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Thermal and structural alternations in CuAlMnNi shape memory alloy by the effect of different pressure applications
Physica B 521 (2017) 331–338
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
Thermal and structural alternations in CuAlMnNi shape memory alloy by the effect of different pressure applications
MARK
⁎
Canan Aksu Canbay , Tercan Polat Department of Physics, Faculty of Science, University of Firat, 23119 Elazig, Turkey
A R T I C L E I N F O
A BS T RAC T
Keywords: Cu-based shape memory alloys Martensitic transformation Shape memory effect Pressure
In this work the effects of the applied pressure on the characteristic transformation temperatures, the high temperature order-disorder phase transitions, the variation in diffraction peaks and the surface morphology of the CuAlMnNi shape memory alloy was investigated. The evolution of the transformation temperatures was studied by differential scanning calorimetry (DSC) with different heating and cooling rates. The differential thermal analysis measurements were performed to obtain the ordered-disordered phase transformations from room temperature to 900 °C. The characteristic transformation temperatures and the thermodynamic parameters were highly sensitive to variations in the applied pressure and also the applied pressure affected the thermodynamic parameters. The activation energy of the sample according to applied pressure values calculated by Kissinger method. The structural changes of the samples were studied by X-ray diffraction (XRD) measurements and by optical microscope observations at room temperature.
1. Introduction Shape memory alloys (SMAs) are a typical class of functional materials which can remember and return the original shape after considerable deformation or cooling from high temperatures. They have thermoelastic martensitic transformation mechanism and this mechanism belongs to solid-solid diffusionless transformation. Thermoelastic martensitic transformation is a first-order transformation in solid-solid phase and it is explained by the collective motion of atoms. In Cu-based SMAs, the thermoelastic martensitic transformation depend on crystal structure of the phases involved in transformation process. During the transformation process there are three main mechanisms, which activates the thermoelastic transformation and can be described as: 1) distribution of energy caused by internal friction, 2) storage of energy, 3) heat transfer caused by the latent heat of phase change. From AuCd alloys a number of SMAs have been reported as; Ni-based, Cu-based, Fe-based and also Heusler alloys. Amongst these SMAs, Ni-Ti and Cu-based alloys are commercially attractive for manufacture and practical applications utilizing their shape memory effect (SME) behavior [1–5]. Cu3Al-based SMAs have unique thermomechanical properties such as SME and superelasticity (SE). The CuAlMn and CuAlNi ternary systems are good candidates due to their wide range transformation temperature and small hysterisis which are required for SME and in particular superelastic CuAlMn SMAs exhibit similar mechanical
⁎
Corresponding author. E-mail address: caksu@firat.edu.tr (C.A. Canbay).
http://dx.doi.org/10.1016/j.physb.2017.07.017 Received 11 May 2017; Received in revised form 6 July 2017; Accepted 7 July 2017 Available online 08 July 2017 0921-4526/ © 2017 Elsevier B.V. All rights reserved.
properties as NiTi, but have better thermal and electrical properties. Cu-based SMAs exhibit a martensitic transformation on cooling (Martensite phase), and austenite transformation on heating (Austenite phase). On cooling process, the close-packed structures are characterized by long period stacking order. The characteristic transformation temperatures of austenite and martensite phases are called as As, Af, Ms and Mf, respectively. Cu-based SMAs have A2-type disordered structure at high temperatures and undergo ordered structure by cooling. The ordered structures in different SMAs can be determined as B2, DO3 or L21 and these structures transform into martensites with further cooling. In Cu3Al-based alloys, by the further cooling we observe these structures as bcc (beta) → (9R, 18R). The phases that can occur for high temperatures and low temperatures close to the composition of Cu3Al can be determined as in two steps; i) the disordered β phase (A2 or bcc structure) is the equilibrium structure at high temperatures, ii) the equilibrium phases at lower temperatures are α-Cu, T3-Cu3Mn2Al, γ2-Cu9Al4 and β-Mn. The metastable disordered bcc β phase undergoes two ordering transitions during cooling, the first one is B2 ordered structure occured on nearest neighbor and the second transition is DO3 structure of Cu3Al (β1 phase) or the L21 structure of the Cu2AlMn occured on next nearest neighbor. So if we summarize this information, a first order DO3 ↔ A2 transition is obtained for the stoichiometric Cu3Al alloy, while for concentrations near Cu2AlMn two second order transitions are found: L21 ↔ B2 and B2 ↔ A2 [5–10].
Physica B 521 (2017) 331–338
C.A. Canbay, T. Polat
was determined by Bruker Model energy dispersive X-ray (EDX) as Cu22.90 Al-4.87 Mn-0.90 Ni (at%). The ingot cut into small pieces and then the specimens were solution-treated at 900 °C for 1 h and quenched in iced-brine water. Then pressure was applied with different sizes on these samples as 280 MPa, 560, MPa, 840 MPa and 1120 MPa. The characteristic transformation temperatures and thermodynamic parameters of the homogenous and pressure applied samples were determined by Shimadzu DSC-60A differential scanning calorimetry with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). The TG/DTA (Shimadzu TA-60 WS) measurements were performed from room temperature to 900 at a heating rate of 25 °C/min to identify order-disorder phase transitions at high temperatures. X-Ray diffraction patterns of the samples were taken by Rigaku RadB-DMAX II diffractometer. The surfaces of the samples were cleaned by an etching solution, consisting of FeCl3-6H2O-methanol-HCl, for optical microscope analysis. Microstructural observations were made using Nikon MA200 model optical metallographic microscope. 3. Results and discussion The X-ray analysis of the homogenous sample and the pressure applied samples were made by CuKα radiation at room temperature. The lattice parameters of the samples are a = 4494 Å; b = 5189 Å; c = 46,610 Å. and the calculated a/b ratio was 0.86. The calculated a/b ratio confirms that parent phase of the samples are ordered and has long period stacking order structure, which means that during the rapid cooling high temperature β-phase is expanded into a L21 ordered structure and can further be transformed into a martensite phase (M18R) at room temperature. The diffraction peaks of the samples are given in Fig. 1. The main diffraction peaks of martensite phase observed in all samples are (202), (0022) and (1210) [11–13]. The average crystallite size (D) for each sample was calculated by the Debye-Scherer equation [14,15].
D= Fig. 1. X-ray diffraction patterns of the homogenous and pressure applied samples.
Crystallite size (nm)
Homogenous 280 560 840 1120
15,70 23,03 46,18 19,05 18,96
(1)
where λ is the wavelength of X-ray (CuKα radiation), B is the peak full width at half maximum and θ is the Bragg angle. The crystallite size of the homogenous sample and pressure applied samples are calculated as and given in Table 1. Along with the increase of the pressure values, besides the martensite plates formed in the material, lattice defects also occured. These lattice defects caused deformation and the deformation in material accumulates in grain boundaries. In order for the deformations not to pass to other grains, the material increases the grain size due to the elastic stretching, there by ensuring that the deformation is dispersed. Because of this, we observed a maximum grain size at 560 MPa applied pressure value with the increase of pressure (Table 1). If the pressure continues to increase, the material reaches maximum flexural tensile strength, and after that point, the material structure deteriorates. These disorders are aimed at preventing breakage and martensite formation at grain boundaries. Other experimental observations show that martensite formations decreased at values after the applied pressure value of 560 MPa and martensite formations not observed at the last applied pressure value. After the maximum flexural tensile strength is exceeded, the grain size is again decreased in order to impart rigidity to the material. The decrease in the particle size values at the final pressure values is also due to this effect. The surface morphology of the samples was observed by optical microscope analysis and given in Fig. 2. The optical microscope observations of the samples revealed that the parent phase is martensite at room temperature as expected from X-ray results. Martensite variants, grains and grain boundaries occured in the structure of the
Table 1 The calculated crystallite size of the samples. Applied pressure value (MPa)
0.9λ B1/2 cos θ
Recently, our work focused on CuAlMnNi SMA that attracted considerable attention due to the ductility of the material. In this work, the effect of applied pressure on thermodynamic parameters such as characteristic transformation temperatures, enthalpy and entropy values on CuAlMnNi SMA was investigated. The structural variations are observed by X-ray analysis and optical microscope observations. 2. Experimental The polycrystalline Cu-Al-Mn-Ni shape memory alloy was produced by arc melting in an argon atmosphere. The composition of the alloy
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Fig. 2. Optical micrographs of the homogenous and pressure applied samples.
performed with different heating/cooling rates as; 5, 15, 25, 35 and 45 °C/min and given in Figs. 3–7. The transformation temperatures and the hysteresis can be determined by the corresponding peak positions of the DSC measurements so the characteristic transformation temperatures of reverse and forward transformations were determined from DSC curves by tangent method. The characteristic austenite and martensite phase start and finish temperatures, equilibrium temperatures, enthalpy and entropy values of each sample with different heating/cooling rates are given in Tables 2–5.
samples are clear and the orientations of martensite variants of each grain is different. These observations supported the X-ray analysis results of the samples. The thermal analysis of the samples was made to determine the effect of pressure on characteristic austenite and martensite transformation temperatures and also thermodynamic parameters. Differential scanning calorimetry (DSC) measurements has been used to follow the reverse and forward transformations from martensite to austenite (M → A) or austenite to martensite (A → M) phases. DSC measurements
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Fig. 3. The DSC curve of the homogenous sample with different heating/cooling rates.
Fig. 4. The DSC curve of the 280 MPa pressure applied sample with different heating/cooling rates.
the equilibrium temperature. The variation in thermodynamic parameters such as, equilibrium temperature (T0), enthalpy (ΔH), and Af-Ms hysteresis value by the effect of applied pressure is given by the graphics with a heating/ cooling rate of 25 °C/min to observe the alternations clearly These curves are given in Figs. 8–10, respectively. The activation energies of the samples calculated by Kissinger method and the activation energy curve of the samples is given in Fig. 11. The calculated activation energy values of samples according to Kissinger method are processed by the following equation [19];
The equilibrium temperature between the reverse and forward phases according to Salzbrenner and Cohen is determined by the following relation [16];
T0 =
1 (Ms + Af ) 2
(2)
where Ms is the martensite start temperature, Af is the austenite finish temperature. The entropy change during the reverse transformation can be determined by the following equation [17,18];
ΔSM→ A =
ΔHM→ A T0
d [ ln(ϕ / T 2m )]/ d (1/ Tm ) = −E / R
(3)
where ΔS is the entropy change, ΔH is the enthalpy change, and T0 is
(4)
where ϕ is the heating rate, Tm the maximum temperature of the DSC
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Fig. 5. The DSC curve of the 560 MPa pressure applied sample with different heating/cooling rates.
Fig. 6. The DSC curve of the 840 MPa pressure applied sample with different heating/cooling rates.
transition (Cu2MnAl type ordered), iii) L21 → 9R or 18R transition.
peak, R is the universal gas constant and E is the activation energy. The calculated activation energy values for homogenous sample, 280 MPa pressure applied sample, 560 MPa pressure applied sample and 840 MPa pressure applied sample are; 166.45, 79.08, 108.25 and 111.32 kJ/mol, respectively. The order-disorder phase transitions at high temperatures determined by differential thermal analysis TG/DTA measurements with a heating rate of 25 °C/min and given in Fig. 12. Cu-based SMAs show three phase transition with the increase of the heating, and these transitions are described as β (A2) → β2 (B2) → β1 (L21) [20–22]. The transformation steps observed in this work for all the samples can be described as; i) A2 → B2 transition (CsCl type ordered), ii) B2 → L21
4. Conclusion The effect of pressure on the transformation temperatures, thermodynamic parameters such as enthalpy, entropy and (Af-Ms) hysteresis values of the Cu-Al-Mn-Ni SMA was investigated. The XRD analysis of the samples and the calculated a/b ratio confirms that the structure is M18R martensite, i.e., the parent phase is ordered and has 18R long period stacking order structure. Although the diffractograms of the samples exhibit similar characteristics, some changes were observed in the peak intensities with the variation of the pressure and also we observe cracks in
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Fig. 7. The DSC curve of the 1120 MPa pressure applied sample with a heating/cooling rate of 25 °C/min. Table 2 Transformation temperatures and thermodynamic parameters of homogenous sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
58,24 45,01 52,92 50,35 52,29
62,08 68,09 72,53 75,77 81,19
59,75 59,97 63,25 67,45 68,82
54,18 56,82 59,30 62,45 63,28
1,93 0,27 2,27 2,52 2,24
0,035 0,004 0,038 0,040 0,035
46,29 45,55 46,08 49,13 45,37
43,48 43,30 43,10 48,01 42,89
−1,82 −0,72 −0,53 −0,86 −0,46
−0,033 −0,012 −0,008 −0,013 −0,007
Table 3 Transformation temperatures and thermodynamic parameters of 280 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
68,35 69,51 70,57 73,75 72,74
83,20 92,04 100,10 105,76 113,73
74,29 79,65 83,46 87,69 100,21
76,63 79,96 83,43 85,39 90,89
6,87 7,24 8,02 7,30 6,98
0,089 0,090 0,096 0,085 0,069
70,07 67,89 66,77 65,03 67,99
45,51 49,31 49,51 50,09 48,46
−12,35 −1,47 −3,40 −3,28 −3,70
−0,161 −0,018 −0,040 −0,038 −0,036
Table 4 Transformation temperatures and thermodynamic parameters of 560 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
65,73 65,30 67,44 71,28 68,96
72,29 78,58 85,29 91,57 99,96
68,65 72,62 75,81 79,15 87,26
63,21 66,27 69,22 72,45 76,52
6,93 7,35 7.70 10,37 7,54
0,109 0,110 0,111 0,143 0,098
54,14 53,97 53,15 53,34 53,08
45,47 45,01 44,94 44,55 45,67
−7,43 −4.69 −5,04 −5,15 −4,86
−0,117 −0,070 −0,072 −0,071 −0,063
Table 5 Transformation temperatures and thermodynamic parameters of 840 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
70,70 70,93 71,83 73,58 79,90
81,03 89,62 94,29 101,75 111,04
74,48 80,58 83,43 87,48 94,39
74,65 82,37 85,73 88,91 93,83
6,43 7,45 8,57 9,62 8,82
0,086 0,090 0,099 0,108 0,093
68,27 75,12 77,17 76,07 76,63
53,43 57,39 54,28 53,20 54,21
−13,57 −2,50 −3,20 −3,87 −4,68
−0,181 −0,030 −0,037 −0,043 −0,049
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Fig. 8. The T0 versus pressure curve of the samples with a heating/cooling rate of 25 °C/ min.
Fig. 10. The Af-Ms hysterisis variation versus pressure curve of the samples with a heating/cooling rate of 25 °C/min.
Fig. 9. The enthalpy variation versus pressure curve of the samples with a heating/ cooling rate of 25 °C/min.
Fig. 11. The activation energy curve of the samples according to Kissinger method.
Fig. 12. The DTA curve of the samples with a heating rate of 25 °C/min.
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References
peaks. This is the result of pressure applied on samples. The intensity of the 1120 MPa pressure applied sample is the smallest amongst the others and this also supported the optical and thermal measurements. The complete transformation of the austenite phase to martensite can be observed in the micrographs of the alloys except 1120 MPa pressure applied sample. The martensite plates observed in optical micrographs of all samples were identified as β1 martensite at room temperature. The thermal measurements showed that characteristic transformation temperatures of the samples increased by the increase of applied pressure value. So, the thermodynamic parameters such as enthalpy, entropy, hysterisis value and activation energies varied according to the characteristic transformation temperatures. The increase observed in the characteristic transformation temperatures with the increase of applied pressure is supported by literature. To calculate activation energy of austenite phase transformation, we plotted the curves of ln(ϕ /Tm2) vs. 1000/T and the calculated activation energies changed by the applied pressure as expected. In DTA measurements we observe the order-disorder phase transitions for the homogenous and pressure applied samples as mentioned in literature for Cu-based SMAs. For the homogenous sample, 280 MPa and 560 MPa pressure applied samples, we can not observe the B2 → L21 transition or Cu2MnAl type ordered transition. But 840 MPa and 1120 MPa pressure applied samples we observed B2 → L21 between 400–470 °C. This is mainly due to the effect of the size of applied pressure on samples. In this work we both observed the thermal and structural behavior of the order-disorder phase transition in Cu3Al-based SMAs by the effect of pressure. The order-disorder transition has an influence over the stability of bcc phase and over the martensitic transformation. Martensitic transformation properties are strongly connected with the degree of disorder retained by the samples and also ordering has the effect of stabilization on bcc phase of Cu-based SMAs. We observe two continuous transitions occurred during thermal measurements as L21 ↔ B2 and B2 ↔ A2; these two transitions combine with a single first order DO3 ↔ A2 transition.
[1] D. Dunne, M. Morin, C. Gonzalez, G. Guenin, The effect of quenching treatment on the reversible martensitic transformation in CuAlBe alloys, Mater. Sci. Eng. A 378 (2004) 257–262. [2] J.P. Oliveira, Z. Zeng, T. Omori, N. Zhou, R.M. Miranda, F.M. Braz Fernandes, Improvement of damping properties in laser processed superelastic Cu-Al-Mn shape meory alloys, Mater. Des. 98 (2016) 280–284. [3] D. Yang, Shape meory alloy and smart hybrid composites-advanced materials for 21st century, Mater. Des. 21 (2000) 503–505. [4] M.O. Moroni, R. Saldivia, M. Sarrazin, A. Sepulveda, Damping characteristics of a Cu-Zn-Al-Ni shape memory alloy, Mater. Sci. Eng. A 335 (1–2) (2002) 313–319. [5] U. Sari, I. Aksoy, Microstructural analysis of self-accomodating martensites in Cu11.92 wt% Al-3.78wt% Ni shape memory alloy, J. Mater. Process. Technol. 195 (2008) 72–76. [6] T. Omori, N. Koeda, Y. Sutou, R. Kainuma, K. Ishida, Superplasticity of Cu-Al-MnNi shape memory alloy, Mater. Trans. 11 (2007) 2914–2918. [7] Y. Sutou, N. Koeda, T. Omori, R. Kainuma, K. Ishida, Effects of ageing on bainitic and thermally induced martensitic transformations in ductile Cu-Al-Mn based shape memory alloys, Acta Mater. 57 (2009) 5748–5758. [8] Kainuma, S. Takahashi, K. Ishida, Thermoelastic martensite and shape memory effect in ductile Cu-Al-Mn alloys, Metall. Mater. Trans. A 27A (1996) 2187–2195. [10] C. Aksu Canbay, Z. Karagoz, The effect of quaternary element on the thermodynamic parameters and structure of CuAlMn shape memory alloys, Appl. Phys. A (2013). http://dx.doi.org/10.1007/s00339-013-7880-3. [11] S. Stanciu, L.G. Bujoreanu, Formation of β1 stress-induced martensite in the presence of γ-phase, in a Cu-Al-Ni-Mn-Fe shape memory alloy, Mater. Sci. Eng. A 481–482 (2008) 494–499. [12] H.Y. Peng, Y.D. Yu, D.X. Li, High resolution electron microscopy studies of martensite around precipitates in a Cu-Al-Ni-Mn-Ti shape memory alloy, Acta Mater. 45 (12) (1997) 5153–5161. [13] U.S. Mallik, V. Sampath, Influence of quaternary alloying additions on transformation temperatures and shape memory properties of Cu-Al-Mn shape memory alloy, J. Alloy. Compd. 469 (2009) 156–163. [14] N. Suresh, U. Ramamurty, Aging response and its effect on the functional properties of Cu-Al-Ni shape memory alloys, J. Alloy. Compd. 449 (2008) 113–118. [15] D. Sonia, P. Rotaru, S. Rizescu, N.G. Bizdoaca, Thermal study of a shape memory alloy (SMA) spring actuator designed to insure the motion of a barrier structure, J. Therm. Anal. Calorim. (2013). http://dx.doi.org/10.1007/s10973-012-2369-4 (111-1255-1262). [16] R.J. Salzbrenner, M. Cohen, On the thermodynamics of thermoelastic martensitic transformations, Acta Metall. 27 (1979) 739–748. [17] R.A. Portier, P. Ochin, A. Pasko, G.E. Monastyrsky, A.V. Gilchuk, V.I. Kolomytsev, Y.N. Koval, Spark plasma sintering of Cu–Al– Ni shape memory alloy, J. Alloys Compd. (2012). http://dx.doi.org/10.1016/j.jallcom.2012.02.145. [18] M.O. Prado, P.M. Decarte, F. Lovey, Martensitic transformation in Cu-Mn-Al alloys, Scr. Mater. 33 (1995) 878–883. [19] H.E. Kissinger, Reaction kinetics in differential thermal analysis, Anal. Chem. 29– 11 (1957) 1702–1706. [20] Y. Sutou, R. Kainuma, K. Ishida, Effect of alloying elements on the shape memory properties of ductile Cu-Al-Mn alloys, Mater. Sci. Eng. A 273–275 (1999) 375–379. [21] U.S. Mallik, V. Sampath, Influence of quaternary alloying additions on transformation temperatures and shape memory properties of Cu-Al-Mn shape memory alloy, J. Alloy. Compd. 469 (2009) 156–163. [22] J. Fernandez, A.V. Benedetti, J.M. Guilemany, X.M. Zhang, Thermal stability of the martensitic transformation of Cu-Al-Ni-Mn-Ti, Mater. Sci. Eng. A 438–440 (2006) 723–725.
Acknowledgements This work is financially supported by FÜBAP, Project No: FF.15.14. Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at http://doi:10.1016/j.physb.2017.07.017.
338
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
Thermal and structural alternations in CuAlMnNi shape memory alloy by the effect of different pressure applications
MARK
⁎
Canan Aksu Canbay , Tercan Polat Department of Physics, Faculty of Science, University of Firat, 23119 Elazig, Turkey
A R T I C L E I N F O
A BS T RAC T
Keywords: Cu-based shape memory alloys Martensitic transformation Shape memory effect Pressure
In this work the effects of the applied pressure on the characteristic transformation temperatures, the high temperature order-disorder phase transitions, the variation in diffraction peaks and the surface morphology of the CuAlMnNi shape memory alloy was investigated. The evolution of the transformation temperatures was studied by differential scanning calorimetry (DSC) with different heating and cooling rates. The differential thermal analysis measurements were performed to obtain the ordered-disordered phase transformations from room temperature to 900 °C. The characteristic transformation temperatures and the thermodynamic parameters were highly sensitive to variations in the applied pressure and also the applied pressure affected the thermodynamic parameters. The activation energy of the sample according to applied pressure values calculated by Kissinger method. The structural changes of the samples were studied by X-ray diffraction (XRD) measurements and by optical microscope observations at room temperature.
1. Introduction Shape memory alloys (SMAs) are a typical class of functional materials which can remember and return the original shape after considerable deformation or cooling from high temperatures. They have thermoelastic martensitic transformation mechanism and this mechanism belongs to solid-solid diffusionless transformation. Thermoelastic martensitic transformation is a first-order transformation in solid-solid phase and it is explained by the collective motion of atoms. In Cu-based SMAs, the thermoelastic martensitic transformation depend on crystal structure of the phases involved in transformation process. During the transformation process there are three main mechanisms, which activates the thermoelastic transformation and can be described as: 1) distribution of energy caused by internal friction, 2) storage of energy, 3) heat transfer caused by the latent heat of phase change. From AuCd alloys a number of SMAs have been reported as; Ni-based, Cu-based, Fe-based and also Heusler alloys. Amongst these SMAs, Ni-Ti and Cu-based alloys are commercially attractive for manufacture and practical applications utilizing their shape memory effect (SME) behavior [1–5]. Cu3Al-based SMAs have unique thermomechanical properties such as SME and superelasticity (SE). The CuAlMn and CuAlNi ternary systems are good candidates due to their wide range transformation temperature and small hysterisis which are required for SME and in particular superelastic CuAlMn SMAs exhibit similar mechanical
⁎
Corresponding author. E-mail address: caksu@firat.edu.tr (C.A. Canbay).
http://dx.doi.org/10.1016/j.physb.2017.07.017 Received 11 May 2017; Received in revised form 6 July 2017; Accepted 7 July 2017 Available online 08 July 2017 0921-4526/ © 2017 Elsevier B.V. All rights reserved.
properties as NiTi, but have better thermal and electrical properties. Cu-based SMAs exhibit a martensitic transformation on cooling (Martensite phase), and austenite transformation on heating (Austenite phase). On cooling process, the close-packed structures are characterized by long period stacking order. The characteristic transformation temperatures of austenite and martensite phases are called as As, Af, Ms and Mf, respectively. Cu-based SMAs have A2-type disordered structure at high temperatures and undergo ordered structure by cooling. The ordered structures in different SMAs can be determined as B2, DO3 or L21 and these structures transform into martensites with further cooling. In Cu3Al-based alloys, by the further cooling we observe these structures as bcc (beta) → (9R, 18R). The phases that can occur for high temperatures and low temperatures close to the composition of Cu3Al can be determined as in two steps; i) the disordered β phase (A2 or bcc structure) is the equilibrium structure at high temperatures, ii) the equilibrium phases at lower temperatures are α-Cu, T3-Cu3Mn2Al, γ2-Cu9Al4 and β-Mn. The metastable disordered bcc β phase undergoes two ordering transitions during cooling, the first one is B2 ordered structure occured on nearest neighbor and the second transition is DO3 structure of Cu3Al (β1 phase) or the L21 structure of the Cu2AlMn occured on next nearest neighbor. So if we summarize this information, a first order DO3 ↔ A2 transition is obtained for the stoichiometric Cu3Al alloy, while for concentrations near Cu2AlMn two second order transitions are found: L21 ↔ B2 and B2 ↔ A2 [5–10].
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was determined by Bruker Model energy dispersive X-ray (EDX) as Cu22.90 Al-4.87 Mn-0.90 Ni (at%). The ingot cut into small pieces and then the specimens were solution-treated at 900 °C for 1 h and quenched in iced-brine water. Then pressure was applied with different sizes on these samples as 280 MPa, 560, MPa, 840 MPa and 1120 MPa. The characteristic transformation temperatures and thermodynamic parameters of the homogenous and pressure applied samples were determined by Shimadzu DSC-60A differential scanning calorimetry with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). The TG/DTA (Shimadzu TA-60 WS) measurements were performed from room temperature to 900 at a heating rate of 25 °C/min to identify order-disorder phase transitions at high temperatures. X-Ray diffraction patterns of the samples were taken by Rigaku RadB-DMAX II diffractometer. The surfaces of the samples were cleaned by an etching solution, consisting of FeCl3-6H2O-methanol-HCl, for optical microscope analysis. Microstructural observations were made using Nikon MA200 model optical metallographic microscope. 3. Results and discussion The X-ray analysis of the homogenous sample and the pressure applied samples were made by CuKα radiation at room temperature. The lattice parameters of the samples are a = 4494 Å; b = 5189 Å; c = 46,610 Å. and the calculated a/b ratio was 0.86. The calculated a/b ratio confirms that parent phase of the samples are ordered and has long period stacking order structure, which means that during the rapid cooling high temperature β-phase is expanded into a L21 ordered structure and can further be transformed into a martensite phase (M18R) at room temperature. The diffraction peaks of the samples are given in Fig. 1. The main diffraction peaks of martensite phase observed in all samples are (202), (0022) and (1210) [11–13]. The average crystallite size (D) for each sample was calculated by the Debye-Scherer equation [14,15].
D= Fig. 1. X-ray diffraction patterns of the homogenous and pressure applied samples.
Crystallite size (nm)
Homogenous 280 560 840 1120
15,70 23,03 46,18 19,05 18,96
(1)
where λ is the wavelength of X-ray (CuKα radiation), B is the peak full width at half maximum and θ is the Bragg angle. The crystallite size of the homogenous sample and pressure applied samples are calculated as and given in Table 1. Along with the increase of the pressure values, besides the martensite plates formed in the material, lattice defects also occured. These lattice defects caused deformation and the deformation in material accumulates in grain boundaries. In order for the deformations not to pass to other grains, the material increases the grain size due to the elastic stretching, there by ensuring that the deformation is dispersed. Because of this, we observed a maximum grain size at 560 MPa applied pressure value with the increase of pressure (Table 1). If the pressure continues to increase, the material reaches maximum flexural tensile strength, and after that point, the material structure deteriorates. These disorders are aimed at preventing breakage and martensite formation at grain boundaries. Other experimental observations show that martensite formations decreased at values after the applied pressure value of 560 MPa and martensite formations not observed at the last applied pressure value. After the maximum flexural tensile strength is exceeded, the grain size is again decreased in order to impart rigidity to the material. The decrease in the particle size values at the final pressure values is also due to this effect. The surface morphology of the samples was observed by optical microscope analysis and given in Fig. 2. The optical microscope observations of the samples revealed that the parent phase is martensite at room temperature as expected from X-ray results. Martensite variants, grains and grain boundaries occured in the structure of the
Table 1 The calculated crystallite size of the samples. Applied pressure value (MPa)
0.9λ B1/2 cos θ
Recently, our work focused on CuAlMnNi SMA that attracted considerable attention due to the ductility of the material. In this work, the effect of applied pressure on thermodynamic parameters such as characteristic transformation temperatures, enthalpy and entropy values on CuAlMnNi SMA was investigated. The structural variations are observed by X-ray analysis and optical microscope observations. 2. Experimental The polycrystalline Cu-Al-Mn-Ni shape memory alloy was produced by arc melting in an argon atmosphere. The composition of the alloy
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Fig. 2. Optical micrographs of the homogenous and pressure applied samples.
performed with different heating/cooling rates as; 5, 15, 25, 35 and 45 °C/min and given in Figs. 3–7. The transformation temperatures and the hysteresis can be determined by the corresponding peak positions of the DSC measurements so the characteristic transformation temperatures of reverse and forward transformations were determined from DSC curves by tangent method. The characteristic austenite and martensite phase start and finish temperatures, equilibrium temperatures, enthalpy and entropy values of each sample with different heating/cooling rates are given in Tables 2–5.
samples are clear and the orientations of martensite variants of each grain is different. These observations supported the X-ray analysis results of the samples. The thermal analysis of the samples was made to determine the effect of pressure on characteristic austenite and martensite transformation temperatures and also thermodynamic parameters. Differential scanning calorimetry (DSC) measurements has been used to follow the reverse and forward transformations from martensite to austenite (M → A) or austenite to martensite (A → M) phases. DSC measurements
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Fig. 3. The DSC curve of the homogenous sample with different heating/cooling rates.
Fig. 4. The DSC curve of the 280 MPa pressure applied sample with different heating/cooling rates.
the equilibrium temperature. The variation in thermodynamic parameters such as, equilibrium temperature (T0), enthalpy (ΔH), and Af-Ms hysteresis value by the effect of applied pressure is given by the graphics with a heating/ cooling rate of 25 °C/min to observe the alternations clearly These curves are given in Figs. 8–10, respectively. The activation energies of the samples calculated by Kissinger method and the activation energy curve of the samples is given in Fig. 11. The calculated activation energy values of samples according to Kissinger method are processed by the following equation [19];
The equilibrium temperature between the reverse and forward phases according to Salzbrenner and Cohen is determined by the following relation [16];
T0 =
1 (Ms + Af ) 2
(2)
where Ms is the martensite start temperature, Af is the austenite finish temperature. The entropy change during the reverse transformation can be determined by the following equation [17,18];
ΔSM→ A =
ΔHM→ A T0
d [ ln(ϕ / T 2m )]/ d (1/ Tm ) = −E / R
(3)
where ΔS is the entropy change, ΔH is the enthalpy change, and T0 is
(4)
where ϕ is the heating rate, Tm the maximum temperature of the DSC
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Fig. 5. The DSC curve of the 560 MPa pressure applied sample with different heating/cooling rates.
Fig. 6. The DSC curve of the 840 MPa pressure applied sample with different heating/cooling rates.
transition (Cu2MnAl type ordered), iii) L21 → 9R or 18R transition.
peak, R is the universal gas constant and E is the activation energy. The calculated activation energy values for homogenous sample, 280 MPa pressure applied sample, 560 MPa pressure applied sample and 840 MPa pressure applied sample are; 166.45, 79.08, 108.25 and 111.32 kJ/mol, respectively. The order-disorder phase transitions at high temperatures determined by differential thermal analysis TG/DTA measurements with a heating rate of 25 °C/min and given in Fig. 12. Cu-based SMAs show three phase transition with the increase of the heating, and these transitions are described as β (A2) → β2 (B2) → β1 (L21) [20–22]. The transformation steps observed in this work for all the samples can be described as; i) A2 → B2 transition (CsCl type ordered), ii) B2 → L21
4. Conclusion The effect of pressure on the transformation temperatures, thermodynamic parameters such as enthalpy, entropy and (Af-Ms) hysteresis values of the Cu-Al-Mn-Ni SMA was investigated. The XRD analysis of the samples and the calculated a/b ratio confirms that the structure is M18R martensite, i.e., the parent phase is ordered and has 18R long period stacking order structure. Although the diffractograms of the samples exhibit similar characteristics, some changes were observed in the peak intensities with the variation of the pressure and also we observe cracks in
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Fig. 7. The DSC curve of the 1120 MPa pressure applied sample with a heating/cooling rate of 25 °C/min. Table 2 Transformation temperatures and thermodynamic parameters of homogenous sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
58,24 45,01 52,92 50,35 52,29
62,08 68,09 72,53 75,77 81,19
59,75 59,97 63,25 67,45 68,82
54,18 56,82 59,30 62,45 63,28
1,93 0,27 2,27 2,52 2,24
0,035 0,004 0,038 0,040 0,035
46,29 45,55 46,08 49,13 45,37
43,48 43,30 43,10 48,01 42,89
−1,82 −0,72 −0,53 −0,86 −0,46
−0,033 −0,012 −0,008 −0,013 −0,007
Table 3 Transformation temperatures and thermodynamic parameters of 280 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
68,35 69,51 70,57 73,75 72,74
83,20 92,04 100,10 105,76 113,73
74,29 79,65 83,46 87,69 100,21
76,63 79,96 83,43 85,39 90,89
6,87 7,24 8,02 7,30 6,98
0,089 0,090 0,096 0,085 0,069
70,07 67,89 66,77 65,03 67,99
45,51 49,31 49,51 50,09 48,46
−12,35 −1,47 −3,40 −3,28 −3,70
−0,161 −0,018 −0,040 −0,038 −0,036
Table 4 Transformation temperatures and thermodynamic parameters of 560 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
65,73 65,30 67,44 71,28 68,96
72,29 78,58 85,29 91,57 99,96
68,65 72,62 75,81 79,15 87,26
63,21 66,27 69,22 72,45 76,52
6,93 7,35 7.70 10,37 7,54
0,109 0,110 0,111 0,143 0,098
54,14 53,97 53,15 53,34 53,08
45,47 45,01 44,94 44,55 45,67
−7,43 −4.69 −5,04 −5,15 −4,86
−0,117 −0,070 −0,072 −0,071 −0,063
Table 5 Transformation temperatures and thermodynamic parameters of 840 MPa pressure applied sample with different heating/cooling rates (5, 15, 25, 35 and 45 °C/min). Heating/cooling rate (°C/min)
As (℃)
Af (℃)
Amax (℃)
T0 (℃)
ΔHM→A (J/g)
ΔSM→A (J/g °C)
Ms (℃)
Mf (℃)
ΔHA→M (J/g)
ΔSA→M (J/g °C)
5 15 25 35 45
70,70 70,93 71,83 73,58 79,90
81,03 89,62 94,29 101,75 111,04
74,48 80,58 83,43 87,48 94,39
74,65 82,37 85,73 88,91 93,83
6,43 7,45 8,57 9,62 8,82
0,086 0,090 0,099 0,108 0,093
68,27 75,12 77,17 76,07 76,63
53,43 57,39 54,28 53,20 54,21
−13,57 −2,50 −3,20 −3,87 −4,68
−0,181 −0,030 −0,037 −0,043 −0,049
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Fig. 8. The T0 versus pressure curve of the samples with a heating/cooling rate of 25 °C/ min.
Fig. 10. The Af-Ms hysterisis variation versus pressure curve of the samples with a heating/cooling rate of 25 °C/min.
Fig. 9. The enthalpy variation versus pressure curve of the samples with a heating/ cooling rate of 25 °C/min.
Fig. 11. The activation energy curve of the samples according to Kissinger method.
Fig. 12. The DTA curve of the samples with a heating rate of 25 °C/min.
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References
peaks. This is the result of pressure applied on samples. The intensity of the 1120 MPa pressure applied sample is the smallest amongst the others and this also supported the optical and thermal measurements. The complete transformation of the austenite phase to martensite can be observed in the micrographs of the alloys except 1120 MPa pressure applied sample. The martensite plates observed in optical micrographs of all samples were identified as β1 martensite at room temperature. The thermal measurements showed that characteristic transformation temperatures of the samples increased by the increase of applied pressure value. So, the thermodynamic parameters such as enthalpy, entropy, hysterisis value and activation energies varied according to the characteristic transformation temperatures. The increase observed in the characteristic transformation temperatures with the increase of applied pressure is supported by literature. To calculate activation energy of austenite phase transformation, we plotted the curves of ln(ϕ /Tm2) vs. 1000/T and the calculated activation energies changed by the applied pressure as expected. In DTA measurements we observe the order-disorder phase transitions for the homogenous and pressure applied samples as mentioned in literature for Cu-based SMAs. For the homogenous sample, 280 MPa and 560 MPa pressure applied samples, we can not observe the B2 → L21 transition or Cu2MnAl type ordered transition. But 840 MPa and 1120 MPa pressure applied samples we observed B2 → L21 between 400–470 °C. This is mainly due to the effect of the size of applied pressure on samples. In this work we both observed the thermal and structural behavior of the order-disorder phase transition in Cu3Al-based SMAs by the effect of pressure. The order-disorder transition has an influence over the stability of bcc phase and over the martensitic transformation. Martensitic transformation properties are strongly connected with the degree of disorder retained by the samples and also ordering has the effect of stabilization on bcc phase of Cu-based SMAs. We observe two continuous transitions occurred during thermal measurements as L21 ↔ B2 and B2 ↔ A2; these two transitions combine with a single first order DO3 ↔ A2 transition.
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Acknowledgements This work is financially supported by FÜBAP, Project No: FF.15.14. Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at http://doi:10.1016/j.physb.2017.07.017.
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