Shape memory effect and magnetic properties of Co-Fe ferromagnetic shape memory alloys

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1 Shape memory effect and magnetic properties of Co-Fe ferromagnetic shape memory alloys Yunqing Ma*, Shuiyuan Yang, Yuxia Deng, Cuiping Wang, Xingjun Liu Department of Materials Science and Engineering, Xiamen University, Xiamen, 422 Siming South Road, Fujian rovince,. R. China ABSTRACT In 1996, after Ullakko et al. first reported a measurement of 0.2% strain along the [001] direction of a NiB2BMnGa single crystal when subjected to a magnetic field of 8KOe at 265K, the researches of ferromagnetic shape memory alloys became popular. The reversible strain induced by magnetic-field in NiB2BMnGa has been proved to be nearly 10% in a magnetic field of less than 1T, which is much higher than that of the rare-earth giant magnetostrictive alloys. But the high brittleness of NiB2BMnGa hinders its practical application. Therefore, the development of ferromagnetic shape memory alloys with good ductility is of primary importance. In this paper, the microstructure, martensitic transformation temperature, shape memory effect, as well as magnetic and mechanical properties of Co-Fe alloys were investigated by optical observation, X-ray diffraction, DSC, bending test and vibrating sample magnetometer. It was confirmed that the shape memory effect in Co-Fe alloys is associated with the fcc/hcp martensitic transformation. Moreover, Co-Fe alloys exhibit high saturation magnetization with the values of above 170 emu/g, which are much higher than that of NiB2BMnGa (66 emu/g). So the driving force under magnetic field will be large for Co-Fe alloys. Additionally, Co-Fe alloys possess good ductility for practical application with tensile elongation higher than 17.5%. All these results indicate that Co-Fe alloys are promising candidates for developing as ferromagnetic shape memory alloys. Keywords: Martensitic transformation, ferromagnetic shape memory alloy, Co-Fe alloy, ductility 1. INTRODUCTION Since the discovery of shape memory phenomenon, many kinds of shape memory alloys (SMAs) have been developed [1-2]. Until now, Ti-Ni-based alloys exhibit the best performance. Not only do they bear the shape memory effect (SME) and superelasticity similar with other SMAs, they also exhibit good mechanical properties (especially ductility), excellent corrosion and abrasion resistance etc. So most of the commercial applications of SMAs have been done for Ti- Ni-based alloys up to now, such as various sensors and actuators in automobile and oil industry and in safety devices, a flap in air-conditioner, coffee maker, antenna for mobile phones, medical applications such as orthodontics wire, guide wire and various stents etc [3]. The outstanding characters of SMAs are their superelasticity and large thermally recoverable strains and stresses that can be as large as 10% and 800 Ma, respectively. However, heating or especially cooling through a transformation temperature is not a satisfactorily fast method of actuation for many applications. There is a pressing need to increase the response speed, i.e. shorten the feedback time of SMA. In 1996, after Ullakko et al. first reported a measurement of 0.2% strain along the [001] direction of a NiB2BMnGa single crystal when subjected to a magnetic field of 8KOe at 265K [4], substantial efforts have been devoted to developing ferromagnetic shape memory alloys (FSMAs) as they open the possibilities of changing the shapes and dimensions of a SMA by an external magnetic field in addition to stress and temperature. Owing to the fact that their driving force is provided by the outside magnetic field, FSMAs exhibit much higher response speed than conventional SMAs, which are expected to greatly extend the potentials of SMAs. Such fascinating phenomenon can be achieved either by the magnetic-filed-induced reversible redistribution of martensitic variants or the magnetic-filed-induced martensitic transformation from the parent phase. Many FSMAs systems have been developed, including NiMnGa [5-7], Fed [8], Fet [9], CoNi(Al, Ga) [10,11], NiFeGa [12], NiMn(In, Sn, Sb) [13] etc. Until now, the reversible strain induced by magnetic field in FSMAs has been proved to be nearly 10% in a magnetic field of less than 1T [14], which is much *myq@xmu.edu.cnt;t phone ; fax International Conference on Smart Materials and Nanotechnology in Engineering edited by Shanyi Du, Jinsong Leng, Anand K. Asundi roc. of SIE Vol. 6423, 64233I, (2007) X/07/$18 doi: / roc. of SIE Vol I-1

2 the radiation. = [18], higher than that of the rare-earth giant magnetostrictive alloys. However, there are still some problems involved in the practical application of these FSMAs, where good ductility, high martensitic and magnetic transformation temperatures are required. Several Co-based alloys such as CoNi [15], CoAl [16], CoSi [17] etc. have been reported to exhibit SME rooted from the martensitic transformation between the γ phase with a face centered cubic (fcc) structure at high temperature and the ε phase with a hexagonal closepacked (hcp) structure at low temperature. One important property of these Co-based SMAs is their magnetisms with high Curie temperatures (TBcB), hence they are considered to be attractive candidates for FSMAs. Recently, the present authors have found that Co-Fe alloys also exhibit such SME, and possess good ductility and high saturation magnetization, which means that the driving force under magnetic field will be large. The intention of this paper is to report the microstructure, SME, magnetic and mechanical properties in polycrystalline Co-xFe (x=2-6 wt.%) alloys. 2. EXERIMENTAL Co-xFe (x=2-6 wt.%) binary alloy buttons were prepared by melting pure cobalt (99.9%) and iron (99.7%) under argon atmosphere using an arc melter. Each specimen of about 50g was remelted four times to ensure uniformity. To facilitate description, the composition of Co-Fe alloys will be denoted in weight percent hereafter in this paper. The specimens were hot-rolled to make thin plates of 1 mm thickness, and then solution-treated at 1200 for 1 h followed by quenching in water. The microstructure was observed by optical microscopy (OM) and the phase structure was identified at room temperature by a Regaku D/Max 2200 C x-ray diffractometer with Cu KBaB The martensitic transformation temperatures of the alloys were determined by the differential scanning calorimetry (DSC) (Netzsch STA 449) at the cooling and heating rates of 10 C/min, where the transformation temperatures are defined as a cross point of the base line and the tangent of the maximum or minimum inclination. The magnetic transition temperatures were measured by the modified thermogravimetry (TG) in the Netzsch STA 449 with a couple of small NdFeB magnets below the balance outside the chamber that monitors the changes of the sample s nominal weight near the magnetic transitions, which reflects the variations of the magnetic force. The magnetization curves were determined by a LDJ9600 vibrating sample magnetometer (VSM). The mechanical properties were measured by tensile tests at ambient temperature using a Galdabini Sun2500 machine at a crosshead speed of 0.2 mm/min. The tensile direction was parallel to the rolling direction. The size of the gauge part of 1/2 the tensile specimen was 3 mm wide, 1.0 mm thick, and 10 mm long according to the relationship of LB0B 5.65 A where A is the cross-sectional area and LB0B length of the gauge part. 3 The SME was evaluated by bending a sheet specimen with dimensions of mm into a round shape at room temperature (surface strain εbib), unloading it (surface strain εbcb) and then heating it (surface strain εbhb). The surface strain is defined as ε = t/(2r+t), where t and r are the specimen thickness and the radius of curvature, respectively. The recovered strain εbsmeb are evaluated by εbsmeb =εbcb-εbhb. 3. RESULTS AND DISCUSSION 3.1 Microstructure X-ray diffraction patterns of Co-xFe (x=2-6) alloys tested at room temperature are shown in Fig. 1. All the diffraction peaks could be indexed by two phases, i.e. γ phase with fcc structure and ε phase with hcp structure. The diffraction peak of γ(111) overlaps with ε(002), and the characteristic peaks of γ and ε phase are γ(200) and ε(211), respectively. Co-2Fe alloy exhibits mainly ε martensite with little residual γ phase, implying most of the parent phase have transformed to martensite, as shown in Fig. 1. When Fe content increased from 2 to 5.6, the intensities of the residual γ phase peaks increased gradually, implying the amount of ε martensite decrease gradually with the increase of Fe content. When Fe content is higher than 5.65 wt.%, all the reflection peaks can be indexed with γ phase, and no other phases could be roc. of SIE Vol I-2

3 are monitored. This means that the martensitic transformation start temperatures of Co-xFe alloys are lower than room temperature (25 ) when Fe content is higher than 5.65 wt.%. Fig. 2 shows the optical micrograph of Co-4Fe alloy quenched from 1200 in the γ single-phase region. A typical band microstructure of the ε martensite is observed. The white region around the band structure is the residual γ phase. The characters of the microstructure are similar with that of other Co-based alloys [16,17] and Fe-Mn-based alloys [19]. γ (111) ε(002) γ (200) γ (220) γ (311) γ (222) X=6 Intensity (A.U.) X=5.8 X=5.7 X=5.65 X=5.6 X=5.4 ε (210) ε (211) ε (212) ε (213) X=4 X= theta (deg) Fig. 1 X-ray diffraction patterns of Co-xFe (x=2-6) alloys quenched from 1200 Fig.2 Optical micrograph of Co-4Fe alloy quenched from Martensitic transformation & SME Heat flow (mw/mg) exo Heating Cooling Temperature ( O C) Fig. 4 Relationship between martensitic transformation Fig. 3 DSC curves of Co-2Fe alloy temperatures and Fe contents in Co-Fe alloys Fig. 3 shows the DSC curves of Co-2Fe alloy during heating and cooling. The clear endothermic peak appeared on the heating DSC curve is associated with the reverse martensitic transformation from martensite to cubic austenite, where the austenite transformation starting temperature ABsB, finishing temperature B ABf 359.4, and 378.6, respectively. The exothermic peak indicating the forward martensitic transformation from austenite to martensite occurs on the cooling M s temperature ( O C) O C 276 O C Ms ( O C) = X (wt.% Fe) Fe content (wt.%) Ms temperature 25 O C roc. of SIE Vol I-3

4 are DSC curve, and the martensite starting temperature MBsB, finishing temperature MBfB and 236.4, respectively. In the same way, the transformation temperatures were determined for other Co-xFe (x=2-6) alloys. However, no peaks could be detected on the cooling and heating curves of Co-Fe alloys with high Fe content. This may due to two facts: (1) the amount of the martensite decreased with the increasement of Fe content, as indicated in the previous part. (2) The γ/ε martensitic transformation is one of the simplest transformations crystallographically, since both the fcc and hcp phases have closed-packed structures and the martensitic transformation is achieved by shearing along <112>BγB on every other {111}BγB plane [16]. So the latent heat associated with the phase transformation will be small, which may be beyond the sensitivity of the DSC equipment. More studies should be conducted on this item. However, we can estimate the martensitic transformation temperatures using the known data. The martensitic transformation starting temperature of pure Co is 417. The martensitic transformation starting temperature of Co-2Fe is 276.6, and the martensitic transformation starting temperature of Co-5.6Fe is 25, as indicated from the measurement of X-ray diffraction. From these data, the relationship between the martensitic transformation temperatures and the Fe content is estimated and plotted in Fig.4. The martensitic transformation temperatures are almost linearly decreased with increasing Fe content. MBsB ( ) = (wt.% Fe) 1, as Shape recovery (OMS%) 0,2 - OO J!I!I!F! I l Temperature ( C) Fig. 5 Shape recovery of Co-4Fe alloy gradually heated from room temperature to 600 The SME of Co-4Fe alloy was illustrated in Fig. 5. After a surface strain of 1.47% was initially conducted by bending at room temperature, the specimen was heated to 600 to investigate the change of temperature-dependent surface strain. The surface strain was measured at each heating step, and the heating temperature being increased from 100 to 600 at intervals of 100. The shape recovery of Co-4Fe alloy occurs slightly by heating up to 200, then drastically changes in the temperature range from 200 to 300, which corresponding to the reverse martensitic transformation. The total recoverable strain, 0.86%, is lower than that of CoAl alloys (1.1%) [16], but higher than that of CoSi alloys (0.34%) [17]. It was well known that good SME appeared in alloys with thermoelastic martensitic transformation, and highly reversible crystallogphy. The phase transformation between disordered fcc and hcp in Co-based alloys belongs to the non-thermoelastic martensitic transformation and the recoverable strain are caused by the movement of stacking fault. It is understandable for Co-Fe alloys exhibit inferior SME to NiTi alloys, considering the return path of the basal planes during the transition from ε to γ has a multiplicity [20]. In other words, the mechanism of SME in Co-based alloys is similar to that of Fe-Mn-Si alloys, which is also a disordered structure with non-thermoelastic martensite transformation. Since Fe-Mn-Si alloys exhibit good SME associated with stress-induced martensite after training [1], it is natural to roc. of SIE Vol I-4

5 temperatures for values of think good SME may be also obtained through the stress-induced martensitic transformation in Co-Fe alloys. Further investigations are necessary to demonstrate it. 3.3 Magnetic properties The magnetization curves of two as-quenched Co-Fe alloys measured at room temperature are presented in Fig. 6. One can see that the magnetization curves exhibit typical characteristics of ferromagnetic materials. It is noted that Co-5.3Fe trends to saturate easily due to the fact that it contains more γ phase with high symmetry, as compared with Co-4Fe. The saturation magnetization and magnetic anisotropy constant KB1B Co-4Fe at room temperature calculated from the law of approach as MBsB=171.3 emu/g and KB1B= erg/cm, respectively. These two values are emu/g and erg/cm Co-5.3Fe alloy. It can be seen the MBsB are much higher than that of NiB2BMnGa (66 emu/g) alloy [21], CoNi (124 emu/g) alloy [15] and CoAl (120 emu/g) alloy [22]. So the driving force under magnetic field will be large for Co-Fe alloys. The TBcB of Co-Fe alloys were determined to be 1085 for Co-4Fe and 1075 for Co-5.3Fe, as shown in Fig.7, suggesting a wider temperature range for Co-Fe alloys in practical application. The addition of the Fe seems to decrease the Curie temperatures of Co-based alloys. 200 Magnetization (emu/g) Co-4Fe Co-5.3Fe TG ( a.u. ) Co-4Fe Co-5.3Fe Tc = 1085 O C Tc = 1075 O C Magnetic Field (Oe) Temperature ( O C) Fig. 6 Magnetization curves of Co-Fe alloys Fig. 7 TG curves of Co-Fe alloy 3.4 Mechanical properties 400 Co-4Fe Stress (Ma) Co-5.3Fe Strain (%) Fig. 8 Stress-strain curves of Co-Fe alloys Fig. 9 SEM microscopy of the fracture surface of Co-5.3Fe alloy roc. of SIE Vol I-5

6 Fig. 8 shows the stress-strain curves of Co-4Fe and Co-5.3Fe alloys at room temperature. The symbol ( ) represents the fracture point. The tensile stress and the strain were measured to be 383Ma and 17.5% for Co-4Fe alloy, and 382Ma and 30.2% for Co-5.3Fe alloy, respectively. The yield strength (σb0.2b) was measured to be Ma for Co-4Fe alloy, and Ma for Co-5.3Fe alloy. Co-4Fe alloy exhibits higher strength and Co-5.3Fe alloy is more ductile, which should be due to the fact the former containing more ε martensite and the latter containing more γ phase, as described in the previous part. Generally, the stress-strain curves of NiTi and other SMAs exhibit a stress plateau with little work hardening till about 8% before rapid work hardening, which is associated with the reorientation of martensitic variants or the stress-induced martensitic transformation depending on the martensitic or parent-phase state of the alloy [1,2]. However, in the case of Co-Fe alloys, the stress plateau completely disappears, and high work hardening is constantly observed in the stress-strain curves, demonstrating the prematurely dislocation slip during the reversible movements of martensitic variants. It is well known that the strain recovered due to SME occurs solely through the motion of the intervariant boundaries, without any contribution from normal slip. From this point, methods should be tried to increase the critical stress of slip in Co-Fe alloys, so as to improve the SME. Anyway, CoFe alloys are much ductile than present FSMAs. This can be also proved by the SEM microscopy of the fracture surface of Co-5.3Fe alloy after tensile, as shown in Fig. 9. The fracture surface shows clear ductile cup and cone features on a fine scale. 4. CONCLUSION In this paper, the microstructure, martensitic transformation temperature, shape memory effect as well as magnetic and mechanical properties of Co-xFe (x=2-6 wt.%) alloys were investigated to reveal their potentials as ferromagnetic shape memory alloys. The following conclusions are made: 1. Co-xFe (x=2-6 wt.%) alloys exhibit a single γ phase of fcc for x 5.65 and dual phases containing ε martensite and γ for x 5.6. The martensitic transformation temperatures of these alloys are almost linearly decreased with increasing Fe content. 2. The recoverable strain due to the γ(fcc)/ε(hcp) martensitic transformation of Co-4Fe (wt.%) alloys was measured to be 0.86% upon pre-strained to 1.47%. 3. The Curie temperature and the saturation magnetization at room temperature are 1085 and emu/g for Co-4Fe alloy, and 1075 and emu/g for Co-5.3Fe alloy, respectively. 4. Co-Fe alloys are ductile as compared with other ferromagnetic shape memory alloys. The tensile stress and the strain were measured to be 383Ma and 17.5% for Co-4Fe alloy, and 382Ma and 30.2% for Co-5.3Fe alloy, respectively. ACKNOWLEDGMENTS This work was supported by rogram for New Century Excellent Talents in Fujian rovince University (NCETFJ), the Youth Science Foundation of Fujian rovince of China (No.2006F3119), and the Natural Science Foundation of China (No ). The authors are grateful to rof. Y Li of Beihang University for the help in carrying out some experiments. REFERENCES 1. K. Otsuka and C.M. Wayman, Shape Memory Materials, Cambridge University ress, K. Otsuka and X. Ren, Recent development in the research of shape memory alloys, Intermetallics 7, (1999). 3. K. Otsuka and X. Ren, hysical metallurgy of Ti-Ni-based shape memory alloys, rogr. Mater. Sci. 50, (2005). 4. K. Ullakko, J. K. Huang, C. Kantner, R.C. O'Handley and V. V. Kokorin, Large magnetic-field-induced strains in NiB2BMnGa single crystals, Appl. hys. Lett. 69, (1996). roc. of SIE Vol I-6

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