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A new ultrahigh strength Cu–Ni–Si alloy
Intermetallics 42 (2013) 77e84
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Short communication
A new ultrahigh strength CueNieSi alloy Q. Lei a, Z. Li a, b, *, T. Xiao a, Y. Pang a, Z.Q. Xiang a, W.T. Qiu a, Z. Xiao c a
School of Materials Science and Engineering, Central South University, Changsha 410083, China State Key Laboratory of Powder Metallurgy, Changsha 410083, China c Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 4100083, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 February 2013 Received in revised form 2 May 2013 Accepted 19 May 2013 Available online
A new ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy has been developed by alloying design and thermal mechanical treatment. Electrical conductivity, mechanical properties of the designed alloy such as hardness, tensile strength, yield strength, elongation and anti-stress relaxation resistance were tested. The mechanical properties and electrical properties of designed alloy were comparable to those of CueBe alloys, and its anti-stress relaxation resistance even was better than that of CueBe alloys at evaluated temperature. Microstructure observation revealed that b-Ni3Si phase precipitated in the initial stages ageing process, with further increasing the ageing time, b-Ni3Si, and d-Ni2Si phase precipitates appeared and contributed to the ultrahigh strength by Orowan strengthening. The satellites spots around diffraction spots of Cu matrix symmetrically in [112]Cu zone crystal axis have been determined, resulting that satellites spots are from the electron diffraction of d-Ni2Si precipitates. The crystal orientation relationship between matrix and precipitates is that: ð111Þcu ==ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d . Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: B. Alloy design B. Precipitates C. Melting C. Thermomechanical treatment D. Microstructure G. Automotive uses
1. Introduction High strength materials play significant role on the automaton industry, automobile industry, electrical and electronics industry. Copper alloys are widely used due to their high strength and good electrical conductivity [1,2]. CueBe alloys is the most intensive used ones among present Cu-based elastic materials, owing to their good mechanical and electrical properties [3]. However, the toxicity of beryllium element in CueBe alloy was harmful to human body in the production process [4], and anti-stress relaxation properties of CueBe alloys are poor as used at elevated temperature. A lot of attempts have been made by other researchers to developed novel copper alloys for replacing the CueBe alloys before, such as CueNie Sn, CueNieZn and CueNieAl system alloys [5e7]. These alloys exhibit ultrahigh strength (>1000 MPa). However, their electrical conductivities (<15 %IACS) are less than those of CueBe alloys (w25 %IACS). Encouragingly, CueNieSi alloys have aroused considerable interest because of their high strength and good electrical conductivity [8e12]. The high strength of CueNieSi alloys is attributed to the nano-scale precipitates during precipitation [13e16]. In our * Corresponding author. School of Materials Science and Engineering, Central South University, Changsha 410083, China. Tel.: þ86 731 88830264; fax: þ86 731 88876692. E-mail address: [email protected] (Z. Li). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.05.013
previous investigations, strength could be effectively enhanced through increasing the content of Ni and Si. Cu-8.0Ni-1.8Si system alloys were designed and they showed good mechanical properties [17e19], however, some precipitates composed by Ni and Si could not be re-dissolved back to matrix even at very high solution temperature [17], as a result that it is hard to achieve the desired performance. Therefore, it is a key issue to lower the content of Ni and Si but not to decrease the strength of the alloy. One of the effective ways is addition of other suitable alloying elements [20]. Al and Mg could affect the precipitation kinetics and increased the formation rate of precipitates [9,21]. Effects of additions of Mg [21] and Cr [22] on the microstructure and property of CueNieSi alloys have been reported, showing that they could effectively improve the strength of alloy. Therefore, Al, Mg and Cr elements were added in the designed CueNieSi alloys. Copper base alloys with high strength and good electrical conductivity are generally solution treated by high temperature quenching and subsequent ageing. Nano-scale precipitates formed in copper matrix through these processes, as a result that the alloy is strengthened and its electrical conductivity is increased. Studies on the mechanical properties and microstructure of ultrahigh strength copperenickelesilicon (CueNieSi) alloys have been carried out thus far [11,16e19], and some interesting results on precipitation have been obtained. However, the appropriate thermomechanical treatment in the Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1
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Cr alloy has not been investigated yet. Not only the chemical composition but also the process condition play important role in the comprehensive properties. Therefore, thermo-mechanical treatment of the ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloys should be investigated, by which an optimum process condition for improving the properties can be obtained. In the present work, a novel copper alloy with nominal composition of Cu-6.0% Ni-1.0% Si-0.5% Al-0.15% Mg-0.1% Cr (weight percentage) has been designed and its microstructure evolution and properties have been investigated.
time of 15 s. Electrical conductivity was measured using a doublearm electrical bridge device. Tensile tests were performed using an Instron 8019 tester type machine with a constant strain rate of 103 s1. Microstructure observation was operated using a JEM2100F transmission electron microscope with operation voltage of 200 kV and point resolution of 0.19 nm. Thin foils for TEM and HRTEM observations were prepared by ion beam milling. Stress relaxation experiment was carried out on a home-made device designed according to ASTM E328-02(2008) [23]. 3. Results and discussion
2. Experimental details 3.1. Thermo-mechanical treatment Firstly, pure Cu, pure Ni, and pure Cr blocks were melted in a medium-frequency induction furnace. Then, Cu-14 wt.% Si, Cu20 wt.% Mg master alloys and pure Al blocks were added into the melt. Melting and casting were carried out in N2 atmosphere to prevent the melt and ingot from oxidization. After removing the surface oxides, a 30 mm thickness ingot was homogenization treated at 940 C for 4 h, followed by hot rolled at 850 C with total deformation of 80%, thickness from 30 mm to 6 mm. Supersaturated solid solution is the basis of precipitation hardening, therefore, solution treatment at 950 C, 960 C, 970 C and 980 C for 6 h were performed on the hot-rolling specimens to determine appropriate solution treatment temperature, respectively. Then the specimens were cold rolled to decrease the thickness from 6 mm to 3 mm, finally ageing treated at 450 C and 500 C for different time. Metallographic observation was carried out using a Leica EC3 optical microscope. Hardness tests were operated using an HV-5 typical Vickers hardness tester, with a load of 3 kg and holding
The initial microstructure of as-cast ingot is shown in Fig. 1(a). Developed dendrite grains are observed. Meanwhile, some nonequilibrium precipitates appeared between the dendrite arms (Fig. 1(a)), which are harmful to subsequent thermal-mechanical heat treatment. To eliminate the dendrite structures, homogenization treatment at 940 C for 4 h was conducted, by which the dendrite structures were eliminated (Fig. 1(b)). Backscatter electron image and elemental mapping images of the homogenization specimens were taken. The elemental mapping images of Fig. 1(c) is shown in Fig. 1(d)e(i), which indicates that the particles were rich in Ni and Si but poor in Cu. Al, Mg and Cr distributed inside the grains evenly. The elemental line distribution image also agreed well with the elemental mapping results (Fig. 1(j)). Fig. 2(a) shows microstructures of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after hot rolling at 850 C with a total deformation of 80%. After hot rolling deformation, initial large grains were
Fig. 1. Optical microstructure, SEM micrograph and the elemental mapping images of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy. (a) As cast; (b) Homogenized at 940 C for 4 h; (c) Backscatter electron image; Elemental mapping images: (d) Cu; (e) Ni; (f) Si, (g) Al; (h) Mg; (i) Cr; (j) Line distribution of Cu, Ni, Si and Al elements.
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Fig. 2. The microstructure of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy as hot rolled at 850 C and solution treatment at different temperatures for 6 h. Hot rolling: (a) Rolling surface; (b) Rolling side; Solution treatment at different temperature: (c) 950 C; (d) 960 C; (e) 970 C; (f) 980 C.
replaced by fine recrystallized grains. Dynamic recrystallization finished completely, annealing twins appeared in the rolling surface (as shown in Fig. 2(a)) and some recrystallization grains were elongated along the rolling direction in the rolling slide (Fig. 2(b)). Fig. 2(c)e(f) shows microstructures of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after solution treatments. The content of residual intermetallic compounds decreased with increase of temperature. As the specimens were solution treated at 950 C or 960 C for 6 h, a lot of intermetallic compounds still can be observed in the matrix (Fig. 2(c) and (d)). While, for the specimens solution treatment at 970 C for 6 h, intermetallic compounds re-dissolved into the matrix completely, indicating that supersaturated solid solution formed (Fig. 2(e)). Further increase the temperature to 980 C, the grains coarsened seriously (Fig. 2(f)). Based on the microstructure observations, intermetallic compounds and grain size, the appropriate solution treatment was determined, 970 C 6 h.
3.2. Hardness and electrical conductivity After solution treatment at 970 C for 6 h, the solution treated bulk materials was cold rolled with total deformation of 50%. Then it was cut into standard specimens for hardness and electrical conductivity, respectively. Fig. 3 shows the variation of hardness and electrical conductivity of alloy with ageing time. With ageing time, hardness of specimens increased rapidly and reached peak value, then decreased. This decrease of hardness is a thermally activated process [24]. When specimens were aged for more than 30 min, hardness of specimens decreased more quickly as aged at 500 C, than that of specimens aged at 450 C (Fig. 3(a)). When the ageing temperature was 500 C, the hardness decreased due to the coarsening and dissolution of the precipitates. Electrical conductivity of specimens increased rapidly in the initial ageing process, then increased slowly (Fig. 3(b)). With the same ageing time, the
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Fig. 3. Properties of designed alloy. (a) Hardness; (b) Electrical conductivity; (c) Stressestrain curve of designed alloy aged at 450 C for 60 min.
electrical conductivities of the specimens aged at 500 C were higher than those of specimens aged at 450 C. Considering the combination of properties, the more appropriate ageing method was 450 C 60 min for the designed alloy. Under this ageing treatment, hardness of the alloy was 341 HV and the corresponding electrical conductivity was 26.5% IACS. Electrical resistivity is sensitive to a number of microstructural factors including vacancy concentration, solute concentration in the matrix, precipitate or cluster size, and precipitate volume fraction [25]. The results from tensile test of specimens aged at 450 C for 60 min (Fig. 3(c)), indicating that the tensile strength, yield strength and elongation of designed alloy were 1090 MPa, 940 MPa and 3.5%, respectively. 3.3. Precipitation Microstructures and selected-area diffraction patterns (SADPs) of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 15 min are shown in Fig. 4. Some nano-scale precipitating particles can be seen (Fig. 4(a)). The mean dimension of particles was w8 nm, as shown in Fig. 4(b). The SADPs are shown in Fig. 4(c) and (d), some additional diffraction spots different from the spots from the electron diffraction of Cu matrix could be detected. The indexing results of Fig. 4(c) and (d) illustrate that additional spots are derived from the electron diffraction of bNi3Si phase precipitates. Cu matrix has a face central cubic structure, a ¼ 0.3615 nm, b-Ni3Si phase precipitate has a prime cubic structure, a ¼ 0.3351 nm [26,27]. Based on the analysis results of SADPs in Fig. 4, the crystal orientation relationship between matrix and b-Ni3Si precipitate can be indexed as follows: ð001Þcu == ð001Þb ; ½100cu ==½100b ; ð110Þcu ==ð220Þb ; ½112cu ==½112b . Fig. 5 showed microstructures and SADPs of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy aged at 450 C for 60 min. The precipitates coarsened slightly and the mean size of precipitates was w10 nm (Fig. 5(a) and (b)). Except some new faint diffraction spots from
other precipitates, the diffraction spots and their indexation results were similar to Fig. 4. In order to detect all precipitates and determine the crystal orientation relationship using SADPs with beam along [112]Cu, the ageing time was further increased to 480 min Fig. 6 shows the brightfield micrograph, central dark-field micrograph, HRTEM micrograph and SADPs of specimens as aged at 450 C for 480 min. The precipitates obviously coarsened compared to that ageing treated at 450 C for 60 min d-Ni2Si precipitates with mutual perpendicular directions formed and determined, as shown in the micrographs images (Fig. 6(a),(b),(d) and (f)). The index results in Fig. 6(c) and (e) illustrated that two kinds of precipitate phases of b-Ni3Si and d-Ni2Si appeared in the specimens. The d-Ni2Si precipitates had two variations in mutual perpendicular directions, as shown in central darkfield images (Fig. 6(d) and (f)). d-Ni2Si and d0 -Ni2Si were marked to distinguish the two orientations of d-Ni2Si particles. Both of d0 -Ni2Si and d-Ni2Si had orthorhombic structure with the same structure parameters: a ¼ 0.706 nm, b ¼ 0.499 nm, and c ¼ 0.372 nm [28]. From Fig. 6(d), the crystal orientation relationship between copper matrix and precipitates (b-Ni3Si, d-Ni2Si and d0 -Ni2Si) has been determined as following: ð022Þcu ==ð011Þb ==ð010Þd ; ½100cu ==½100b ==½001d ; ð022Þcu ==ð011Þb ==ð010Þd0 ; ½100cu ==½100b ==½001d0 . The results of Fourier transform also show that two orientations of d-Ni2Si precipitates with mutual perpendicular growth direction contributed the two sets of mutual perpendicular diffraction spectrums (Fig. 6(b)). The SADP with electron beam parallel to [112]Cu were indexed in some papers [17,29], but the satellites spots arrowed diffraction spots of Cu matrix symmetrically in [112]Cu crystal zone axis have not been indexed. Here, the diffraction spots marked by white arrows in Fig. 6(e) were firstly determined, which were due to the diffraction of d-Ni2Si precipitates. The corresponding crystal orientation relationship between copper matrix and precipitates has been determined as following: ð111Þcu ==ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d .
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Fig. 4. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 15 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [112]Cu; (d) Selected-area diffraction pattern of (b), beam direction along [001]Cu.
3.4. Stress relaxation behaviour The home-made device for stress relaxation test and the stress relaxation behaviours of designed alloy aged at 450 C for 60 min are shown in Fig. 7. The home-made device is shown in Fig. 7(a). “S”
is the specimen. “B” is a bolt, and one end of the specimen was fixed by blocks “C” and “D”. First of all, the specimen lies on a location with an initial height of H0, and then the specimen is jacked up to Hmax by the bolt in an incubator with temperatures of 20 C, 100 C, 150 C, 200 C for specified time of tn (n ¼ 1,2.n). After stress
Fig. 5. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 60 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [112]Cu; (d) Selected-area diffraction pattern of (b), beam direction along [001]Cu.
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Fig. 6. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 480 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [100]Cu; (d) Central dark-field micrograph of precipitates, diffraction spots were selected from the marked round zone in (c); (e) Selected-area diffraction pattern, beam direction along [112]Cu; (f) Central dark-field micrograph of precipitates.
relaxation experiment, the specimen remains at a new height of Hn (n ¼ 1, 2.n, respectively). H0, Hmax, Hn (n ¼ 1, 2.n) are measured by tool microscope. The stress relaxation rate is expressed as Rn ¼ (Hmax Hn)/(Hmax H0), (n ¼ 1, 2.n). Stress relaxation occurs in the elastic materials at elevated temperature environments. Stress relaxed rapidly in the initial test stage, and then approached stable value (Fig. 7(b) and (c)). After loading for 100 h, stress relaxation rates were 4.03% at 20 C, 6.55% at 100 C, 8.52% at 150 C and 9.81% at 200 C, respectively. The stress relaxation was considered to occur by logarithmic creep caused by a relativelyshort range motion of dislocations. Mg solute atoms increased the density of mobile dislocation, which blocked the dislocation movement [21]. Meanwhile, nano-scale precipitates did not coarsen obviously with ageing time due to the addition of Al element [30]. Precipitates blocked dislocation movement and improved the anti-stress relaxation property. 3.5. Strength mechanism The increase of Orowan stress is inversely proportional to the inter-precipitate spacing l. The increase of strength by the alloying
element can be discussed qualitatively by estimating l, which is pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ffi pffiffiffiffiffiffiffiffi expressed as l ¼ ð1:23 2p=ð3f Þ 2 2=3Þr [31]. Where r is the average radius of precipitates and f is the volume fraction of precipitates. Addition of Al element slowed down the growth of precipitates [30]. As showed in Fig. 4e6, dimension of precipitate in specimens aged at 450 C for 480 min only increased by 6.5 nm, i.e, r in the equation increased slowly with ageing time. In the initial stage of ageing process, the f increased obviously with the ageing time. Therefore, the increase of Orowan stress and strength of specimens increased rapidly. Further increasing the ageing time, l increased slightly with f and r, the strength of the specimens decreased slowly. 3.6. Property comparison The comparison of properties of various elastic coppers and designed CueNieSi alloy are shown in Fig. 8. The C17200 (CueBe alloy) shows good properties but bad anti-stress relaxation resistance, stress relaxation rate of Cu-1.8Be-0.21Co alloy(1/2 HM) was 25% as performed at 150 C for 100 h [32]. While stress relaxation rate of designed CueNieSi alloy is only 8.52% under the same
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Fig. 7. Home-made device for stress relaxation measurement and stress relaxation rate curves of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy as tested at different temperatures. (a) Home-made device; (b) Experimental curves of stress relaxation. (c) Stress relaxation rate.
Fig. 8. Comparison of properties of various elastic coppers and designed alloy. (a) Hardness; (b) Ultimate tensile strength; (c) Electrical conductivity; (d) Stress relaxation rate, 150 C for 100 h.
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condition. CueNieSn, CueNieZn and CueNieAl alloys have good tensile strength and anti-stress relaxation property, but their electrical conductivity are low than 15% IACS [33,34], while that of the designed CueNieSi alloy is 26.5% IACS, which also is higher than that of the C17200 (CueBe alloy). Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy would be an ultrahigh strength copper and it exhibits a promising prospect in replacing CueBe alloys in some applications. 4. Conclusions This study focused on designing a novel Cu-6.0 Ni-1.0 Si-0.5 Al0.15 Mg-0.1 Cr alloy by adding suitable alloying elements and applying thermal-mechanical treatments. The results show that the designed alloy achieves good comprehensive properties and owns a promising prospect in replacing CueBe alloys. The main conclusions can be drawn as follows: 1) A new environment friendly, ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr copper alloy has been obtained by designing suitable alloy system and appropriate thermalmechanical treatments, resulting in desirable combination of mechanical properties, electrical conductivity and anti-stress relaxation resistance. 2) Appropriate thermal mechanical treatment method for designed alloy ingot is that: homogenization treated at 940 C for 4 h, hot rolled at 850 C by 80%, solution treated at 970 C for 6 h, cold rolled by 50%, and finally aged at 450 C for 1 h. 3) The designed alloy achieved ultrahigh strength (1090 MPa), good electrical conductivity (26.5% IACS), and low stress relaxation rate (8.52%) as performed at 150 C for 100 h, which are comparable to those of CueBe alloys. 4) The satellites spots arrowed diffraction spots of Cu matrix symmetrically in [112]Cu zone crystal axis have been determined firstly, which are due to the electron diffraction of nanoscale d-Ni2Si precipitates. The corresponding crystal orientations between matrix and precipitates is that: ð111Þcu == ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d . 5) b-Ni3Si phase precipitated in the initial ageing stages. With further increasing the ageing time, b-Ni3Si and d-Ni2Si phase precipitates appeared, which contribute the ultrahigh strength by Orowan strengthening. Acknowledgements The authors are pleased to acknowledge the financial supply supported by the Project supported by the National Natural Science Foundation of China (51271203), Hunan Provincial Natural Science Foundation of China (11JJ2025), and the Aid program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province. Q. Lei is grateful for the Hunan Provincial Innovation Foundation for Postgraduate (CX2011B107), the Excellent Doctor Degree Thesis Support Foundation of Central South University (2012ybjz 003) and the Scholarship Award for Excellent Doctoral Student granted by Ministry of Education. References [1] Pérez-Landazábal JI, Recarte V, Nó ML, San Juan J. Determination of the order in g1 intermetallic phase in Cu-Al-Ni shape memory alloys. Intermetallics 2003;11(9):927e30.
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Short communication
A new ultrahigh strength CueNieSi alloy Q. Lei a, Z. Li a, b, *, T. Xiao a, Y. Pang a, Z.Q. Xiang a, W.T. Qiu a, Z. Xiao c a
School of Materials Science and Engineering, Central South University, Changsha 410083, China State Key Laboratory of Powder Metallurgy, Changsha 410083, China c Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 4100083, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 February 2013 Received in revised form 2 May 2013 Accepted 19 May 2013 Available online
A new ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy has been developed by alloying design and thermal mechanical treatment. Electrical conductivity, mechanical properties of the designed alloy such as hardness, tensile strength, yield strength, elongation and anti-stress relaxation resistance were tested. The mechanical properties and electrical properties of designed alloy were comparable to those of CueBe alloys, and its anti-stress relaxation resistance even was better than that of CueBe alloys at evaluated temperature. Microstructure observation revealed that b-Ni3Si phase precipitated in the initial stages ageing process, with further increasing the ageing time, b-Ni3Si, and d-Ni2Si phase precipitates appeared and contributed to the ultrahigh strength by Orowan strengthening. The satellites spots around diffraction spots of Cu matrix symmetrically in [112]Cu zone crystal axis have been determined, resulting that satellites spots are from the electron diffraction of d-Ni2Si precipitates. The crystal orientation relationship between matrix and precipitates is that: ð111Þcu ==ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d . Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: B. Alloy design B. Precipitates C. Melting C. Thermomechanical treatment D. Microstructure G. Automotive uses
1. Introduction High strength materials play significant role on the automaton industry, automobile industry, electrical and electronics industry. Copper alloys are widely used due to their high strength and good electrical conductivity [1,2]. CueBe alloys is the most intensive used ones among present Cu-based elastic materials, owing to their good mechanical and electrical properties [3]. However, the toxicity of beryllium element in CueBe alloy was harmful to human body in the production process [4], and anti-stress relaxation properties of CueBe alloys are poor as used at elevated temperature. A lot of attempts have been made by other researchers to developed novel copper alloys for replacing the CueBe alloys before, such as CueNie Sn, CueNieZn and CueNieAl system alloys [5e7]. These alloys exhibit ultrahigh strength (>1000 MPa). However, their electrical conductivities (<15 %IACS) are less than those of CueBe alloys (w25 %IACS). Encouragingly, CueNieSi alloys have aroused considerable interest because of their high strength and good electrical conductivity [8e12]. The high strength of CueNieSi alloys is attributed to the nano-scale precipitates during precipitation [13e16]. In our * Corresponding author. School of Materials Science and Engineering, Central South University, Changsha 410083, China. Tel.: þ86 731 88830264; fax: þ86 731 88876692. E-mail address: [email protected] (Z. Li). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.05.013
previous investigations, strength could be effectively enhanced through increasing the content of Ni and Si. Cu-8.0Ni-1.8Si system alloys were designed and they showed good mechanical properties [17e19], however, some precipitates composed by Ni and Si could not be re-dissolved back to matrix even at very high solution temperature [17], as a result that it is hard to achieve the desired performance. Therefore, it is a key issue to lower the content of Ni and Si but not to decrease the strength of the alloy. One of the effective ways is addition of other suitable alloying elements [20]. Al and Mg could affect the precipitation kinetics and increased the formation rate of precipitates [9,21]. Effects of additions of Mg [21] and Cr [22] on the microstructure and property of CueNieSi alloys have been reported, showing that they could effectively improve the strength of alloy. Therefore, Al, Mg and Cr elements were added in the designed CueNieSi alloys. Copper base alloys with high strength and good electrical conductivity are generally solution treated by high temperature quenching and subsequent ageing. Nano-scale precipitates formed in copper matrix through these processes, as a result that the alloy is strengthened and its electrical conductivity is increased. Studies on the mechanical properties and microstructure of ultrahigh strength copperenickelesilicon (CueNieSi) alloys have been carried out thus far [11,16e19], and some interesting results on precipitation have been obtained. However, the appropriate thermomechanical treatment in the Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1
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Cr alloy has not been investigated yet. Not only the chemical composition but also the process condition play important role in the comprehensive properties. Therefore, thermo-mechanical treatment of the ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloys should be investigated, by which an optimum process condition for improving the properties can be obtained. In the present work, a novel copper alloy with nominal composition of Cu-6.0% Ni-1.0% Si-0.5% Al-0.15% Mg-0.1% Cr (weight percentage) has been designed and its microstructure evolution and properties have been investigated.
time of 15 s. Electrical conductivity was measured using a doublearm electrical bridge device. Tensile tests were performed using an Instron 8019 tester type machine with a constant strain rate of 103 s1. Microstructure observation was operated using a JEM2100F transmission electron microscope with operation voltage of 200 kV and point resolution of 0.19 nm. Thin foils for TEM and HRTEM observations were prepared by ion beam milling. Stress relaxation experiment was carried out on a home-made device designed according to ASTM E328-02(2008) [23]. 3. Results and discussion
2. Experimental details 3.1. Thermo-mechanical treatment Firstly, pure Cu, pure Ni, and pure Cr blocks were melted in a medium-frequency induction furnace. Then, Cu-14 wt.% Si, Cu20 wt.% Mg master alloys and pure Al blocks were added into the melt. Melting and casting were carried out in N2 atmosphere to prevent the melt and ingot from oxidization. After removing the surface oxides, a 30 mm thickness ingot was homogenization treated at 940 C for 4 h, followed by hot rolled at 850 C with total deformation of 80%, thickness from 30 mm to 6 mm. Supersaturated solid solution is the basis of precipitation hardening, therefore, solution treatment at 950 C, 960 C, 970 C and 980 C for 6 h were performed on the hot-rolling specimens to determine appropriate solution treatment temperature, respectively. Then the specimens were cold rolled to decrease the thickness from 6 mm to 3 mm, finally ageing treated at 450 C and 500 C for different time. Metallographic observation was carried out using a Leica EC3 optical microscope. Hardness tests were operated using an HV-5 typical Vickers hardness tester, with a load of 3 kg and holding
The initial microstructure of as-cast ingot is shown in Fig. 1(a). Developed dendrite grains are observed. Meanwhile, some nonequilibrium precipitates appeared between the dendrite arms (Fig. 1(a)), which are harmful to subsequent thermal-mechanical heat treatment. To eliminate the dendrite structures, homogenization treatment at 940 C for 4 h was conducted, by which the dendrite structures were eliminated (Fig. 1(b)). Backscatter electron image and elemental mapping images of the homogenization specimens were taken. The elemental mapping images of Fig. 1(c) is shown in Fig. 1(d)e(i), which indicates that the particles were rich in Ni and Si but poor in Cu. Al, Mg and Cr distributed inside the grains evenly. The elemental line distribution image also agreed well with the elemental mapping results (Fig. 1(j)). Fig. 2(a) shows microstructures of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after hot rolling at 850 C with a total deformation of 80%. After hot rolling deformation, initial large grains were
Fig. 1. Optical microstructure, SEM micrograph and the elemental mapping images of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy. (a) As cast; (b) Homogenized at 940 C for 4 h; (c) Backscatter electron image; Elemental mapping images: (d) Cu; (e) Ni; (f) Si, (g) Al; (h) Mg; (i) Cr; (j) Line distribution of Cu, Ni, Si and Al elements.
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Fig. 2. The microstructure of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy as hot rolled at 850 C and solution treatment at different temperatures for 6 h. Hot rolling: (a) Rolling surface; (b) Rolling side; Solution treatment at different temperature: (c) 950 C; (d) 960 C; (e) 970 C; (f) 980 C.
replaced by fine recrystallized grains. Dynamic recrystallization finished completely, annealing twins appeared in the rolling surface (as shown in Fig. 2(a)) and some recrystallization grains were elongated along the rolling direction in the rolling slide (Fig. 2(b)). Fig. 2(c)e(f) shows microstructures of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after solution treatments. The content of residual intermetallic compounds decreased with increase of temperature. As the specimens were solution treated at 950 C or 960 C for 6 h, a lot of intermetallic compounds still can be observed in the matrix (Fig. 2(c) and (d)). While, for the specimens solution treatment at 970 C for 6 h, intermetallic compounds re-dissolved into the matrix completely, indicating that supersaturated solid solution formed (Fig. 2(e)). Further increase the temperature to 980 C, the grains coarsened seriously (Fig. 2(f)). Based on the microstructure observations, intermetallic compounds and grain size, the appropriate solution treatment was determined, 970 C 6 h.
3.2. Hardness and electrical conductivity After solution treatment at 970 C for 6 h, the solution treated bulk materials was cold rolled with total deformation of 50%. Then it was cut into standard specimens for hardness and electrical conductivity, respectively. Fig. 3 shows the variation of hardness and electrical conductivity of alloy with ageing time. With ageing time, hardness of specimens increased rapidly and reached peak value, then decreased. This decrease of hardness is a thermally activated process [24]. When specimens were aged for more than 30 min, hardness of specimens decreased more quickly as aged at 500 C, than that of specimens aged at 450 C (Fig. 3(a)). When the ageing temperature was 500 C, the hardness decreased due to the coarsening and dissolution of the precipitates. Electrical conductivity of specimens increased rapidly in the initial ageing process, then increased slowly (Fig. 3(b)). With the same ageing time, the
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Fig. 3. Properties of designed alloy. (a) Hardness; (b) Electrical conductivity; (c) Stressestrain curve of designed alloy aged at 450 C for 60 min.
electrical conductivities of the specimens aged at 500 C were higher than those of specimens aged at 450 C. Considering the combination of properties, the more appropriate ageing method was 450 C 60 min for the designed alloy. Under this ageing treatment, hardness of the alloy was 341 HV and the corresponding electrical conductivity was 26.5% IACS. Electrical resistivity is sensitive to a number of microstructural factors including vacancy concentration, solute concentration in the matrix, precipitate or cluster size, and precipitate volume fraction [25]. The results from tensile test of specimens aged at 450 C for 60 min (Fig. 3(c)), indicating that the tensile strength, yield strength and elongation of designed alloy were 1090 MPa, 940 MPa and 3.5%, respectively. 3.3. Precipitation Microstructures and selected-area diffraction patterns (SADPs) of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 15 min are shown in Fig. 4. Some nano-scale precipitating particles can be seen (Fig. 4(a)). The mean dimension of particles was w8 nm, as shown in Fig. 4(b). The SADPs are shown in Fig. 4(c) and (d), some additional diffraction spots different from the spots from the electron diffraction of Cu matrix could be detected. The indexing results of Fig. 4(c) and (d) illustrate that additional spots are derived from the electron diffraction of bNi3Si phase precipitates. Cu matrix has a face central cubic structure, a ¼ 0.3615 nm, b-Ni3Si phase precipitate has a prime cubic structure, a ¼ 0.3351 nm [26,27]. Based on the analysis results of SADPs in Fig. 4, the crystal orientation relationship between matrix and b-Ni3Si precipitate can be indexed as follows: ð001Þcu == ð001Þb ; ½100cu ==½100b ; ð110Þcu ==ð220Þb ; ½112cu ==½112b . Fig. 5 showed microstructures and SADPs of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy aged at 450 C for 60 min. The precipitates coarsened slightly and the mean size of precipitates was w10 nm (Fig. 5(a) and (b)). Except some new faint diffraction spots from
other precipitates, the diffraction spots and their indexation results were similar to Fig. 4. In order to detect all precipitates and determine the crystal orientation relationship using SADPs with beam along [112]Cu, the ageing time was further increased to 480 min Fig. 6 shows the brightfield micrograph, central dark-field micrograph, HRTEM micrograph and SADPs of specimens as aged at 450 C for 480 min. The precipitates obviously coarsened compared to that ageing treated at 450 C for 60 min d-Ni2Si precipitates with mutual perpendicular directions formed and determined, as shown in the micrographs images (Fig. 6(a),(b),(d) and (f)). The index results in Fig. 6(c) and (e) illustrated that two kinds of precipitate phases of b-Ni3Si and d-Ni2Si appeared in the specimens. The d-Ni2Si precipitates had two variations in mutual perpendicular directions, as shown in central darkfield images (Fig. 6(d) and (f)). d-Ni2Si and d0 -Ni2Si were marked to distinguish the two orientations of d-Ni2Si particles. Both of d0 -Ni2Si and d-Ni2Si had orthorhombic structure with the same structure parameters: a ¼ 0.706 nm, b ¼ 0.499 nm, and c ¼ 0.372 nm [28]. From Fig. 6(d), the crystal orientation relationship between copper matrix and precipitates (b-Ni3Si, d-Ni2Si and d0 -Ni2Si) has been determined as following: ð022Þcu ==ð011Þb ==ð010Þd ; ½100cu ==½100b ==½001d ; ð022Þcu ==ð011Þb ==ð010Þd0 ; ½100cu ==½100b ==½001d0 . The results of Fourier transform also show that two orientations of d-Ni2Si precipitates with mutual perpendicular growth direction contributed the two sets of mutual perpendicular diffraction spectrums (Fig. 6(b)). The SADP with electron beam parallel to [112]Cu were indexed in some papers [17,29], but the satellites spots arrowed diffraction spots of Cu matrix symmetrically in [112]Cu crystal zone axis have not been indexed. Here, the diffraction spots marked by white arrows in Fig. 6(e) were firstly determined, which were due to the diffraction of d-Ni2Si precipitates. The corresponding crystal orientation relationship between copper matrix and precipitates has been determined as following: ð111Þcu ==ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d .
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Fig. 4. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 15 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [112]Cu; (d) Selected-area diffraction pattern of (b), beam direction along [001]Cu.
3.4. Stress relaxation behaviour The home-made device for stress relaxation test and the stress relaxation behaviours of designed alloy aged at 450 C for 60 min are shown in Fig. 7. The home-made device is shown in Fig. 7(a). “S”
is the specimen. “B” is a bolt, and one end of the specimen was fixed by blocks “C” and “D”. First of all, the specimen lies on a location with an initial height of H0, and then the specimen is jacked up to Hmax by the bolt in an incubator with temperatures of 20 C, 100 C, 150 C, 200 C for specified time of tn (n ¼ 1,2.n). After stress
Fig. 5. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 60 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [112]Cu; (d) Selected-area diffraction pattern of (b), beam direction along [001]Cu.
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Fig. 6. Microstructure evolution of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy after ageing treatment at 450 C for 480 min. (a) Bright-field image; (b) High resolution image; (c) Selected-area diffraction pattern of (a), beam direction along [100]Cu; (d) Central dark-field micrograph of precipitates, diffraction spots were selected from the marked round zone in (c); (e) Selected-area diffraction pattern, beam direction along [112]Cu; (f) Central dark-field micrograph of precipitates.
relaxation experiment, the specimen remains at a new height of Hn (n ¼ 1, 2.n, respectively). H0, Hmax, Hn (n ¼ 1, 2.n) are measured by tool microscope. The stress relaxation rate is expressed as Rn ¼ (Hmax Hn)/(Hmax H0), (n ¼ 1, 2.n). Stress relaxation occurs in the elastic materials at elevated temperature environments. Stress relaxed rapidly in the initial test stage, and then approached stable value (Fig. 7(b) and (c)). After loading for 100 h, stress relaxation rates were 4.03% at 20 C, 6.55% at 100 C, 8.52% at 150 C and 9.81% at 200 C, respectively. The stress relaxation was considered to occur by logarithmic creep caused by a relativelyshort range motion of dislocations. Mg solute atoms increased the density of mobile dislocation, which blocked the dislocation movement [21]. Meanwhile, nano-scale precipitates did not coarsen obviously with ageing time due to the addition of Al element [30]. Precipitates blocked dislocation movement and improved the anti-stress relaxation property. 3.5. Strength mechanism The increase of Orowan stress is inversely proportional to the inter-precipitate spacing l. The increase of strength by the alloying
element can be discussed qualitatively by estimating l, which is pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ffi pffiffiffiffiffiffiffiffi expressed as l ¼ ð1:23 2p=ð3f Þ 2 2=3Þr [31]. Where r is the average radius of precipitates and f is the volume fraction of precipitates. Addition of Al element slowed down the growth of precipitates [30]. As showed in Fig. 4e6, dimension of precipitate in specimens aged at 450 C for 480 min only increased by 6.5 nm, i.e, r in the equation increased slowly with ageing time. In the initial stage of ageing process, the f increased obviously with the ageing time. Therefore, the increase of Orowan stress and strength of specimens increased rapidly. Further increasing the ageing time, l increased slightly with f and r, the strength of the specimens decreased slowly. 3.6. Property comparison The comparison of properties of various elastic coppers and designed CueNieSi alloy are shown in Fig. 8. The C17200 (CueBe alloy) shows good properties but bad anti-stress relaxation resistance, stress relaxation rate of Cu-1.8Be-0.21Co alloy(1/2 HM) was 25% as performed at 150 C for 100 h [32]. While stress relaxation rate of designed CueNieSi alloy is only 8.52% under the same
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Fig. 7. Home-made device for stress relaxation measurement and stress relaxation rate curves of Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy as tested at different temperatures. (a) Home-made device; (b) Experimental curves of stress relaxation. (c) Stress relaxation rate.
Fig. 8. Comparison of properties of various elastic coppers and designed alloy. (a) Hardness; (b) Ultimate tensile strength; (c) Electrical conductivity; (d) Stress relaxation rate, 150 C for 100 h.
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condition. CueNieSn, CueNieZn and CueNieAl alloys have good tensile strength and anti-stress relaxation property, but their electrical conductivity are low than 15% IACS [33,34], while that of the designed CueNieSi alloy is 26.5% IACS, which also is higher than that of the C17200 (CueBe alloy). Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr alloy would be an ultrahigh strength copper and it exhibits a promising prospect in replacing CueBe alloys in some applications. 4. Conclusions This study focused on designing a novel Cu-6.0 Ni-1.0 Si-0.5 Al0.15 Mg-0.1 Cr alloy by adding suitable alloying elements and applying thermal-mechanical treatments. The results show that the designed alloy achieves good comprehensive properties and owns a promising prospect in replacing CueBe alloys. The main conclusions can be drawn as follows: 1) A new environment friendly, ultrahigh strength Cu-6.0 Ni-1.0 Si-0.5 Al-0.15 Mg-0.1 Cr copper alloy has been obtained by designing suitable alloy system and appropriate thermalmechanical treatments, resulting in desirable combination of mechanical properties, electrical conductivity and anti-stress relaxation resistance. 2) Appropriate thermal mechanical treatment method for designed alloy ingot is that: homogenization treated at 940 C for 4 h, hot rolled at 850 C by 80%, solution treated at 970 C for 6 h, cold rolled by 50%, and finally aged at 450 C for 1 h. 3) The designed alloy achieved ultrahigh strength (1090 MPa), good electrical conductivity (26.5% IACS), and low stress relaxation rate (8.52%) as performed at 150 C for 100 h, which are comparable to those of CueBe alloys. 4) The satellites spots arrowed diffraction spots of Cu matrix symmetrically in [112]Cu zone crystal axis have been determined firstly, which are due to the electron diffraction of nanoscale d-Ni2Si precipitates. The corresponding crystal orientations between matrix and precipitates is that: ð111Þcu == ð111Þb ==ð021Þd ; ½112cu ==½112b ==½012d . 5) b-Ni3Si phase precipitated in the initial ageing stages. With further increasing the ageing time, b-Ni3Si and d-Ni2Si phase precipitates appeared, which contribute the ultrahigh strength by Orowan strengthening. Acknowledgements The authors are pleased to acknowledge the financial supply supported by the Project supported by the National Natural Science Foundation of China (51271203), Hunan Provincial Natural Science Foundation of China (11JJ2025), and the Aid program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province. Q. Lei is grateful for the Hunan Provincial Innovation Foundation for Postgraduate (CX2011B107), the Excellent Doctor Degree Thesis Support Foundation of Central South University (2012ybjz 003) and the Scholarship Award for Excellent Doctoral Student granted by Ministry of Education. References [1] Pérez-Landazábal JI, Recarte V, Nó ML, San Juan J. Determination of the order in g1 intermetallic phase in Cu-Al-Ni shape memory alloys. Intermetallics 2003;11(9):927e30.
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