition method [10] or thinned single crystals [11] show promising results, but can barely reach these thicknesses due to the low deposition rates and the time-consuming fabrication process. Therefore appropriate methods are needed to produce FSMA materials in the desired design. In this work Fe70Pd30 splats with intermediate thickness were fabricated and investigated.
2. Experimental details 2.1. Fabrication of the Fe70Pd30 splats
* Corresponding author. Tel.: þ49 551 39 5553; fax: þ49 551 39 2328. ** Corresponding author. E-mail addresses: [email protected] (I. Kock), [email protected] (S. Hamann). 0966-9795/$ – see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2009.12.019
An Fe–Pd ingot of a nominal composition Fe70Pd30 was prepared by arc melting in an Ar atmosphere. To ensure homogenization, the
878
I. Kock et al. / Intermetallics 18 (2010) 877–882
ingot was turned over and melted six times. Subsequently, the ingot was divided into samples of about 180 mg each. These were remelted in the arc melter to form small spheres. These samples were prepared for splat-quenching [12,13]. The chamber of the splat quencher was purged six times with Ar (purity 99,998%) to avoid contamination and oxidation of the material. The samples were melted inductively in an Ar atmosphere of 600 mbar and subsequently splatted between two Cu pistons, leading to rapid solidification. The resultant splats have a thickness of about 60 mm. A typical splat is depicted in Fig. 1. It has a rough surface and shows structural differences between inner and outer parts. After polishing the splat, a homogeneous surface was obtained. The splats have a strong columnar structure as shown in Fig. 2 by optical microscopy on a cross section. In the lower part of Fig. 2 a splatlamella, fabricated by focussed ion beam (FIB) and investigated by transmission electron microscope (TEM), is presented. A clear columnar structure and martensitic twins can be observed in this picture. Due to the high quenching rate, the material crystallizes in the desired fcc austenite phase, which can transform to the martensite fct phase upon cooling. However, for the same reason, there is a great amount of stress and defects incorporated in the sample, which can pin the martensitic phase boundaries. Therefore, the samples were annealed at various temperatures (Tan) to heal out these defects and thus improve the transformation behavior. For this purpose, the samples were sealed into quartz tubes, which were purged with Ar several times, and then filled with Ar to a pressure of 600 mbar. After annealing at different Tan (600 C, 700 C, 800 C, 900 C and 1000 C) for 15 min the quartz tubes with the samples were dropped into cold water. A Bu¨hler diamond compound with a grain size of 1 mm was used to polish the samples after annealing. Especially for the temperature-dependent resistance measurements a smooth surface is important to avoid measurement errors. After polishing and cleaning, a depth profile (50 nm) was measured using an HR-Auger FEG-SEM (PHI-AES 690) in order to check that there was no diamond contamination in the samples. For visual investigations of cross sections, trenches and lamellae from various splats were fabricated using a FEI Quanta 3D FIB system. The visual observations of the lamellae were performed using a FEI Tecnai F20-G2 TEM. 2.2. Characterization of the Fe70Pd30 splats The composition of each splat was determined by energy dispersive X-ray spectroscopy (EDX) to exclude possible
Fig. 1. Photograph of a splat (as-splatted) showing structural differences between inner and outer parts.
Fig. 2. Sectional view of a splat by optical microscopy (upper part, as-splatted) and TEM (lower part, annealed at 1000 C). A strong columnar structure is visible.
composition changes during the fabrication process: The composition of ten points of each sample was measured using a LEO Supra 55 SEM with an Oxford Instruments EDX detector. To minimize the total measurement error (z0.5 at.%) an Fe70Pd30 bulk standard was used for calibration. The results of the EDX measurements showed a deviation smaller than the measurement error from the nominal composition. Automated temperature-dependent resistance measurements (R(T)) were performed using a custom-designed test stand [14]. The samples were positioned on a variable temperature vacuum stage and measured with a four point spring-probe resistance measurement head. The temperature of the samples was varied with a rate of 5 C/min and cycled two times between 40 C and 120 C. For the R(T) measurement a source current of 100 mA was applied using a Multi Source Meter Keithley 2000. XRD measurements were performed using a Bruker-AXS (with General Area Detector Diffraction System, Cu Ka radiation, spot size 0.5 mm, integration time 300 s, 2Q area range from 35 to 65 ) to determine the crystal structure of the samples at room temperature. Before measurement, an Al2O3 standard was used to calibrate the system and minimize errors to values < 0.05 . Temperaturedependent XRD measurements in the range from 25 C to 40 C were used to verify the martensitic transformation previously determined by the R(T) measurement. The temperature-dependent measurements from 25 C up to 20 C were conducted using an evacuated (p < 3$103 mbar) Be dome to prevent ice from forming on the samples. Additional measurements at 20 C were performed in air without a dome to identify diffraction peaks originating from Be. The temperature during the XRD measurement was determined using a NiCr–Ni thermocouple, located between the samples and the Peltier element. The accuracy of the temperature measurement was better than 0.5 C. Magnetization of the samples was measured with a Quantum Design Squid magnetometer. The samples were embedded in epoxy to avoid movement during the measurement. The magnetic moment of the epoxy was confirmed to be more than three orders of magnitude lower than that of the samples, so that it could be neglected. Temperature-dependent measurements of the magnetic moment were performed at a field of 300 Oe. First, the temperature was cycled from 27 C to 73 C and back in steps of 5 C. To determine the transformation temperatures more exactly, a second measurement was performed between 27 C and 2 C with a step width of 1 C.
I. Kock et al. / Intermetallics 18 (2010) 877–882
3. Results 3.1. Temperature-dependent resistance measurements R(T) measurements were used to identify martensitic transformations of the splats and their transformation temperatures. In Fig. 3 the normalized R(T) measurements (arbitrary offset for clarity) of Fe70Pd30 splats, as-splatted and annealed at different temperatures, are shown. The samples as-splatted and annealed at 600 C and 700 C show only a slight non-linearity (hardly visible here due to scaling) in the R(T) curve. The transformation temperatures were determined from the non-linearity using the tangential method (light grey lines). The other splats show a much clearer non-linearity in the R(T) curve. In contrast to R(T) measurements of conventional shape memory alloys (SMAs) [15] and extremely rapidly quenched Fe70Pd30 thin films using microhotplates [16], these samples show only weak changes in the R(T) curve with no determinable hysteresis width. Therefore it is not possible to distinguish between martensite start (Ms) and austenite finish (Af) or martensite finish (Mf) and austenite start (As) temperature. Thus only one temperature is determined for the start (Af and Ms temperature) and the end (As and Mf temperature) of the martensitic transformation. An increase of the transformation temperatures occurs from the as-splatted to annealed samples with increasing Tan as shown in Fig. 6. 3.2. Temperature-dependent XRD measurements Temperature-dependent XRD spectra of the splats are presented in Fig. 4. At 25 C, the splat annealed at 800 C shows the (111) and the (200) fct peaks of the transforming Fe70Pd30 phase. With increasing temperature, the (200) peak of the fct phase shifts to higher 2Q angles and the (200) fcc peak appears and increases in intensity. The transformation is reversible as confirmed by heating and cooling the splats several times. This shift of the fct peak and the development of the (200) fcc peak is observed for all six splats (even the one annealed at 700 C shows a slight shift and development of the (200) fcc peak). Some splats (as-splatted, annealed at 600 C and at 700 C) show next to the (111) fct/fcc peak one broad peak at the (200) fct and (200) fcc peak position at low temperatures. With increasing temperature this broadened peak shifts more and more to the (200) fcc peak position while the austenitic state appears. As-
879
splatted splats and those annealed at 600 C and 700 C show additional phases due to a decomposition of the austenite Fe70Pd30 phase into Fe50Pd50 and Fe. A (110) peak of the a-Fe phase (splats annealed at 600 C and 700 C) and a (111) Fe50Pd50 peak (assplatted and annealed at 600 C and 700 C) occurs for various splats. The ones annealed at 600 C and 700 C show a weak (200) fcc peak. This can be correlated to the decreased amount of the Fe70Pd30 phase due to decomposition during annealing below 800 C. All splats annealed at 800 C and higher start to develop a well defined (200) fcc peak when they are heated over 6 C. It is noted that the position of the fct peaks are slightly different in comparison to literature values [7]. This can be explained by the temperature dependence of the lattice parameters for Fe70Pd30 reported by [3]. Thus the splats have a martensitic state at 25 C with a small amount of austenite remainders. At temperatures above 20 C, the splats annealed at temperatures >700 C transform into the austenitic state without any martensite remainders. Splats annealed at 700 C and below have an intermediate state, where both the fct martensitic and the fcc austenitic structure occur even at low temperatures (25 C). From the occurence of the peaks during heating, the transformation temperatures (only Ms/Af and Mf/As), were estimated with an error <5 C. 3.3. Temperature-dependent magnetization measurements Fig. 5 shows the temperature-dependent magnetization measurements M(T) (at a constant magnetic field of 300 Oe) for different splats. At temperatures >18 C, where the splats are austenitic, the magnetization does not vary significantly with temperature. At the beginning of the phase transition, the magnetization starts to decrease rapidly. Measurements between 27 C and 73 C (not shown here) indicate, that the decrease of magnetization continues below 0 C. For splats annealed at 600 C, 800 C, 900 C and 1000 C, a hysteresis was observed in the M(T) cycles, and the martensite start temperature Ms (determined by the tangential method), is lower than the austenite finish temperature Af. The hysteresis extends to lower temperatures in the cycles that have been continued to lower temperatures. As there is no distinct change in magnetization at the end of the phase transformation, Mf and As could not be determined from the magnetization curves. For splats annealed at 900 C and 1000 C, the decrease in magnetization and the hysteretic behavior are most distinct. The sample annealed at 700 C does not show a clear transformation, while the one annealed at 600 C transforms in two steps (not shown here). In Fig. 6 the transformation temperatures (Af/Ms and As/Mf) determined by R(T), XRD and M(T) measurements, are presented as a function of Tan. The As/Mf curve shows a steeper increase than the Af/Ms curve. The guide to the eye fitted to the experimental data accentuates the Arrhenius like increase of the Af/Ms temperatures with increasing Tan. The reason for this behaviour will be discussed in the following chapter. 4. Discussion
Fig. 3. Normalized resistance of different splats as function of temperature. In the inset, black dots indicate the transformation temperatures determined by the tangential method of the splat annealed at 1000 C.
The crystal structure, resistance and magnetization of the splats show distinct changes upon phase transformation. The transformation temperatures – determined with the different methods – match within error margins and increase with increasing Tan. Only the sample annealed at 700 C shows significant deviations in all measurements, as no pronounced transformation can be confirmed. The XRD diffractogramms show strong Fe precipitates and the EDX analysis of this sample also shows inhomogeneities in the Pd content which is assumed to be caused by some irregularity during the splatting or annealing process.
880
I. Kock et al. / Intermetallics 18 (2010) 877–882
Fig. 4. Temperature-dependent XRD spectra of as-splatted and annealed samples (offset for clarity).
A change in band structure, and especially in the density of electronic d-states near the Fermi energy as calculated by Opahle et al. [17], would explain both the change in conductivity as well as the decrease in magnetic moment. An alternative explanation for the decrease of magnetization and resistance is the increase of magnetic anisotropy energy with decreasing temperature [18]. In a low magnetic field of 300 Oe, the individual magnetic moments rotate out of the field direction and into the direction of the magnetic easy axis when the magnetic anisotropy energy is sufficiently high. This leads to demagnetization of the sample and a change in magnetic domain structure, which influences the electronic scattering at domain walls [19]. To understand the increase of transformation temperature with increasing Tan, (Fig. 6) a closer examination of the parameters influencing the transformation behavior is necessary. In previous works by Cui et al., Guenin et al. and Kato et al. it was found that the transformation temperature is mainly influenced by the Pd content [3], the defect density [20] and the stress inside the sample [21].
These quantities depend for example on homogeneity, quenching rate, grain size and shape or precipitates, which also influence each other. Moreover, it is necessary to distinguish between factors influencing the temperature T0 at which austenite and martensite are in thermodynamic equilibrium and other factors influencing the actual transformation temperature via the necessary undercooling DT ¼ T0 – Ms that acts as a driving force. In the present experiment, great care was taken to ensure adequate homogeneity of the Pd content in the splats. However, the possibility of slight deviations in Pd content (<0.3 at.%) can not be ruled out by EDX. As the thermodynamic equilibrium temperature T0 strongly depends on Pd content [3] (z20 C/at.%), this could lead to a change in transformation temperature of several degrees, independent of any influences from the annealing process. The defect density and distribution in the samples is not identical, either. Even within one sample there are great differences between inner and outer parts of the splat due to the splatting of an initially spherical droplet (see Fig. 1). A high defect density provides
I. Kock et al. / Intermetallics 18 (2010) 877–882
Fig. 5. Temperature-dependent normalized magnetization of the splats measured at a magnetic field of 300 Oe, (curves are offset for clarity).
many nucleation sites for the martensite phase [20], leading to high transformation temperatures. During the annealing process, some defects can heal out, which should lead to a decrease in Ms with increasing Tan. On the other hand, defects can hinder the movements of the phase boundary through the material. Especially if precipitates have formed during the annealing process at Tan < 800 C, the phase boundary is pinned, the necessary undercooling is increased and the material transforms at lower temperatures [22]. As the increase of Mf with increasing Tan is much more distinct than that of Ms and there is hardly any hysteresis between heating and cooling, the amount of martensite phase is assumed to increase by movement of the phase boundary rather than by nucleation of new martensite embryos. The stress state inside the sample can also influence the phase transformation. The ocurrence of stress induced martensite is well known [21]. However, as the splats are not bound to a rigid substrate, most of the stress incorporated during the splatting process can relax via bending or expansion of the sample. Still, in a rapidly solidified polycrystalline sample there are always contributions to the intrinsic stress by lattice distortions at defects. A quantification of the intrinsic microstress from the XRD diffractograms is difficult, as the peak width w is not only influenced by the
881
s [24] (E is microstress s inside the sample according to w ¼ E$4tan q the elastic modulus of the sample), but also by the grain size according to the Scherrer-formula. The dependence of transformation temperature on grain size was discussed in several works. Seki et al. [23] and Kang et al. [25] found a strong increase of the transformation temperature with grain size. As the grain size R is expected to increase with Tan according to R ¼ R0 þ a$t$exp(Q/kBTan) (with the initial grain size R0, the activation energy for grain boundary movement Q and a proportionality factor a) the transformation temperature should increase in the same way. The data of Fig. 6 supports this assumption, if it is considered that the annealing time is short (15 min). Deviations from the Arrhenius like behavior can be explained by variations in the initial grain size due to the splatting and small deviations in the Pd content. An exact determination of the grain size by metallography was not possible. Therefore the grain structure was determined by SEM and TEM using cross sectional samples of the splats prepared by FIB. The expected increase of columnar grain width from 150–200 nm in the assplatted sample to 500–600 nm in the sample annealed at 1000 C was confirmed (see Supplementary data). As mentioned above, the martensite phase is assumed to increase by growth rather than by nucleation. In a coarse grained structure, there are less grain boundaries pinning the phase boundary, which leads to an increase of transformation temperature, especially of Mf, as the elastic energy stored during transformation is smaller [26]. From the discussion, it can be seen that the transformation temperature can be strongly influenced by slight fluctuations in the Pd content and initial grain size, while an increase in transformation temperature with increasing annealing temperature due to the increase in grain size and changes in stress state and defect density during the annealing process is still visible.
5. Conclusions Splat-quenching of the ferromagnetic shape memory alloy Fe70Pd30 was found to be an excellent method to obtain thin foils in the fcc austenite phase, which show a martensitic transformation upon cooling. The temperature-dependent behavior of crystal structure, resistance and magnetization of these splats was examined. All three properties show distinct changes upon phase transformation and thus are useful to determine the transformation temperature and behavior. Using these results, the transformation temperature was correlated to the annealing temperature and the transformation behavior was optimized by an appropriate annealing treatment at high temperatures. The optimum annealing temperature was found to be around 800 C as the fct/fcc peaks are well defined and there are no hints for decomposition and precipitates. Acknowledgments The authors gratefully acknowledge funding by the Deutsche Forschungsgemeinschaft (DFG) in the priority program (SPP) 1239 TP C4 and C8, and would like to thank Prof. K. Samwer for support, A. Kru¨ger for preparation of the splats and D. Ko¨nig for TEM investigations. Appendix. Supplementary data
Fig. 6. Transformation temperatures of the splats determined by different measurement methods as function of Tan.
The supplementary data associated with this article can be found in the on-line version at doi:10.1016/j.intermet.2009.12.019.
882
I. Kock et al. / Intermetallics 18 (2010) 877–882
References [1] Ullakko K, Huang JK, Kantner CC, O’Handley RC, Kokorin VV. Appl Phys Lett 1996;69:1966. [2] Kakeshita T, Fukuda T. Mater Sci Forum 2002;394:531. [3] Cui J, Shield TW, James RD. Acta Mater 2004;52:35–47. [4] Massalski B, Okamoto H, Subramanian PR, Kacprzak L. ASM international’s binary alloy phase diagrams. 2nd ed.; 1996. p. 35–47. [5] Spasova B, Ruffert C, Wurz M, Gatzen H. Phys Stat Sol A 2008;205:2307. [6] Vokoun D, Shih J, Chin T, Hu C. J Magn Magn Mater 2004;281:105. [7] Inoue S, Inoue K, Koterazawa K, Mizuuchi K. Mater Sci Eng 2003;A339:29. [8] Sugimura Y, Cohen-Karni I, McCluskey P, Vlassak JJ. J Mater Res 2005;20:2279. [9] Kock I, Edler T, Mayr SG. J Appl Phys 2008;103:046108. [10] Edler T, Buschbeck J, Mickel C, Fa¨hler S, Mayr SG. New J Phys 2008;10:063007. [11] Heczko 0, Soroka A, Hannula S-P. Appl Phys Lett 2008;93:022503. [12] Duwez P, Willens RH, Klement W. J Appl Phys 1960;31:1136. [13] Jones H. Rep Prog Phys 1973;36:1425.
[14] Thienhaus S, Zamponi C, Rumpf H, Hattrick-Simpers J, Takeuchiand I, Ludwig A. Materials Research Society Symposium Proceeding 894, Warrendale Pa 0894-LL06-06.1; 2006. [15] Zarnetta R, Savan A, Thienhaus S, Ludwig A. Appl Surf Sci 2007;254:743–8. [16] Hamann S, Ehmann M, Thienhaus S, Savan A, Ludwig A. Sensor Actuator: A Phys 2008;147(2):576–82. [17] Opahle I, Kopernik K, Nitzsche U, Richter M. Appl Phys Lett 2009;94:072508. [18] Kakeshita T, Fukuda T. Sci Tech Adv Mater 2006;7:350. [19] Viret M, Samson Y, Warin P, Marty A, Ott F, Søndergård E, et al. Phys Rev Lett 2000;85:3962. [20] Guenin G, Gobin PF. Metall Mater Trans 1982;13:1127–34. [21] Kato H, Liang Y, Taya M. Scripta Mater 2002;46:471–5. [22] Dondl P, Bhattacharya K. Proc Appl Math Mech 2007;7:1151207. [23] Seki K, Kura H, Sato T, Taniyama T. J Appl Phys 2008;103:063910. [24] Deryavko I, Lanin A, Taubini M. Powder Metall Met Ceram 1978;17:137–40. [25] Kang S, Lee Y, Lim Y, Nam J, Nam T, Kim Y. Scripta Mater 2008;59:1186–9. [26] Wollants P, Roos J, Delaely L. Progr Mater Sci 1993;37:227–88.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Development and characterization of Fe70Pd30 ferromagnetic shape memory splats Iris Kock a, *, Sven Hamann b, **, Hayo Brunken b, Tobias Edler a, S.G. Mayr a, c, Alfred Ludwig b a
¨t Go ¨ttingen, Friedrich-Hund-Platz 1, 37077 Go ¨ttingen, Germany I. Physikalisches Institut, Georg-August-Universita Institute for Materials, Ruhr-University Bochum, 44780 Bochum, Germany c ¨chenmodifizierung e.V., Translationszentrum fu ¨t fu ¨t Leipzig, 04318 ¨ r Oberfla ¨ r regenerative Medizin und Fakulta ¨ r Physik und Geowissenschaften der Universita Leibniz-Institut fu Leipzig, Germany b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 18 March 2009 Received in revised form 23 October 2009 Accepted 12 December 2009 Available online 12 January 2010
Freestanding Fe70Pd30 foils with a thickness of about 60 mm were fabricated using the splat-quenching technique. A shift of the martensitic transformation temperatures as a function of different annealing treatments (600 C, 700 C, 800 C, 900 C, 1000 C for 15 min) was observed by temperature-dependent X-ray diffraction (XRD), resistance and magnetization measurements. The sample annealed at 800 C showed the highest degree of crystallinity for the (200) fcc austenite peak and no secondary phases. Samples annealed below 800 C kept austenite remainders even at 25 C. The transformation temperatures, determined by all three-measurement methods, showed an increase with increasing annealing temperature. Ó 2009 Elsevier Ltd. All rights reserved.
Keywords: B. Martensitic transformation C. Rapid solidification processing C. Heat treatment B. Magnetic properties B. Electrical resistance
1. Introduction Ferromagnetic shape memory alloys (FSMAs) have attracted tremendous scientific interest during the past years – primarily due to their high potential for actuators in engineering and applied sciences. The magnetic shape memory effect was first observed in Ni2MnGa bulk samples [1] and subsequently other compounds including Fe70Pd30. In Fe–Pd the actuation mechanism is based on the growth of martensite variants aligned by an external magnetic field in expense of other variants, which leads to a magnetic field induced strain (MFIS) up to 3% [2]. Application in the form of functional thin foils offers a promising route for microscale actuation. In contrast to conventional SMAs the speed of actuation is controlled by the switching of the magnetic field and the rearrangement of variants and is not limited by heat conduction. This leads to higher switching frequencies and thus to an expanded field of application. Generally there are some advantages of the Fe–Pd system in comparison to Ni–Mn–Ga. Fe70Pd30 has a higher saturation magnetization, a higher Curie temperature (490 C), better ductility and higher blocking stress [3]. Therefore Fe–Pd is much more applicable for sensor and actuator systems. From the Fe–Pd
phase diagram is known that the transforming fcc (Fe70Pd30) austenite phase occurs at 800 C or higher [4]. Therefore a rapid cooling after annealing is necessary to keep this metastable fcc austenite phase, needed for the MFIS effect, at room temperature. Slow cooling from the annealing temperature leads to decomposition into the stable, ordered L10 phase (Fe50Pd50) and the stable, iron rich, bcc Fe(Pd) solid solution of the equillibrium state [3]. FSMA application requires freestanding materials with thicknesses ranging from a few mm to several 100 mm [5]. Films produced by sputtering [6–8], thermal evaporation [9] and pulsed laser deposition method [10] or thinned single crystals [11] show promising results, but can barely reach these thicknesses due to the low deposition rates and the time-consuming fabrication process. Therefore appropriate methods are needed to produce FSMA materials in the desired design. In this work Fe70Pd30 splats with intermediate thickness were fabricated and investigated.
2. Experimental details 2.1. Fabrication of the Fe70Pd30 splats
* Corresponding author. Tel.: þ49 551 39 5553; fax: þ49 551 39 2328. ** Corresponding author. E-mail addresses: [email protected] (I. Kock), [email protected] (S. Hamann). 0966-9795/$ – see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2009.12.019
An Fe–Pd ingot of a nominal composition Fe70Pd30 was prepared by arc melting in an Ar atmosphere. To ensure homogenization, the
878
I. Kock et al. / Intermetallics 18 (2010) 877–882
ingot was turned over and melted six times. Subsequently, the ingot was divided into samples of about 180 mg each. These were remelted in the arc melter to form small spheres. These samples were prepared for splat-quenching [12,13]. The chamber of the splat quencher was purged six times with Ar (purity 99,998%) to avoid contamination and oxidation of the material. The samples were melted inductively in an Ar atmosphere of 600 mbar and subsequently splatted between two Cu pistons, leading to rapid solidification. The resultant splats have a thickness of about 60 mm. A typical splat is depicted in Fig. 1. It has a rough surface and shows structural differences between inner and outer parts. After polishing the splat, a homogeneous surface was obtained. The splats have a strong columnar structure as shown in Fig. 2 by optical microscopy on a cross section. In the lower part of Fig. 2 a splatlamella, fabricated by focussed ion beam (FIB) and investigated by transmission electron microscope (TEM), is presented. A clear columnar structure and martensitic twins can be observed in this picture. Due to the high quenching rate, the material crystallizes in the desired fcc austenite phase, which can transform to the martensite fct phase upon cooling. However, for the same reason, there is a great amount of stress and defects incorporated in the sample, which can pin the martensitic phase boundaries. Therefore, the samples were annealed at various temperatures (Tan) to heal out these defects and thus improve the transformation behavior. For this purpose, the samples were sealed into quartz tubes, which were purged with Ar several times, and then filled with Ar to a pressure of 600 mbar. After annealing at different Tan (600 C, 700 C, 800 C, 900 C and 1000 C) for 15 min the quartz tubes with the samples were dropped into cold water. A Bu¨hler diamond compound with a grain size of 1 mm was used to polish the samples after annealing. Especially for the temperature-dependent resistance measurements a smooth surface is important to avoid measurement errors. After polishing and cleaning, a depth profile (50 nm) was measured using an HR-Auger FEG-SEM (PHI-AES 690) in order to check that there was no diamond contamination in the samples. For visual investigations of cross sections, trenches and lamellae from various splats were fabricated using a FEI Quanta 3D FIB system. The visual observations of the lamellae were performed using a FEI Tecnai F20-G2 TEM. 2.2. Characterization of the Fe70Pd30 splats The composition of each splat was determined by energy dispersive X-ray spectroscopy (EDX) to exclude possible
Fig. 1. Photograph of a splat (as-splatted) showing structural differences between inner and outer parts.
Fig. 2. Sectional view of a splat by optical microscopy (upper part, as-splatted) and TEM (lower part, annealed at 1000 C). A strong columnar structure is visible.
composition changes during the fabrication process: The composition of ten points of each sample was measured using a LEO Supra 55 SEM with an Oxford Instruments EDX detector. To minimize the total measurement error (z0.5 at.%) an Fe70Pd30 bulk standard was used for calibration. The results of the EDX measurements showed a deviation smaller than the measurement error from the nominal composition. Automated temperature-dependent resistance measurements (R(T)) were performed using a custom-designed test stand [14]. The samples were positioned on a variable temperature vacuum stage and measured with a four point spring-probe resistance measurement head. The temperature of the samples was varied with a rate of 5 C/min and cycled two times between 40 C and 120 C. For the R(T) measurement a source current of 100 mA was applied using a Multi Source Meter Keithley 2000. XRD measurements were performed using a Bruker-AXS (with General Area Detector Diffraction System, Cu Ka radiation, spot size 0.5 mm, integration time 300 s, 2Q area range from 35 to 65 ) to determine the crystal structure of the samples at room temperature. Before measurement, an Al2O3 standard was used to calibrate the system and minimize errors to values < 0.05 . Temperaturedependent XRD measurements in the range from 25 C to 40 C were used to verify the martensitic transformation previously determined by the R(T) measurement. The temperature-dependent measurements from 25 C up to 20 C were conducted using an evacuated (p < 3$103 mbar) Be dome to prevent ice from forming on the samples. Additional measurements at 20 C were performed in air without a dome to identify diffraction peaks originating from Be. The temperature during the XRD measurement was determined using a NiCr–Ni thermocouple, located between the samples and the Peltier element. The accuracy of the temperature measurement was better than 0.5 C. Magnetization of the samples was measured with a Quantum Design Squid magnetometer. The samples were embedded in epoxy to avoid movement during the measurement. The magnetic moment of the epoxy was confirmed to be more than three orders of magnitude lower than that of the samples, so that it could be neglected. Temperature-dependent measurements of the magnetic moment were performed at a field of 300 Oe. First, the temperature was cycled from 27 C to 73 C and back in steps of 5 C. To determine the transformation temperatures more exactly, a second measurement was performed between 27 C and 2 C with a step width of 1 C.
I. Kock et al. / Intermetallics 18 (2010) 877–882
3. Results 3.1. Temperature-dependent resistance measurements R(T) measurements were used to identify martensitic transformations of the splats and their transformation temperatures. In Fig. 3 the normalized R(T) measurements (arbitrary offset for clarity) of Fe70Pd30 splats, as-splatted and annealed at different temperatures, are shown. The samples as-splatted and annealed at 600 C and 700 C show only a slight non-linearity (hardly visible here due to scaling) in the R(T) curve. The transformation temperatures were determined from the non-linearity using the tangential method (light grey lines). The other splats show a much clearer non-linearity in the R(T) curve. In contrast to R(T) measurements of conventional shape memory alloys (SMAs) [15] and extremely rapidly quenched Fe70Pd30 thin films using microhotplates [16], these samples show only weak changes in the R(T) curve with no determinable hysteresis width. Therefore it is not possible to distinguish between martensite start (Ms) and austenite finish (Af) or martensite finish (Mf) and austenite start (As) temperature. Thus only one temperature is determined for the start (Af and Ms temperature) and the end (As and Mf temperature) of the martensitic transformation. An increase of the transformation temperatures occurs from the as-splatted to annealed samples with increasing Tan as shown in Fig. 6. 3.2. Temperature-dependent XRD measurements Temperature-dependent XRD spectra of the splats are presented in Fig. 4. At 25 C, the splat annealed at 800 C shows the (111) and the (200) fct peaks of the transforming Fe70Pd30 phase. With increasing temperature, the (200) peak of the fct phase shifts to higher 2Q angles and the (200) fcc peak appears and increases in intensity. The transformation is reversible as confirmed by heating and cooling the splats several times. This shift of the fct peak and the development of the (200) fcc peak is observed for all six splats (even the one annealed at 700 C shows a slight shift and development of the (200) fcc peak). Some splats (as-splatted, annealed at 600 C and at 700 C) show next to the (111) fct/fcc peak one broad peak at the (200) fct and (200) fcc peak position at low temperatures. With increasing temperature this broadened peak shifts more and more to the (200) fcc peak position while the austenitic state appears. As-
879
splatted splats and those annealed at 600 C and 700 C show additional phases due to a decomposition of the austenite Fe70Pd30 phase into Fe50Pd50 and Fe. A (110) peak of the a-Fe phase (splats annealed at 600 C and 700 C) and a (111) Fe50Pd50 peak (assplatted and annealed at 600 C and 700 C) occurs for various splats. The ones annealed at 600 C and 700 C show a weak (200) fcc peak. This can be correlated to the decreased amount of the Fe70Pd30 phase due to decomposition during annealing below 800 C. All splats annealed at 800 C and higher start to develop a well defined (200) fcc peak when they are heated over 6 C. It is noted that the position of the fct peaks are slightly different in comparison to literature values [7]. This can be explained by the temperature dependence of the lattice parameters for Fe70Pd30 reported by [3]. Thus the splats have a martensitic state at 25 C with a small amount of austenite remainders. At temperatures above 20 C, the splats annealed at temperatures >700 C transform into the austenitic state without any martensite remainders. Splats annealed at 700 C and below have an intermediate state, where both the fct martensitic and the fcc austenitic structure occur even at low temperatures (25 C). From the occurence of the peaks during heating, the transformation temperatures (only Ms/Af and Mf/As), were estimated with an error <5 C. 3.3. Temperature-dependent magnetization measurements Fig. 5 shows the temperature-dependent magnetization measurements M(T) (at a constant magnetic field of 300 Oe) for different splats. At temperatures >18 C, where the splats are austenitic, the magnetization does not vary significantly with temperature. At the beginning of the phase transition, the magnetization starts to decrease rapidly. Measurements between 27 C and 73 C (not shown here) indicate, that the decrease of magnetization continues below 0 C. For splats annealed at 600 C, 800 C, 900 C and 1000 C, a hysteresis was observed in the M(T) cycles, and the martensite start temperature Ms (determined by the tangential method), is lower than the austenite finish temperature Af. The hysteresis extends to lower temperatures in the cycles that have been continued to lower temperatures. As there is no distinct change in magnetization at the end of the phase transformation, Mf and As could not be determined from the magnetization curves. For splats annealed at 900 C and 1000 C, the decrease in magnetization and the hysteretic behavior are most distinct. The sample annealed at 700 C does not show a clear transformation, while the one annealed at 600 C transforms in two steps (not shown here). In Fig. 6 the transformation temperatures (Af/Ms and As/Mf) determined by R(T), XRD and M(T) measurements, are presented as a function of Tan. The As/Mf curve shows a steeper increase than the Af/Ms curve. The guide to the eye fitted to the experimental data accentuates the Arrhenius like increase of the Af/Ms temperatures with increasing Tan. The reason for this behaviour will be discussed in the following chapter. 4. Discussion
Fig. 3. Normalized resistance of different splats as function of temperature. In the inset, black dots indicate the transformation temperatures determined by the tangential method of the splat annealed at 1000 C.
The crystal structure, resistance and magnetization of the splats show distinct changes upon phase transformation. The transformation temperatures – determined with the different methods – match within error margins and increase with increasing Tan. Only the sample annealed at 700 C shows significant deviations in all measurements, as no pronounced transformation can be confirmed. The XRD diffractogramms show strong Fe precipitates and the EDX analysis of this sample also shows inhomogeneities in the Pd content which is assumed to be caused by some irregularity during the splatting or annealing process.
880
I. Kock et al. / Intermetallics 18 (2010) 877–882
Fig. 4. Temperature-dependent XRD spectra of as-splatted and annealed samples (offset for clarity).
A change in band structure, and especially in the density of electronic d-states near the Fermi energy as calculated by Opahle et al. [17], would explain both the change in conductivity as well as the decrease in magnetic moment. An alternative explanation for the decrease of magnetization and resistance is the increase of magnetic anisotropy energy with decreasing temperature [18]. In a low magnetic field of 300 Oe, the individual magnetic moments rotate out of the field direction and into the direction of the magnetic easy axis when the magnetic anisotropy energy is sufficiently high. This leads to demagnetization of the sample and a change in magnetic domain structure, which influences the electronic scattering at domain walls [19]. To understand the increase of transformation temperature with increasing Tan, (Fig. 6) a closer examination of the parameters influencing the transformation behavior is necessary. In previous works by Cui et al., Guenin et al. and Kato et al. it was found that the transformation temperature is mainly influenced by the Pd content [3], the defect density [20] and the stress inside the sample [21].
These quantities depend for example on homogeneity, quenching rate, grain size and shape or precipitates, which also influence each other. Moreover, it is necessary to distinguish between factors influencing the temperature T0 at which austenite and martensite are in thermodynamic equilibrium and other factors influencing the actual transformation temperature via the necessary undercooling DT ¼ T0 – Ms that acts as a driving force. In the present experiment, great care was taken to ensure adequate homogeneity of the Pd content in the splats. However, the possibility of slight deviations in Pd content (<0.3 at.%) can not be ruled out by EDX. As the thermodynamic equilibrium temperature T0 strongly depends on Pd content [3] (z20 C/at.%), this could lead to a change in transformation temperature of several degrees, independent of any influences from the annealing process. The defect density and distribution in the samples is not identical, either. Even within one sample there are great differences between inner and outer parts of the splat due to the splatting of an initially spherical droplet (see Fig. 1). A high defect density provides
I. Kock et al. / Intermetallics 18 (2010) 877–882
Fig. 5. Temperature-dependent normalized magnetization of the splats measured at a magnetic field of 300 Oe, (curves are offset for clarity).
many nucleation sites for the martensite phase [20], leading to high transformation temperatures. During the annealing process, some defects can heal out, which should lead to a decrease in Ms with increasing Tan. On the other hand, defects can hinder the movements of the phase boundary through the material. Especially if precipitates have formed during the annealing process at Tan < 800 C, the phase boundary is pinned, the necessary undercooling is increased and the material transforms at lower temperatures [22]. As the increase of Mf with increasing Tan is much more distinct than that of Ms and there is hardly any hysteresis between heating and cooling, the amount of martensite phase is assumed to increase by movement of the phase boundary rather than by nucleation of new martensite embryos. The stress state inside the sample can also influence the phase transformation. The ocurrence of stress induced martensite is well known [21]. However, as the splats are not bound to a rigid substrate, most of the stress incorporated during the splatting process can relax via bending or expansion of the sample. Still, in a rapidly solidified polycrystalline sample there are always contributions to the intrinsic stress by lattice distortions at defects. A quantification of the intrinsic microstress from the XRD diffractograms is difficult, as the peak width w is not only influenced by the
881
s [24] (E is microstress s inside the sample according to w ¼ E$4tan q the elastic modulus of the sample), but also by the grain size according to the Scherrer-formula. The dependence of transformation temperature on grain size was discussed in several works. Seki et al. [23] and Kang et al. [25] found a strong increase of the transformation temperature with grain size. As the grain size R is expected to increase with Tan according to R ¼ R0 þ a$t$exp(Q/kBTan) (with the initial grain size R0, the activation energy for grain boundary movement Q and a proportionality factor a) the transformation temperature should increase in the same way. The data of Fig. 6 supports this assumption, if it is considered that the annealing time is short (15 min). Deviations from the Arrhenius like behavior can be explained by variations in the initial grain size due to the splatting and small deviations in the Pd content. An exact determination of the grain size by metallography was not possible. Therefore the grain structure was determined by SEM and TEM using cross sectional samples of the splats prepared by FIB. The expected increase of columnar grain width from 150–200 nm in the assplatted sample to 500–600 nm in the sample annealed at 1000 C was confirmed (see Supplementary data). As mentioned above, the martensite phase is assumed to increase by growth rather than by nucleation. In a coarse grained structure, there are less grain boundaries pinning the phase boundary, which leads to an increase of transformation temperature, especially of Mf, as the elastic energy stored during transformation is smaller [26]. From the discussion, it can be seen that the transformation temperature can be strongly influenced by slight fluctuations in the Pd content and initial grain size, while an increase in transformation temperature with increasing annealing temperature due to the increase in grain size and changes in stress state and defect density during the annealing process is still visible.
5. Conclusions Splat-quenching of the ferromagnetic shape memory alloy Fe70Pd30 was found to be an excellent method to obtain thin foils in the fcc austenite phase, which show a martensitic transformation upon cooling. The temperature-dependent behavior of crystal structure, resistance and magnetization of these splats was examined. All three properties show distinct changes upon phase transformation and thus are useful to determine the transformation temperature and behavior. Using these results, the transformation temperature was correlated to the annealing temperature and the transformation behavior was optimized by an appropriate annealing treatment at high temperatures. The optimum annealing temperature was found to be around 800 C as the fct/fcc peaks are well defined and there are no hints for decomposition and precipitates. Acknowledgments The authors gratefully acknowledge funding by the Deutsche Forschungsgemeinschaft (DFG) in the priority program (SPP) 1239 TP C4 and C8, and would like to thank Prof. K. Samwer for support, A. Kru¨ger for preparation of the splats and D. Ko¨nig for TEM investigations. Appendix. Supplementary data
Fig. 6. Transformation temperatures of the splats determined by different measurement methods as function of Tan.
The supplementary data associated with this article can be found in the on-line version at doi:10.1016/j.intermet.2009.12.019.
882
I. Kock et al. / Intermetallics 18 (2010) 877–882
References [1] Ullakko K, Huang JK, Kantner CC, O’Handley RC, Kokorin VV. Appl Phys Lett 1996;69:1966. [2] Kakeshita T, Fukuda T. Mater Sci Forum 2002;394:531. [3] Cui J, Shield TW, James RD. Acta Mater 2004;52:35–47. [4] Massalski B, Okamoto H, Subramanian PR, Kacprzak L. ASM international’s binary alloy phase diagrams. 2nd ed.; 1996. p. 35–47. [5] Spasova B, Ruffert C, Wurz M, Gatzen H. Phys Stat Sol A 2008;205:2307. [6] Vokoun D, Shih J, Chin T, Hu C. J Magn Magn Mater 2004;281:105. [7] Inoue S, Inoue K, Koterazawa K, Mizuuchi K. Mater Sci Eng 2003;A339:29. [8] Sugimura Y, Cohen-Karni I, McCluskey P, Vlassak JJ. J Mater Res 2005;20:2279. [9] Kock I, Edler T, Mayr SG. J Appl Phys 2008;103:046108. [10] Edler T, Buschbeck J, Mickel C, Fa¨hler S, Mayr SG. New J Phys 2008;10:063007. [11] Heczko 0, Soroka A, Hannula S-P. Appl Phys Lett 2008;93:022503. [12] Duwez P, Willens RH, Klement W. J Appl Phys 1960;31:1136. [13] Jones H. Rep Prog Phys 1973;36:1425.
[14] Thienhaus S, Zamponi C, Rumpf H, Hattrick-Simpers J, Takeuchiand I, Ludwig A. Materials Research Society Symposium Proceeding 894, Warrendale Pa 0894-LL06-06.1; 2006. [15] Zarnetta R, Savan A, Thienhaus S, Ludwig A. Appl Surf Sci 2007;254:743–8. [16] Hamann S, Ehmann M, Thienhaus S, Savan A, Ludwig A. Sensor Actuator: A Phys 2008;147(2):576–82. [17] Opahle I, Kopernik K, Nitzsche U, Richter M. Appl Phys Lett 2009;94:072508. [18] Kakeshita T, Fukuda T. Sci Tech Adv Mater 2006;7:350. [19] Viret M, Samson Y, Warin P, Marty A, Ott F, Søndergård E, et al. Phys Rev Lett 2000;85:3962. [20] Guenin G, Gobin PF. Metall Mater Trans 1982;13:1127–34. [21] Kato H, Liang Y, Taya M. Scripta Mater 2002;46:471–5. [22] Dondl P, Bhattacharya K. Proc Appl Math Mech 2007;7:1151207. [23] Seki K, Kura H, Sato T, Taniyama T. J Appl Phys 2008;103:063910. [24] Deryavko I, Lanin A, Taubini M. Powder Metall Met Ceram 1978;17:137–40. [25] Kang S, Lee Y, Lim Y, Nam J, Nam T, Kim Y. Scripta Mater 2008;59:1186–9. [26] Wollants P, Roos J, Delaely L. Progr Mater Sci 1993;37:227–88.