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Corrosion behaviour of Cu-based shape memory alloys, diffusion bonded
Journal of Alloys and Compounds 387 (2005) 109–114
Corrosion behaviour of Cu-based shape memory alloys, diffusion bonded J.M. G´omez de Salazar∗ , A. Soria, M.I. Barrena Department of Material Science, Faculty of Chemistry, Complutense University of Madrid, 28040 Madrid, Spain Received 28 April 2004; received in revised form 1 June 2004; accepted 1 June 2004
Abstract The metallographic characterization and the corrosion resistance of diffusion bonded Cu–Al–Ag shape memory alloys was performed. An electrochemical study was performed, aimed at the evaluation of corrosion parameters both of the Cu-based alloy and the diffusion bonded alloy, using potentiodynamic techniques (ASTM standard G69-81). The polarisation curves were obtained in aerated marine water (3.5% NaCl solution). The result obtained indicates that bonded alloys exhibit low corrosion potential, inferior to that of Cu–Al–Ag alloys. The corrosion resistance of the bonded zone is worse than in the base material, due to the microsegregation occurring within the interphase bond, during the welding process. © 2004 Elsevier B.V. All rights reserved. Keywords: Metal; Scanning electron microscopy; Cu-based shape memory alloy; Precipitation; Corrosion
1. Introduction Copper-based alloys constitute a class of materials that have a number of applications due to their high conductivity, both thermal and electrical. They also exhibit good resistance to corrosion [1–3]. The aluminium bronzes (Cu–Al) present good strength and are also corrosion resistant. It has been shown that the addition of small amounts of silver to Cu–Al alloys increases hardness [4] and improves their resistance to stress corrosion in steam [5]. These copper-based alloys can also exhibit the shape memory effect within a certain range of compositions which have a disordered body center cubic structure (bcc), called -phase, which are stable at high temperature. They possess two successive ordering transitions during cooling treatment. Shape memory behaviour includes a number of different processes influencing the martensite–austenite reversible phase transformation [6–8]. Ms and Mf are the temperatures at which martensitic transformation begins and finishes, respectively (As and Af in the retransformation) [9,10]. These ∗
Corresponding author. Tel.: +34-91-3944351; fax: +34-91-3944357. E-mail address: [email protected] (J.M. G´omez de Salazar).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.06.028
alloys exhibit superelasticity when deformed by a determined stress, in a temperature range in which thermoelastic martensitic is formed. When the stress is removed the deformation disappears and the material spontaneously returns to the original phase [11,12]. The two way shape memory effect is not an intrisic property of shape memory alloys but can be obtained by several training methods, most of which are based on thermomechanical cycling [13,14]. The stabilized stress-induced martensite can be obtained by producing ␥ precipitates during the training procedure [15–17]. The contribution of each mechanism to the overall properties of shape memory vary according to percentages of alloying elements, heat treatments and stress applications [18–23]. Therefore, these alloys offer a wide range of applications due their high transformation temperatures. It is highly desirable to join shape memory alloys to make up more complex shapes. However due to difficulties in the joining process, their application is limited. Many papers have been concerned with investigating the welding characteristics of shape memory alloys [24–27] and the influence that the addition of the element has on the shape memory effect and corrosion resistance [28–31].
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Table 1 Chemical composition by weight of the Cu–Al–Ag shape memory alloys Ingot
1
2
3
4
5
6
7
8
9
10
11
Cu (%) Al (%) Ag (%)
87.32 11.99 0.69
86.53 12.01 1.46
85.64 12.85 1.51
85.73 13.33 0.94
86.44 12.65 0.91
86.39 13.23 0.38
87.28 12.39 0.33
85.62 12.22 2.16
86.67 12.68 0.65
85.51 11.48 3.01
84.88 11.24 3.88
Table 2 Chemical composition by weight of the Cusil ABA (filler metal) Elements wt.%
Ag 63.67
Cu 34.69
Ti 1.63
Al 0.004
Fe 0.003
Ni 0.002
The main objective of this work has been the metalographic characterization and the corrosion behaviour of the bonded Cu–Al–Ag shape memory alloy, which contain 84–87 wt.% Cu; 11–13 wt.% Al; and 0.3–4 wt.% Ag.
Au <0.001
Ca 0.0006
Ga 0.001
Mg <0.0001
Mn 0.003
Si 0.004
Pd <0.001
Table 3 Diffusion isotherm conditions Joint (◦ C)
Temperature Pressure (MPa)
1
2
3
4
5
750 0.5
800 0.5
850 0.5
750 1
800 1
2. Experimental procedure 2.1. Materials The base materials used in this investigation were 11 Cu–Al–Ag shape memory alloys with different chemical composition, which were obtained by a melting process. The chemical composition of these alloys was determined using Standard Test Methods for Chemical Analysis of Copper Alloys (ASTM 478-89a). These compositions are shown in Table 1. 2.2. Methods 2.2.1. Diffusion bonding process In order to determine the corrosion behavior of these bonded materials, a microstructural and corrosion evaluation of the Cu–Al–Ag alloy was carried out before and after the diffusion bonding process. The filler metals CuAgTi alloy
(Cusil ABA) were used. The filler metal was supplied in the form of a 0.5 mm sheet. Table 2 shows the chemical composition of this material. The dimensions of the Cu–Al–Ag specimens were 10 mm × 5 mm and the superficial roughness was of 0.1 m. The diffusion bonding of the Cu–Al–Ag alloys was carried out in a vacuum furnace (10−3 Pa) using a constants pressure of 0.5 MPa and 1 MPa. The diffusion isotherm conditions are shown in Table 3. The heating rate was 10 ◦ C/ min and the diffusion bonding time was of 30 min for all temperatures. 2.2.2. Electrochemical potentiodynamic measurement The electrochemical and potentiodynamic tests of the Cu–Al–Ag shape memory alloys (11 ingots) and the joints of the ingot-6 (five joints) were carried out using an AgCl/Ag reference electrode in a 3.5% NaCl aerated solution (pH 6.84). The cyclic polarisation curves were obtained using the same electrolyte and reference electrode. A graphite counter-
Fig. 1. The microstructure of Cu–Al–Ag (a) when the silver composition is high, grain (␣1 phase), eutectoid matrix (␣1 + ␥1 ) and silver rich precipitates in the grain boundaries (b) typical air cooled cast structure where the ␣1 phase is observed precipitated in grain boundary with Windmanst¨atten structure.
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111
Fig. 2. (a) The microstructure around the interface of bonding zone, (b) the Ag–Cu eutectic structure.
Fig. 3. The microstructure around the interface of bonding zone, when the joint is obtained at 825 ◦ C.
Fig. 4. Polarization curves for two different alloys in a 3.5% NaCl solution at room temperature.
electrode was used. The tests were carried out using a scanning rate of 1 mV/ s. Scans were initiated by lowering the corrosion potential of the specimen to a pre-set value of −230 mV (with respect to Ag/ClAg). In order to minimize the internal resistance drop in the solution, the reference electrode was positioned as closely as possible to the working electrode. For all tests, three working electrodes of each material were tested in order to ensure a good reproducibility of the results. Electrical connections were obtained by welding a copper wire using a condensater discharger and mounting in polyester resin. The surfaces were ground with emery paper, up to 600 grade, in order to obtain 0.1 m Ra roughness, and finally ultrasonically washed to degrease. 2.2.3. Microstructure and surface analysis Metallographic specimens of the bonded material were cut and embedded in a polyester resin, grinded, polished and etched (using the alcoholic solution of FeCl3 ) on the trans-
Fig. 5. SEM image of the surface of ingot-11 after holding the potential for 60 s at 300 mV.
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Table 4 Ecorr , RP , βcathodic , βanodic , icalc , iexp values obtained from the samples tested in aerated Ag/AgCl/3.5% NaCl solution Ingot
Ecorr (mV)
RP (m)
βcathodic (mV)
βanodic (mV)
icalc (A/cm2 )
iexp (A/cm2 )
1 2 3 4 5 6 7 8 9 10 11
−216 −223 −222 −244 −214 −222 −214 −229 −217 −210 −216
11150 7250 10200 6917 6517 4339 13066 8102 4950 9582 9051
68.029 71.257 73.790 79.486 63.823 69.675 61.154 92.000 68.070 72.033 92.003
−288.870 −173.157 −274.823 −200.537 −90.749 −232.709 −72.929 −112.620 −200.061 −378.03 −112.62
2.064 3.030 2.480 3.578 2.499 5.373 1.107 2.717 4.461 2.745 2.432
6.123 8.851 4.652 5.548 3.581 4.652 3.940 4.115 3.468 5.519 6.966
Table 5 Ecorr , RP , βcathodic , βanodic , icalc , iexp values obtained from the diffusion bonded ingot six samples, which were tested in aerated Ag/AgCl/3.5% NaCl solution Joint
Ecorr (mV)
RP (m)
βcathodic (mV)
βanodic (mV)
icalc (A/cm2 )
iexp (A/cm2 )
1 2 3 4 5
−263 −277 −274 −258 −276
6070 5236 8261 6844 7872
58.658 109.064 98.317 67.270 81.658
−188.545 −149.482 −402.745 −69.260 −233.236
3.205 5.236 4.159 2.168 3.340
2.815 7.568 8.643 4.723 6.636
verse axis of the welds. The loss of alloy elements and precipitated phases were studied. Scanning electron microscopy (SEM) and X-ray diffraction (XRD) studies were performed on the samples using a JEOL 6400 instrument and a Link EDX microanalyzer. After immersion the microstructure of the corroded surfaces were studied.
3. Results and discussion 3.1. Microstructure of Cu–Al–Ag alloys and diffusion-bonded joints The studied ingots all had a silver composition in the range 0.33–3.88. Therefore, the microstructure consists of a solid
solution of Al and Ag in Cu (␣1 phase), a eutectoid matrix (␣1 + ␥1 ) and silver rich precipitates in the grain boundaries. The latter occurring when the silver composition is high (Fig. 1a). At times, the ␣1 phase is observed to be precipitated into grain boundary with Widmannst¨atten structure. The microstructure of these ingots exhibited a typical air cooled cast structure (Fig. 1b). Fig. 2a shows the cross-sectional view of a macrostructure of the bonded specimens in which two bonding interfaces can be distinguished. The Ag–Cu eutectic (Fig. 2b) and Ti and Ag rich precipitates are presented in this interface. When the joints are formed at a higher temperature (775 ◦ C), the interface is more uniform, the Ag–Cu eutectic is not observed and the silver phases are precipitated into the matrix. Fine microinclusions of
Fig. 6. EDX and SEM image of the surface of a joint after electrochemical polarization test.
J.M. G´omez de Salazar et al. / Journal of Alloys and Compounds 387 (2005) 109–114
113
Fig. 7. EDX and SEM images of the surface of a bonding interface after electrochemical polarization test, which are identified as (a) TiAlAgCu phases, (b) AlAg phases and (c) free copper.
Ti–Al–Ag phase can appear both in the interface and in the copper matrix. If the joints are obtained at the 825 ◦ C condition both Ag and TiAlAg phases are not observed, only the ␣1 phase is found in the interface (Fig. 3).
3.2. Corrosion behaviour of Cu–Al–Ag alloys and diffusion-bonded joints Fig. 4 shows the typical polarization curves for two different alloys in a 3.5% NaCl solution at room temperature.
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The silver content of these specimens was 0.3% (ingot-7) and 3.9% (ingot-11). It would be expected that a higher silver content in the alloy would shift the corrosion potential to more positive values. However, this is not observed. For the ingot-11, during the initial positive potential an anodic peak A (Fig. 4) appeared at 130 mV, which is related the formation of a copper oxide layer. Increasing the potential revealed that the measured current did not drop to zero indicating that the passive film was not formed allowing copper to continue to dissolve. This is not observed for the ingot-7, due to a high presence of aluminium. Fig. 5 shows the surface of ingot-11 after holding the potential for 60 s at 300 mV, corresponding to the end of peak A. The eutectoid ␥1 phase was more susceptible to that than the ␣1 phase. This is due to the higher corrosion potential of the copper phase.The cyclic polarisation curves in aerated 3.5% NaCl solution, and polarisation resistance curves were obtained for both the ingots and the joints. Table 4 shows the obtained results from these curves for ingots. From this table, we can deduce that the influence of the composition of the ingot on the Ecorr is not relevant. However, we can deduce that the joints present a higher Ecorr than the ingots (Table 5). This fact is a consequence of the microstructural changes that take place during thermal cycles. In addition, the composition of the used filler and the existence of various different phases supports the formation of local cells, promoting the various possible redox reactions. The corrosion mechanism in the joints takes place mainly through the interface filler-base material. This fact is a consequence of the different chemical composition and the different corrosion behaviour between theses components. As shown on Fig. 6, a corrosion residuum can be observed on the attack surface, formed by a mixture of complex oxides. When the joints have higher quality phases such as TiAlAgCu, AlAg and free copper, it is not unusual to detect intermetallic phases in the interface (Fig. 7a–c, respectively), due to the interfacial reaction between the filler metal and the base metal during the diffusion bonding process. The corrosion behaviour of the phases in the joint is similar to that of the ingots. A preferential attack was associated with the ␥1 phase, which is anodic, from the eutectoid structure. However, the silver precipitates and the ␣1 phase (also of the eutectoid structure) are cathodic areas and hence did not suffer attack.
4. Conclusions 1. Joints by diffusion bonding processes have been obtained of a high microstructural quality at low pressure for temperatures of 775 and 800 ◦ C. 2. Electrochemical and potentiodynamic measurements combined with SEM and EDX allow us determine the effect of the microstructural changes provoked both by diffusion thermal cycles and the different composition, on the electrochemical corrosion behaviour of the Cu–Al–Ag to be understood.
3. The corrosion behaviour of the Cu–Al–Ag is very similar for all the compositions tested. The Ecorr values are between −210 mV and −245 mV. 4. The corrosion behaviour of the joints is affected by the filler material composition and the intermetallic precipitates present in the interface. The Ecorr values in the joints are between −260 mV and −280 mV. Acknowledgment The authors are grateful to CICYT MAT2000-0862-C0301 and MAT2003-05004-C02-01 for their financial support. References [1] E.A. Ashour, B.G. Ateya, Electrochim. Acta 42 (1997) 243–250. [2] S. Longauer, P. Makroczy, G. Jan´ak, M. Longauerov´a, Mater. Sci. Eng. A273–A275 (1999) 415–419. [3] R. Zenging, M. Ceylan, Thermochim. Acta (2004) 414. [4] A.T. Adorno, M.R. Guerreiro, A.V. Benedetti, J. Alloys Compd. 268 (1998) 122–129. [5] S.S. Rosatto, P.L. Cabot, P.T.A. Sumodjo, A.V. Benedetti, Thermochim. Acta 46 (2001) 1043–1051. [6] R. Zenging, M. Ceylan, Mater. Lett. (2003) 58. [7] A. Planes, L.L. Ma˜nosa, Solid State Phys. 55 (2001) 159–165. [8] F. Dagdelen, T. Gokhan, A. Aydogdu, Y. Aydogdu, O. Adig¨uzel, Mater. Lett. 57 (2003) 1079–1085. [9] J.M. Guilemany, J. F´andez, R. Franch, Intermetallics 6 (1998) 15–19. [10] X.M. Zhang, J.M. Guilemany, J. Fern´andez, M. Liu, L. Liu, Intermetallics 8 (2000) 703–707. [11] J. Pe˜na, F.J. Gil, J.M. Guilemany, Acta Mater. 50 (2002) 3115–3124. [12] J. Fern´andez, X.M. Zhang, J.M. Guilemany, J. Mater. Process. Tech. 139 (2003) 117–119. [13] M.A. Morris, Acta Mater. 40 (1992) 1573–1579. [14] R. Stalmans, D. Van Humbeeck, J. Mater. Trans. JIM 33 (1992) 284–289. [15] D.O. Roqueta, F.C. Lovey, M. Sade, Scr. Mater. 36 (1997) 385–391. [16] X.M. Zhang, M. Liu, J. Fern´andez, J.M. Guilemany, Mater. Design 21 (2000) 557–559. [17] J. Pons, R. Portier, Acta Mater. 45 (1997) 2109–2120. [18] J. Lelatko, H. Morawiec, Mater. Chem. Phys. 81 (2003) 472–475. [19] S.N. Balo, M. Ceylan, J. Mater. Process. Tech. 124 (2002) 200–208. [20] S.N. Balo, M. Ceylan, M. Aksoy, Mater. Sci. Eng. A311 (2001) 151–156. [21] V. Nov´ak, P. Sittner, D. Vokoun, N. Z´arubov´a, Mater. Sci. Eng. A273–A275 (1999) 280–285. [22] Y. Sutou, R. Kainuma, K. Ishida, Mater. Sci. Eng. A273–A275 (1999) 375–379. [23] F.J. Gil, J.M. Guilemany, Intermetallics 7 (1999) 699–704. [24] J.M. G´omez de Salazar, F.J. M´endez, A. Ure˜na, J.M. Guilemany, B.G. Mellor, Scr Mater. 37 (6) (1997) 861–867. [25] H.C. Lin, K.M. Lin, Y.C. Chuang, T.S. Chou, J. Alloys Compd. 306 (2000) 186–192. [26] O. Soderberg, X.W. Liu, P.G. Yakovenko, K. Ullakko, V.K. Lindroos, Mater. Sci. Eng. A273–A275 (1999) 543–548. [27] A. Ikai, J. Intell. Mater. Syst. Struct. 7 (1999) 646–651. [28] H.C. Lin, K.M. Lin, C.S. Lin, T.M. Ouyang, Corros. Sci. 44 (2002) 2013–2026. [29] G. Rondelli, B. Vicentini, Biomaterials 23 (2002) 639–644. [30] O. S¨oderberg, X.W. Liu, P.G. Yakovenko, K. Ullakko, V.K. Lindroos, Mater. Sci. Eng. A273–A275 (1999) 543–548. [31] D. Starosvetsky, I. Gotman, Biomaterials 22 (2001) 1853–1859.
Corrosion behaviour of Cu-based shape memory alloys, diffusion bonded J.M. G´omez de Salazar∗ , A. Soria, M.I. Barrena Department of Material Science, Faculty of Chemistry, Complutense University of Madrid, 28040 Madrid, Spain Received 28 April 2004; received in revised form 1 June 2004; accepted 1 June 2004
Abstract The metallographic characterization and the corrosion resistance of diffusion bonded Cu–Al–Ag shape memory alloys was performed. An electrochemical study was performed, aimed at the evaluation of corrosion parameters both of the Cu-based alloy and the diffusion bonded alloy, using potentiodynamic techniques (ASTM standard G69-81). The polarisation curves were obtained in aerated marine water (3.5% NaCl solution). The result obtained indicates that bonded alloys exhibit low corrosion potential, inferior to that of Cu–Al–Ag alloys. The corrosion resistance of the bonded zone is worse than in the base material, due to the microsegregation occurring within the interphase bond, during the welding process. © 2004 Elsevier B.V. All rights reserved. Keywords: Metal; Scanning electron microscopy; Cu-based shape memory alloy; Precipitation; Corrosion
1. Introduction Copper-based alloys constitute a class of materials that have a number of applications due to their high conductivity, both thermal and electrical. They also exhibit good resistance to corrosion [1–3]. The aluminium bronzes (Cu–Al) present good strength and are also corrosion resistant. It has been shown that the addition of small amounts of silver to Cu–Al alloys increases hardness [4] and improves their resistance to stress corrosion in steam [5]. These copper-based alloys can also exhibit the shape memory effect within a certain range of compositions which have a disordered body center cubic structure (bcc), called -phase, which are stable at high temperature. They possess two successive ordering transitions during cooling treatment. Shape memory behaviour includes a number of different processes influencing the martensite–austenite reversible phase transformation [6–8]. Ms and Mf are the temperatures at which martensitic transformation begins and finishes, respectively (As and Af in the retransformation) [9,10]. These ∗
Corresponding author. Tel.: +34-91-3944351; fax: +34-91-3944357. E-mail address: [email protected] (J.M. G´omez de Salazar).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.06.028
alloys exhibit superelasticity when deformed by a determined stress, in a temperature range in which thermoelastic martensitic is formed. When the stress is removed the deformation disappears and the material spontaneously returns to the original phase [11,12]. The two way shape memory effect is not an intrisic property of shape memory alloys but can be obtained by several training methods, most of which are based on thermomechanical cycling [13,14]. The stabilized stress-induced martensite can be obtained by producing ␥ precipitates during the training procedure [15–17]. The contribution of each mechanism to the overall properties of shape memory vary according to percentages of alloying elements, heat treatments and stress applications [18–23]. Therefore, these alloys offer a wide range of applications due their high transformation temperatures. It is highly desirable to join shape memory alloys to make up more complex shapes. However due to difficulties in the joining process, their application is limited. Many papers have been concerned with investigating the welding characteristics of shape memory alloys [24–27] and the influence that the addition of the element has on the shape memory effect and corrosion resistance [28–31].
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Table 1 Chemical composition by weight of the Cu–Al–Ag shape memory alloys Ingot
1
2
3
4
5
6
7
8
9
10
11
Cu (%) Al (%) Ag (%)
87.32 11.99 0.69
86.53 12.01 1.46
85.64 12.85 1.51
85.73 13.33 0.94
86.44 12.65 0.91
86.39 13.23 0.38
87.28 12.39 0.33
85.62 12.22 2.16
86.67 12.68 0.65
85.51 11.48 3.01
84.88 11.24 3.88
Table 2 Chemical composition by weight of the Cusil ABA (filler metal) Elements wt.%
Ag 63.67
Cu 34.69
Ti 1.63
Al 0.004
Fe 0.003
Ni 0.002
The main objective of this work has been the metalographic characterization and the corrosion behaviour of the bonded Cu–Al–Ag shape memory alloy, which contain 84–87 wt.% Cu; 11–13 wt.% Al; and 0.3–4 wt.% Ag.
Au <0.001
Ca 0.0006
Ga 0.001
Mg <0.0001
Mn 0.003
Si 0.004
Pd <0.001
Table 3 Diffusion isotherm conditions Joint (◦ C)
Temperature Pressure (MPa)
1
2
3
4
5
750 0.5
800 0.5
850 0.5
750 1
800 1
2. Experimental procedure 2.1. Materials The base materials used in this investigation were 11 Cu–Al–Ag shape memory alloys with different chemical composition, which were obtained by a melting process. The chemical composition of these alloys was determined using Standard Test Methods for Chemical Analysis of Copper Alloys (ASTM 478-89a). These compositions are shown in Table 1. 2.2. Methods 2.2.1. Diffusion bonding process In order to determine the corrosion behavior of these bonded materials, a microstructural and corrosion evaluation of the Cu–Al–Ag alloy was carried out before and after the diffusion bonding process. The filler metals CuAgTi alloy
(Cusil ABA) were used. The filler metal was supplied in the form of a 0.5 mm sheet. Table 2 shows the chemical composition of this material. The dimensions of the Cu–Al–Ag specimens were 10 mm × 5 mm and the superficial roughness was of 0.1 m. The diffusion bonding of the Cu–Al–Ag alloys was carried out in a vacuum furnace (10−3 Pa) using a constants pressure of 0.5 MPa and 1 MPa. The diffusion isotherm conditions are shown in Table 3. The heating rate was 10 ◦ C/ min and the diffusion bonding time was of 30 min for all temperatures. 2.2.2. Electrochemical potentiodynamic measurement The electrochemical and potentiodynamic tests of the Cu–Al–Ag shape memory alloys (11 ingots) and the joints of the ingot-6 (five joints) were carried out using an AgCl/Ag reference electrode in a 3.5% NaCl aerated solution (pH 6.84). The cyclic polarisation curves were obtained using the same electrolyte and reference electrode. A graphite counter-
Fig. 1. The microstructure of Cu–Al–Ag (a) when the silver composition is high, grain (␣1 phase), eutectoid matrix (␣1 + ␥1 ) and silver rich precipitates in the grain boundaries (b) typical air cooled cast structure where the ␣1 phase is observed precipitated in grain boundary with Windmanst¨atten structure.
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111
Fig. 2. (a) The microstructure around the interface of bonding zone, (b) the Ag–Cu eutectic structure.
Fig. 3. The microstructure around the interface of bonding zone, when the joint is obtained at 825 ◦ C.
Fig. 4. Polarization curves for two different alloys in a 3.5% NaCl solution at room temperature.
electrode was used. The tests were carried out using a scanning rate of 1 mV/ s. Scans were initiated by lowering the corrosion potential of the specimen to a pre-set value of −230 mV (with respect to Ag/ClAg). In order to minimize the internal resistance drop in the solution, the reference electrode was positioned as closely as possible to the working electrode. For all tests, three working electrodes of each material were tested in order to ensure a good reproducibility of the results. Electrical connections were obtained by welding a copper wire using a condensater discharger and mounting in polyester resin. The surfaces were ground with emery paper, up to 600 grade, in order to obtain 0.1 m Ra roughness, and finally ultrasonically washed to degrease. 2.2.3. Microstructure and surface analysis Metallographic specimens of the bonded material were cut and embedded in a polyester resin, grinded, polished and etched (using the alcoholic solution of FeCl3 ) on the trans-
Fig. 5. SEM image of the surface of ingot-11 after holding the potential for 60 s at 300 mV.
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Table 4 Ecorr , RP , βcathodic , βanodic , icalc , iexp values obtained from the samples tested in aerated Ag/AgCl/3.5% NaCl solution Ingot
Ecorr (mV)
RP (m)
βcathodic (mV)
βanodic (mV)
icalc (A/cm2 )
iexp (A/cm2 )
1 2 3 4 5 6 7 8 9 10 11
−216 −223 −222 −244 −214 −222 −214 −229 −217 −210 −216
11150 7250 10200 6917 6517 4339 13066 8102 4950 9582 9051
68.029 71.257 73.790 79.486 63.823 69.675 61.154 92.000 68.070 72.033 92.003
−288.870 −173.157 −274.823 −200.537 −90.749 −232.709 −72.929 −112.620 −200.061 −378.03 −112.62
2.064 3.030 2.480 3.578 2.499 5.373 1.107 2.717 4.461 2.745 2.432
6.123 8.851 4.652 5.548 3.581 4.652 3.940 4.115 3.468 5.519 6.966
Table 5 Ecorr , RP , βcathodic , βanodic , icalc , iexp values obtained from the diffusion bonded ingot six samples, which were tested in aerated Ag/AgCl/3.5% NaCl solution Joint
Ecorr (mV)
RP (m)
βcathodic (mV)
βanodic (mV)
icalc (A/cm2 )
iexp (A/cm2 )
1 2 3 4 5
−263 −277 −274 −258 −276
6070 5236 8261 6844 7872
58.658 109.064 98.317 67.270 81.658
−188.545 −149.482 −402.745 −69.260 −233.236
3.205 5.236 4.159 2.168 3.340
2.815 7.568 8.643 4.723 6.636
verse axis of the welds. The loss of alloy elements and precipitated phases were studied. Scanning electron microscopy (SEM) and X-ray diffraction (XRD) studies were performed on the samples using a JEOL 6400 instrument and a Link EDX microanalyzer. After immersion the microstructure of the corroded surfaces were studied.
3. Results and discussion 3.1. Microstructure of Cu–Al–Ag alloys and diffusion-bonded joints The studied ingots all had a silver composition in the range 0.33–3.88. Therefore, the microstructure consists of a solid
solution of Al and Ag in Cu (␣1 phase), a eutectoid matrix (␣1 + ␥1 ) and silver rich precipitates in the grain boundaries. The latter occurring when the silver composition is high (Fig. 1a). At times, the ␣1 phase is observed to be precipitated into grain boundary with Widmannst¨atten structure. The microstructure of these ingots exhibited a typical air cooled cast structure (Fig. 1b). Fig. 2a shows the cross-sectional view of a macrostructure of the bonded specimens in which two bonding interfaces can be distinguished. The Ag–Cu eutectic (Fig. 2b) and Ti and Ag rich precipitates are presented in this interface. When the joints are formed at a higher temperature (775 ◦ C), the interface is more uniform, the Ag–Cu eutectic is not observed and the silver phases are precipitated into the matrix. Fine microinclusions of
Fig. 6. EDX and SEM image of the surface of a joint after electrochemical polarization test.
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Fig. 7. EDX and SEM images of the surface of a bonding interface after electrochemical polarization test, which are identified as (a) TiAlAgCu phases, (b) AlAg phases and (c) free copper.
Ti–Al–Ag phase can appear both in the interface and in the copper matrix. If the joints are obtained at the 825 ◦ C condition both Ag and TiAlAg phases are not observed, only the ␣1 phase is found in the interface (Fig. 3).
3.2. Corrosion behaviour of Cu–Al–Ag alloys and diffusion-bonded joints Fig. 4 shows the typical polarization curves for two different alloys in a 3.5% NaCl solution at room temperature.
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The silver content of these specimens was 0.3% (ingot-7) and 3.9% (ingot-11). It would be expected that a higher silver content in the alloy would shift the corrosion potential to more positive values. However, this is not observed. For the ingot-11, during the initial positive potential an anodic peak A (Fig. 4) appeared at 130 mV, which is related the formation of a copper oxide layer. Increasing the potential revealed that the measured current did not drop to zero indicating that the passive film was not formed allowing copper to continue to dissolve. This is not observed for the ingot-7, due to a high presence of aluminium. Fig. 5 shows the surface of ingot-11 after holding the potential for 60 s at 300 mV, corresponding to the end of peak A. The eutectoid ␥1 phase was more susceptible to that than the ␣1 phase. This is due to the higher corrosion potential of the copper phase.The cyclic polarisation curves in aerated 3.5% NaCl solution, and polarisation resistance curves were obtained for both the ingots and the joints. Table 4 shows the obtained results from these curves for ingots. From this table, we can deduce that the influence of the composition of the ingot on the Ecorr is not relevant. However, we can deduce that the joints present a higher Ecorr than the ingots (Table 5). This fact is a consequence of the microstructural changes that take place during thermal cycles. In addition, the composition of the used filler and the existence of various different phases supports the formation of local cells, promoting the various possible redox reactions. The corrosion mechanism in the joints takes place mainly through the interface filler-base material. This fact is a consequence of the different chemical composition and the different corrosion behaviour between theses components. As shown on Fig. 6, a corrosion residuum can be observed on the attack surface, formed by a mixture of complex oxides. When the joints have higher quality phases such as TiAlAgCu, AlAg and free copper, it is not unusual to detect intermetallic phases in the interface (Fig. 7a–c, respectively), due to the interfacial reaction between the filler metal and the base metal during the diffusion bonding process. The corrosion behaviour of the phases in the joint is similar to that of the ingots. A preferential attack was associated with the ␥1 phase, which is anodic, from the eutectoid structure. However, the silver precipitates and the ␣1 phase (also of the eutectoid structure) are cathodic areas and hence did not suffer attack.
4. Conclusions 1. Joints by diffusion bonding processes have been obtained of a high microstructural quality at low pressure for temperatures of 775 and 800 ◦ C. 2. Electrochemical and potentiodynamic measurements combined with SEM and EDX allow us determine the effect of the microstructural changes provoked both by diffusion thermal cycles and the different composition, on the electrochemical corrosion behaviour of the Cu–Al–Ag to be understood.
3. The corrosion behaviour of the Cu–Al–Ag is very similar for all the compositions tested. The Ecorr values are between −210 mV and −245 mV. 4. The corrosion behaviour of the joints is affected by the filler material composition and the intermetallic precipitates present in the interface. The Ecorr values in the joints are between −260 mV and −280 mV. Acknowledgment The authors are grateful to CICYT MAT2000-0862-C0301 and MAT2003-05004-C02-01 for their financial support. References [1] E.A. Ashour, B.G. Ateya, Electrochim. Acta 42 (1997) 243–250. [2] S. Longauer, P. Makroczy, G. Jan´ak, M. Longauerov´a, Mater. Sci. Eng. A273–A275 (1999) 415–419. [3] R. Zenging, M. Ceylan, Thermochim. Acta (2004) 414. [4] A.T. Adorno, M.R. Guerreiro, A.V. Benedetti, J. Alloys Compd. 268 (1998) 122–129. [5] S.S. Rosatto, P.L. Cabot, P.T.A. Sumodjo, A.V. Benedetti, Thermochim. Acta 46 (2001) 1043–1051. [6] R. Zenging, M. Ceylan, Mater. Lett. (2003) 58. [7] A. Planes, L.L. Ma˜nosa, Solid State Phys. 55 (2001) 159–165. [8] F. Dagdelen, T. Gokhan, A. Aydogdu, Y. Aydogdu, O. Adig¨uzel, Mater. Lett. 57 (2003) 1079–1085. [9] J.M. Guilemany, J. F´andez, R. Franch, Intermetallics 6 (1998) 15–19. [10] X.M. Zhang, J.M. Guilemany, J. Fern´andez, M. Liu, L. Liu, Intermetallics 8 (2000) 703–707. [11] J. Pe˜na, F.J. Gil, J.M. Guilemany, Acta Mater. 50 (2002) 3115–3124. [12] J. Fern´andez, X.M. Zhang, J.M. Guilemany, J. Mater. Process. Tech. 139 (2003) 117–119. [13] M.A. Morris, Acta Mater. 40 (1992) 1573–1579. [14] R. Stalmans, D. Van Humbeeck, J. Mater. Trans. JIM 33 (1992) 284–289. [15] D.O. Roqueta, F.C. Lovey, M. Sade, Scr. Mater. 36 (1997) 385–391. [16] X.M. Zhang, M. Liu, J. Fern´andez, J.M. Guilemany, Mater. Design 21 (2000) 557–559. [17] J. Pons, R. Portier, Acta Mater. 45 (1997) 2109–2120. [18] J. Lelatko, H. Morawiec, Mater. Chem. Phys. 81 (2003) 472–475. [19] S.N. Balo, M. Ceylan, J. Mater. Process. Tech. 124 (2002) 200–208. [20] S.N. Balo, M. Ceylan, M. Aksoy, Mater. Sci. Eng. A311 (2001) 151–156. [21] V. Nov´ak, P. Sittner, D. Vokoun, N. Z´arubov´a, Mater. Sci. Eng. A273–A275 (1999) 280–285. [22] Y. Sutou, R. Kainuma, K. Ishida, Mater. Sci. Eng. A273–A275 (1999) 375–379. [23] F.J. Gil, J.M. Guilemany, Intermetallics 7 (1999) 699–704. [24] J.M. G´omez de Salazar, F.J. M´endez, A. Ure˜na, J.M. Guilemany, B.G. Mellor, Scr Mater. 37 (6) (1997) 861–867. [25] H.C. Lin, K.M. Lin, Y.C. Chuang, T.S. Chou, J. Alloys Compd. 306 (2000) 186–192. [26] O. Soderberg, X.W. Liu, P.G. Yakovenko, K. Ullakko, V.K. Lindroos, Mater. Sci. Eng. A273–A275 (1999) 543–548. [27] A. Ikai, J. Intell. Mater. Syst. Struct. 7 (1999) 646–651. [28] H.C. Lin, K.M. Lin, C.S. Lin, T.M. Ouyang, Corros. Sci. 44 (2002) 2013–2026. [29] G. Rondelli, B. Vicentini, Biomaterials 23 (2002) 639–644. [30] O. S¨oderberg, X.W. Liu, P.G. Yakovenko, K. Ullakko, V.K. Lindroos, Mater. Sci. Eng. A273–A275 (1999) 543–548. [31] D. Starosvetsky, I. Gotman, Biomaterials 22 (2001) 1853–1859.