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A low Sn content Cu-Ni-Sn alloy with high strength and good ductility
Materials Science & Engineering A 746 (2019) 154–161
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A low Sn content Cu-Ni-Sn alloy with high strength and good ductility a
a
a
Baomin Luo , Daoxi Li , Chao Zhao , Zhi Wang a b
a,b
, Zongqiang Luo
a,b
, Weiwen Zhang
a,b,⁎
T
Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou 510640, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Low Sn content Cu-Ni-Sn alloy Mechanical properties Bimodal structure
Avoiding Sn segregation is vital to produce high performance large scale Cu-Ni-Sn alloys. In this work a low Sncontaining Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy with high strength and ductility was developed by substituting Sn in Cu-Ni-Sn alloy with Al and Si. The results show that Si leads to a bimodal structure, which has coarse unrecrystallization regions with a plenty of < 111 > fiber texture and fine equaixed recrystallization grains. Such a bimodal structure results into good mechanical properties showing a high tensile strength of 861 MPa and a large elongation of 18%, which is comparable to Cu-15Ni-8Sn alloy. The strengthening effects are ascribed to precipitate strengthening, grain refining and twin boundary strengthening in the fine equaixed recrystallization grains and texture strengthening in the coarse unrecrystallization regions.
1. Introduction Cu-Ni-Sn alloys have been attracted increasing attention and become one of the most prominent substitutes for Cu-Be alloys in aerospace, oil and gas exploration areas, due to their outstanding strength, excellent electrical conductivity, exceptional bearing properties and high corrosion resistance [1–3]. Among Cu-Ni-Sn alloys systems, Cu15Ni-8Sn alloy has high strength combined with good ductility due to the formation of coherent DO22 phase from the spinodal decomposition during aging treatment [4,5]. However, it is hardly to avoid severe macro-segregation of Sn during solidification which makes it difficult to produce sheet products when Sn content is high [6,7]. An effective way to avoid Sn macro-segregation is to develop low Sn-containing Cu-Ni-Sn alloys with comparable mechanical properties. However, it is found that when the Sn content in Cu-Ni-Sn alloys is less than 4 wt%, the precipitate strengthening, one of the main strengthening mechanism, will be disappeared or weakened due to the secondary phase was formed by a nucleation process instead of spinodal decomposition [8]. The secondary phases in low Sn content Cu-Ni-Sn alloys show a much larger size (tens nm in diameter and length) than that (~ 2 nm) in high Sn content Cu-Ni-Sn alloys, which results into weaken strengthening effect [8,9]. Micro-alloying has been proved to be an effective way to improve the strength of the low Sn content Cu-NiSn alloys. For example, Ni3Ti/Ni3Nb phase formed in the grain boundaries with the addition of Ti/Nb in the Cu-Ni-Sn alloys, which can give rise to grain refinement and suppress the formation of
discontinuous precipitation [4,10]. It was also found that the addition of small amounts of refractory elements (e.g. Mo, V and Ta) can inhibit the growth of the phase along the grain boundary without inhibiting the spinodal decomposition [11]. It has been found that the strength of low Sn-containing Cu-Ni-Sn alloys can be improved by minor addition of Al [6,12]. For example, Rhu et al. found that Al addition can refine grain and promote the precipitation of CuAl2 particles, which improves the mechanical properties of Cu-6Ni-2Mn-2Sn [6]. Han et al. found that the enhanced tensile strength with the addition of Al is attributed to the solid solution hardening effect [12]. Therefore, Al may sufficiently improve the strength of the low Sn-containing Cu-Ni-Sn alloys owing to grain refinement, CuAl2 nanoparticle precipitation and/or solid solution hardening. In the present work, a novel low Sn-containing Cu-Ni-Sn alloy with both high strength and good ductility was developed by adding minor Al and Si content. As mentioned above, the effect of Al on the mechanical prosperities of low Sn-containing Cu-Ni-Sn alloys were clearly studied, which is helpful for understanding the strengthening effect on the new alloy produced in this work. Furthermore, we found that the minor addition of Si can significantly improve the mechanical properties of Cu-Ni-Sn-Al alloy. Therefore, the effect Si on the microstructure and mechanical properties of a low Sn content Cu-9Ni-2.5Sn-1.5Al alloy was studied.
⁎ Corresponding author at: Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China. E-mail address: [email protected] (W. Zhang).
https://doi.org/10.1016/j.msea.2018.12.120 Received 18 October 2018; Received in revised form 27 December 2018; Accepted 29 December 2018 Available online 31 December 2018 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Table 1 Chemical composition of the alloys. Alloy designation
Cu-9Ni-2.5Sn-1.5Al Cu-9Ni-2.5Sn-1.5Al-0.5Si
wt% Ni
Sn
Al
Si
Cu
9.17 9.72
2.53 2.56
1.38 1.34
̶ 0.45
Balance Balance
2. Processing and experimental procedure The alloys were prepared by electric melting of high purity Cu, Ni, Sn, Al and Si in an intermediate frequency induction furnace and then casting them into a steel mold with a diameter of 60 mm and a length of 160 mm. The chemical compositions of the alloys are summarized in Table 1. The ingots were homogenized at 930 °C for 2.5 h and machined to cylindrical-shaped samples with a diameter of 50 mm and a height of 50 mm for extrusion. Hot extrusion was performed at 950 °C with an extrusion ratio of 17 and a ram speed of 3 mm/s by using a 2000 KN vertical extruding machine. The as-extruded bars were cut into cylindrical tensile samples with a gauge section of 5 mm in diameter and 25 mm in length and the tensile tests were performed using a SANS CMT5105 material test machine in air with an initial tensile velocity of 1 mm/min at room temperature. At least three samples were tested under each condition. Optical micrograph (OM) was observed using a LEICA/DMI 5000 M optical microscope. The mean grain size was measured by the three round intercept method based on the observation of OM according to Chinese GB/T 6394-2002, more than five hundred grains were measured in each specimen. The volume fraction of the coarse unrecrystallized grains was obtained from 15 random areas based on the observation of OM. Texture analyses by EBSD were conducted using a field emission FEI NONA430 scanning electron microscope operated at 30 kV. Samples for EBSD measurement were prepared by mechanical polishing with progressively finer water-based diamond suspension (0.5 µm) followed by electro-polishing using a solution of 50% phosphoric acid and 50% alcohol at 6.5 V for 60 s. TEM specimens were prepared by twin jet electro-polishing in a solution of 30% nitric acid and 70% methanol held at −20 to −30 °C. The specimens were finished for surface cleaning by ion milling using a Gatan precision Ion Polishing System (PIPS) at an operating voltage of 3 kV for 20 min. Microstructure analyses were performed on a JEM-2100F transmission electron microscope. The area fraction of precipitates was measured from five different views of central dark-field TEM micrograph.
Fig. 1. Typical tensile stress-strain curves of as-extruded 0Si and 0.5Si alloys. Table 2 Tensile properties of 0Si and 0.5Si alloys. Alloy
σb (MPa)
σ0.2 (MPa)
Ɛ (%)
0Si 0.5Si
713 ± 10 861 ± 7.8
453 ± 7.5 698 ± 13
23 ± 0.3 18 ± 1.1
Fig. 2. Comparison between the present alloys and other typical Cu-15Ni-8Sn alloys with the tensile strength versus elongation [4,13–17].
3.2. Structural characterization 3. Results and discussion
Fig. 3 shows the microstructure of the dynamically recrystallized grains in the as-extruded 0Si and 0.5Si alloys. The 0Si alloy has uniform and equiaxed dynamic recrystallized grains with an average diameter of 46 ± 2.2 µm (Fig. 3a and c). While the 0.5Si alloy shows a bimodal grain structure consisting of fine dynamic recrystallized grains with an average diameter of 34 ± 4.3 µm and elongated coarse unrecrystallized grains (marked as red arrows in Fig. 3b and d), which is much different from the microstructure of the 0Si alloy. Such microstructure change results into a large difference in mechanical properties, as shown in Fig. 1. This bimodal grain structure was often found in the hot-extruded alloys [18–20], in which the equaixed grains were resulted from the recrystallization started from the grain boundaries and the elongated grains were formed during hot extrusion. Fig. 4 shows the TEM micrographs and the SADPs of the homogenized and as-extruded 0Si alloy. It indicates that fine γ′-Ni3Al particles precipitated in the matrix, and the crystal orientation relationship between Ni3Al and Cu matrix is [100]γ′//[200]Cu, [010] γ′//020]Cu and [001] γ′//[001]Cu. Fig. 5 shows the microstructure of the homogenized and as-extruded 0.5Si alloy. The homogenization treated sample shows
3.1. Mechanical properties Fig. 1 shows tensile properties of the as-extruded Cu-9Ni-2.5Sn1.5Al (marked as 0Si alloy) and Cu-9Ni-2.5Sn-1.5Al-0.5Si (marked as 0.5Si alloy) alloys. The ultimate tensile strength (σb), 0.2% offset yield stress (σ0.2) and strains to failure in tension (Ɛ) are summarized in Table 2. It is clear that the 0.5Si alloy has a much higher strength than the 0Si alloy. The yield and tensile strength of the 0.5Si alloy are 698 MPa and 861 MPa, respectively, while for 0Si alloy they are 453 MPa and 713 MPa, respectively. Although the elongation decreases slightly, the 0.5Si alloy still possesses good ductility of more than 18% plastic strain. Fig. 2 shows the comparison between the present alloys and other typical Cu-Ni-Sn alloys with the tensile strength versus elongation, indicating that the 0.5Si alloy, a low Sn content Cu-Ni-Sn alloy produced in this work, has a good combination of high tensile strength and good ductility.
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Fig. 3. OM images of the as-extruded 0Si and 0.5 Si alloys: (a and b) the plane perpendicular to the extrusion direction (ED); (c and d) the plane parallel to the extrusion direction (TD). The average diameter of dynamic recrystallized grains is given as d d in Fig. 3(a) and (b).
Fig. 4. TEM micrographs and SADPs of 0Si alloy: (a and b) homogenized at 930 °C for 2.5 h; (c and d) hot extruded at 950 °C. 156
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Fig. 5. TEM micrographs of 0.5Si alloy. (a): homogenized at 930 °C for 2.5 h and (b-e): hot extruded at 950 °C; (c) the HRTEM image of the rectangular area in (b) and the inset is the FFT of the red rectangle area in (c); (d) and (e) is bright-field and central dark-field image, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
Fig. 6. The elemental mapping images of the as-extruded 0.5Si alloy.
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Fig. 7. EBSD results showing IPF maps and local misorientation maps for 0Si alloy (a and b) and 0.5Si alloy (c and d), black lines denote high angle grain boundaries (HAGBs, > 10), red lines denote SAGBs (1°~10°); (e): misorientation profiles measured along the yellow arrow in (d), the black line shows point-to-point misorientation, while the red line shows point-to-origin misorientation respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
radians with a distance δ, and b is the Burges vector [24]. Taking b = 0.2556 nm for pure Cu, with values for θ and δ deriving from misorientation profile in Fig. 7e, the estimated local dislocation densities is 0.5–1 × 1013 m−2. Fig. 8 shows the {111} < 110 > Schmid factor distribution maps of the 0.5Si alloy. It shows that unrecrystallizaiton regions have a darker color indicating a lower Schmid factor (Fig. 8a). The average Schmid factor of the {111} < 110 > is 0.315 and 0.436 for the unrecrystallizaion regions (Fig. 8b) and recrystallization regions (Fig. 8c), respectively, indicating that high strength was achieved in the unrecrystallization regions. Fig. 9 shows the distribution of coincidence site lattice (CSL) boundaries and grain boundaries, in which random high-angle grain boundaries and Σ3 boundaries are denoted by the black lines and red lines, respectively. Σ3 boundaries occurred frequently in the microstructure with numerous long and parallel-side annealing twin boundaries, which are observed in recrystallization regions of both 0Si and 0.5Si alloys. However, no Σ3 boundary is observed in the unrecrystallization regions of the 0.5Si alloy (Fig. 9b). The proportions of Σ3 boundaries and HAGBs (55˚–60˚) are decreased, while SAGBs (≤5˚) are increased obviously with the addition of Si. The average number of twin grains within a parent grain was calculated by determining the number of grains in the region both including and excluding twin boundaries as defining individual grains. The calculated value of the 0Si and 0.5Si alloy is 3.17 and 2.70, respectively. During recrystallization, when growth stagnation occurred, the boundary between the recrystallizing grain and the deformed matrix was often in a region of low dislocation density and/or low misorientation angle [25]. The form of Σ3 twin boundaries had initiated new orientations without the rotation of new grain with respect to the parent grain and apparently offers the additional boundary energy required to facilitate the continued grain growth [26,27]. It means that the density of Σ3 twin boundary may be decreased when the grain growth was retarded. The 0.5Si alloy shows a lower density of Ʃ3 twin boundary than the 0Si alloy owing to that Ni3Si nanoparticles were formed with the addition of Si which can suppress the grain growth.
that a few fine precipitates in the matrix (Fig. 5a), while the hot-extruded sample shows that a high density fine precipitates (20 ± 3 nm) distributed in the Cu matrix and along the grain boundary (Figs. 5b-5e), which were identified as β-Ni3Si from the diffraction pattern (the insert in Fig. 5c). β-Ni3Si was commonly observed in Cu-Ni-Si alloys [21,22], which has a L12 ordering structure with structure parameters of a = 0.351 nm, Pm-3 m(221) [23]. The crystal orientation relationship between Ni3Si and Cu matrix is [100]β//[200]Cu, [010]β//020]Cu and [001]β//[001]Cu. The precipitate-free zones (PFZ) with ~227 nm width were observed on the both sides of the grain boundary according to Figs. 5d and 5e, which is further confirmed by the element mapping (Fig. 6). It indicates that Sn, Al and Si are uniform distributed while Ni is significantly depleted near the grain boundaries. The Ni3Si nanoparticles in the grain boundaries retarded the recrystallization process and kept the coarse grains away from recrystallization. Therefore, the bimodal structure shown in as-extruded 0.5Si alloy is the result of the competition between recrystallization (heating process) and non-recrystallization (recrystallization suppressed by the Ni3Si precipitates in the grain boundaries). Fig. 7 shows the texture in the as extruded 0Si and 0.5Si alloys. Their inverse pole figure (IPF) in the form of contour plot shows the texture intensity. The 0Si alloy shows relatively randomly oriented grains with < 111 > fiber texture which is a typical texture in Cu alloys after extrusion (Fig. 7a), and there are almost no local strains in the grains implying that grains are free of dislocations after recrystallization (Fig. 7b). While the 0.5Si alloy shows a mixture of relatively randomly equiaxed grains and the unrecrystallization grains with a strong < 111 > fiber texture (Fig. 7c). The local misorientation maps suggest that there are almost no local strains in the recrystallization regions of 0.5Si alloy (Fig. 7d). But in the unrecrystallization regions, a large number of small angle grain boundaries (SAGBs, 1°–10°) are observed. Unrecrystallized grains keep high dislocation density, which can pin the dislocation movement and result into the formation of SAGBs. Generally, the larger the long-range misorientation gradient in the same grain, the higher the dislocation density is. The equivalent geometrically necessary dislocation density can roughly be estimated by ρ ≈ θ/(bδ), where θ is the accumulated misorientation angle in 158
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Fig. 8. The {111} < 110 > Schmid factor distribution maps of the 0.5Si alloy (a), and the corresponding regions of unrecrystallized grains (b) and recrystallized grains(c); (d) and (e) is the {111} < 110 > Schmid factor distribution obtained from Fig. b and Fig. c, respectively.
grain size) which shows a good agreement with the experiment data in Cu-Al alloys [33]. In present study, the twin boundaries were considered as grain boundaries and the strengthening effect was roughly calculated by a modified Hall-Petch equation. Therefore, the yield strength increment (σGTBS) due to fine grain and twin boundaries can be estimated by Hall-Petch relationship [35]
3.3. Strengthening mechanisms The low Sn-containing Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy shows a super mechanical property with a high strength of 861 MPa and a large elongation of 18%, which is comparable to the high Sn-containing alloy such as Cu-15Ni-8Sn. The strengthening effect can be explained by the following factors: (i) Grain boundary strengthening (ii) Texture strengthening; (iii) Orowan precipitation strengthening; (iv) Others strengthening, such as strengthening by twins and solid solution. Parameters used in the yield strength calculations were shown in Table 3.
−1/2
df ⎡ σGTBS = σ0 + k ⎢Vf ⎜⎛ ⎟⎞ n ⎣ ⎝ ⎠
⎤ + Vc dc−1/2 ⎥ ⎦
(1)
σ0 is the fractional stress, k is slope, df and dc is the grain size of fine recrystallized grains and coarse unrecrystallized grains, respectively. Vf and Vc is the volume fraction of fine recrystallized grains and coarse unrecrystallized grains, respectively. n is the average number of twin grains within a grain of 0.5Si alloy, which was only considered in recrystallization regions. The value of σGTBS was calculated to be 142 MPa. (ii) Texture strengthening The texture is the result of both of hot extrusion and dynamic recovery. The texture strengthening mechanism is familiar with work
(i) Grain and twin boundary strengthening The microstructure of the as-extruded alloy consisted of a bimodal grain structure of fine recrystallized grains and coarse unrecrystallized grains. The grain boundary strengthening effect was taken into account for the enhanced strength in the fine recrystallized grains. In addition, it has reported that deformation twin boundaries can act as effective barriers to dislocation movements [32–34]. The additional strengthening contribution by deformation twins has been calculated by a modified Hall-Petch equation (using inter spacing between the twins as the “effective” 159
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Fig. 9. EBSD orientation maps showing the microstructural topology in (a) 0Si alloy, (b) 0.5Si alloy. (c) and (d) is CSL boundaries and misorientation angle, respectively. Black lines denote random high angle boundaries; red lines denote Σ3 boundaries.
hardening. The yield strength due to work hardening as given by Eq. (2) [36,37].
ΔσTS = MαGb ρ
τp = 0.81
(2)
ln(dp/ b) Gm b 2π (1 − ν )1/2 λ − dp
(3)
where b is the Burgers Vector, G is the shear modulus of the matrix, v is the Poisson’s radio, dp is the average radius of the precipitates, and λ is the spacing between particles in the glide plane. And λ can be calculated in Eq. (4) [39]:
where M is Taylor factor, α is a geometric constant (about 0.3), G is the shear modulus of the copper matrix, b is the Burgers vector, ρ is the dislocation density. The texture strengthening was only considered in unrecrystallization regions, and the yield strength increase from texture strengthening was 8 MPa. (iii) Orowan precipitation strengthening
λ=
1 dp 2
3π 2fp
(4)
where fp is the volume fraction of second-phase particles. The incremental strength increased by Orowan mechanisms can be calculated as [39]:
Orowan precipitation strengthening mechanism can be described by the following Orowan-Ashby equation [38]: Table 3 Parameters used in the yield strength calculations. Parameter
Description
Value
Units
Refs.
σ0 k Gm b dp fp Vc dc ρ M
Fractional stress A constant for Cu-6Sn Shear modulus of copper Burgers vector of copper The average radius of the Precipitates (0.5Si alloy) Volume fraction of the precipitates Volume fraction of coarse unrecrystallized grains Average diameter of coarse unrecrystallized grains Dislocation density in coarse unrecrystallized grains Taylor factor
60 356 46 0.2556 20 8.85% 26.8% 118 0.75 × 1013 3.06
MPa MPa/μm−1/2 GPa nm nm – – μm m−2 –
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Fig. 10. Yield strength increase calculations from different strengthening mechanisms.
ΔσOPS = Mτp
(5)
where M is Taylor factor. The calculated result showed that the increase strength caused by Orowan mechanisms was 470 MPa for 0.5Si alloy. The total yield strength strengthened by different strengthening mechanisms were shown in Fig. 10, and the Orowan precipitation strengthening plays the main role with 67.3% of total strengthening contribution. 4. Conclusions A low Sn-containing Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy was developed by casting and hot extrusion, showing a super mechanical property with a high strength of 861 MPa and a large elongation of 18%, which is comparable to the high Sn-containing alloy such as Cu-15Ni-8Sn. The addition of 0.5 wt% Si in the Cu-9Ni-2.5Sn-1.5Al alloy can significantly change the microstructure and mechanical properties. A bimodal microstructure was formed in the Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy, which has coarse unrecrystallization regions having a plenty of < 111 > fiber texture and fine equaixed recrystallization grains. The bimodal microstructure in the Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy give rises to super combination of mechanical property such as high strength and good ductility. The enhanced strength is contributed by: (i) grain refinement and twin boundary strengthening; (ii) precipitates strengthening; (iii) texture strengthening. Acknowledgements This work was sponsored by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003) and Guangdong Special Found Project of Applied Technology Research and Development (Grant No. 2016B090931002) References [1] Y. Zhang, Z. Xiao, Y. Zhao, Effect of thermo-mechanical treatments on corrosion behavior of Cu-15Ni-8Sn alloy in 3.5 wt% NaCl solution, Mater. Chem. Phys. 199 (2017) 54–66, https://doi.org/10.1016/j.matchemphys.2017.06.041. [2] J.B. Singh, W. Cai, P. Bellon, Dry sliding of Cu–15 wt% Ni–8 wt% Sn bronze: wear behaviour and microstructures, Wear 263 (2007) 830–841, https://doi.org/10.1016/j. wear.2007.01.061. [3] J. Caris, D. Li, J.J. Stephens, Microstructural effects on tension behavior of Cu–15Ni–8Sn sheet, Mater. Sci. Eng. A 527 (2010) 769–781, https://doi.org/10.1016/j.msea.2008.01. 061. [4] C. Zhao, W. Zhang, Z. Wang, Improving the mechanical properties of Cu-15Ni-8Sn alloys by addition of titanium, Materials 10 (2017), https://doi.org/10.3390/ma10091038. [5] M.J. Diánez, E. Donoso, M.J. Sayagués, The calorimetric analysis as a tool for studying the aging hardening mechanism of a Cu-10 wt%Ni-5.5 wt%Sn alloy, J. Alloy. Compd. 688 (2016) 288–294, https://doi.org/10.1016/j.jallcom.2016.07.021.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
A low Sn content Cu-Ni-Sn alloy with high strength and good ductility a
a
a
Baomin Luo , Daoxi Li , Chao Zhao , Zhi Wang a b
a,b
, Zongqiang Luo
a,b
, Weiwen Zhang
a,b,⁎
T
Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou 510640, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Low Sn content Cu-Ni-Sn alloy Mechanical properties Bimodal structure
Avoiding Sn segregation is vital to produce high performance large scale Cu-Ni-Sn alloys. In this work a low Sncontaining Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy with high strength and ductility was developed by substituting Sn in Cu-Ni-Sn alloy with Al and Si. The results show that Si leads to a bimodal structure, which has coarse unrecrystallization regions with a plenty of < 111 > fiber texture and fine equaixed recrystallization grains. Such a bimodal structure results into good mechanical properties showing a high tensile strength of 861 MPa and a large elongation of 18%, which is comparable to Cu-15Ni-8Sn alloy. The strengthening effects are ascribed to precipitate strengthening, grain refining and twin boundary strengthening in the fine equaixed recrystallization grains and texture strengthening in the coarse unrecrystallization regions.
1. Introduction Cu-Ni-Sn alloys have been attracted increasing attention and become one of the most prominent substitutes for Cu-Be alloys in aerospace, oil and gas exploration areas, due to their outstanding strength, excellent electrical conductivity, exceptional bearing properties and high corrosion resistance [1–3]. Among Cu-Ni-Sn alloys systems, Cu15Ni-8Sn alloy has high strength combined with good ductility due to the formation of coherent DO22 phase from the spinodal decomposition during aging treatment [4,5]. However, it is hardly to avoid severe macro-segregation of Sn during solidification which makes it difficult to produce sheet products when Sn content is high [6,7]. An effective way to avoid Sn macro-segregation is to develop low Sn-containing Cu-Ni-Sn alloys with comparable mechanical properties. However, it is found that when the Sn content in Cu-Ni-Sn alloys is less than 4 wt%, the precipitate strengthening, one of the main strengthening mechanism, will be disappeared or weakened due to the secondary phase was formed by a nucleation process instead of spinodal decomposition [8]. The secondary phases in low Sn content Cu-Ni-Sn alloys show a much larger size (tens nm in diameter and length) than that (~ 2 nm) in high Sn content Cu-Ni-Sn alloys, which results into weaken strengthening effect [8,9]. Micro-alloying has been proved to be an effective way to improve the strength of the low Sn content Cu-NiSn alloys. For example, Ni3Ti/Ni3Nb phase formed in the grain boundaries with the addition of Ti/Nb in the Cu-Ni-Sn alloys, which can give rise to grain refinement and suppress the formation of
discontinuous precipitation [4,10]. It was also found that the addition of small amounts of refractory elements (e.g. Mo, V and Ta) can inhibit the growth of the phase along the grain boundary without inhibiting the spinodal decomposition [11]. It has been found that the strength of low Sn-containing Cu-Ni-Sn alloys can be improved by minor addition of Al [6,12]. For example, Rhu et al. found that Al addition can refine grain and promote the precipitation of CuAl2 particles, which improves the mechanical properties of Cu-6Ni-2Mn-2Sn [6]. Han et al. found that the enhanced tensile strength with the addition of Al is attributed to the solid solution hardening effect [12]. Therefore, Al may sufficiently improve the strength of the low Sn-containing Cu-Ni-Sn alloys owing to grain refinement, CuAl2 nanoparticle precipitation and/or solid solution hardening. In the present work, a novel low Sn-containing Cu-Ni-Sn alloy with both high strength and good ductility was developed by adding minor Al and Si content. As mentioned above, the effect of Al on the mechanical prosperities of low Sn-containing Cu-Ni-Sn alloys were clearly studied, which is helpful for understanding the strengthening effect on the new alloy produced in this work. Furthermore, we found that the minor addition of Si can significantly improve the mechanical properties of Cu-Ni-Sn-Al alloy. Therefore, the effect Si on the microstructure and mechanical properties of a low Sn content Cu-9Ni-2.5Sn-1.5Al alloy was studied.
⁎ Corresponding author at: Guangdong Key Laboratory for Processing and Forming of Advanced Metallic Materials, South China University of Technology, Guangzhou 510640, China. E-mail address: [email protected] (W. Zhang).
https://doi.org/10.1016/j.msea.2018.12.120 Received 18 October 2018; Received in revised form 27 December 2018; Accepted 29 December 2018 Available online 31 December 2018 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Table 1 Chemical composition of the alloys. Alloy designation
Cu-9Ni-2.5Sn-1.5Al Cu-9Ni-2.5Sn-1.5Al-0.5Si
wt% Ni
Sn
Al
Si
Cu
9.17 9.72
2.53 2.56
1.38 1.34
̶ 0.45
Balance Balance
2. Processing and experimental procedure The alloys were prepared by electric melting of high purity Cu, Ni, Sn, Al and Si in an intermediate frequency induction furnace and then casting them into a steel mold with a diameter of 60 mm and a length of 160 mm. The chemical compositions of the alloys are summarized in Table 1. The ingots were homogenized at 930 °C for 2.5 h and machined to cylindrical-shaped samples with a diameter of 50 mm and a height of 50 mm for extrusion. Hot extrusion was performed at 950 °C with an extrusion ratio of 17 and a ram speed of 3 mm/s by using a 2000 KN vertical extruding machine. The as-extruded bars were cut into cylindrical tensile samples with a gauge section of 5 mm in diameter and 25 mm in length and the tensile tests were performed using a SANS CMT5105 material test machine in air with an initial tensile velocity of 1 mm/min at room temperature. At least three samples were tested under each condition. Optical micrograph (OM) was observed using a LEICA/DMI 5000 M optical microscope. The mean grain size was measured by the three round intercept method based on the observation of OM according to Chinese GB/T 6394-2002, more than five hundred grains were measured in each specimen. The volume fraction of the coarse unrecrystallized grains was obtained from 15 random areas based on the observation of OM. Texture analyses by EBSD were conducted using a field emission FEI NONA430 scanning electron microscope operated at 30 kV. Samples for EBSD measurement were prepared by mechanical polishing with progressively finer water-based diamond suspension (0.5 µm) followed by electro-polishing using a solution of 50% phosphoric acid and 50% alcohol at 6.5 V for 60 s. TEM specimens were prepared by twin jet electro-polishing in a solution of 30% nitric acid and 70% methanol held at −20 to −30 °C. The specimens were finished for surface cleaning by ion milling using a Gatan precision Ion Polishing System (PIPS) at an operating voltage of 3 kV for 20 min. Microstructure analyses were performed on a JEM-2100F transmission electron microscope. The area fraction of precipitates was measured from five different views of central dark-field TEM micrograph.
Fig. 1. Typical tensile stress-strain curves of as-extruded 0Si and 0.5Si alloys. Table 2 Tensile properties of 0Si and 0.5Si alloys. Alloy
σb (MPa)
σ0.2 (MPa)
Ɛ (%)
0Si 0.5Si
713 ± 10 861 ± 7.8
453 ± 7.5 698 ± 13
23 ± 0.3 18 ± 1.1
Fig. 2. Comparison between the present alloys and other typical Cu-15Ni-8Sn alloys with the tensile strength versus elongation [4,13–17].
3.2. Structural characterization 3. Results and discussion
Fig. 3 shows the microstructure of the dynamically recrystallized grains in the as-extruded 0Si and 0.5Si alloys. The 0Si alloy has uniform and equiaxed dynamic recrystallized grains with an average diameter of 46 ± 2.2 µm (Fig. 3a and c). While the 0.5Si alloy shows a bimodal grain structure consisting of fine dynamic recrystallized grains with an average diameter of 34 ± 4.3 µm and elongated coarse unrecrystallized grains (marked as red arrows in Fig. 3b and d), which is much different from the microstructure of the 0Si alloy. Such microstructure change results into a large difference in mechanical properties, as shown in Fig. 1. This bimodal grain structure was often found in the hot-extruded alloys [18–20], in which the equaixed grains were resulted from the recrystallization started from the grain boundaries and the elongated grains were formed during hot extrusion. Fig. 4 shows the TEM micrographs and the SADPs of the homogenized and as-extruded 0Si alloy. It indicates that fine γ′-Ni3Al particles precipitated in the matrix, and the crystal orientation relationship between Ni3Al and Cu matrix is [100]γ′//[200]Cu, [010] γ′//020]Cu and [001] γ′//[001]Cu. Fig. 5 shows the microstructure of the homogenized and as-extruded 0.5Si alloy. The homogenization treated sample shows
3.1. Mechanical properties Fig. 1 shows tensile properties of the as-extruded Cu-9Ni-2.5Sn1.5Al (marked as 0Si alloy) and Cu-9Ni-2.5Sn-1.5Al-0.5Si (marked as 0.5Si alloy) alloys. The ultimate tensile strength (σb), 0.2% offset yield stress (σ0.2) and strains to failure in tension (Ɛ) are summarized in Table 2. It is clear that the 0.5Si alloy has a much higher strength than the 0Si alloy. The yield and tensile strength of the 0.5Si alloy are 698 MPa and 861 MPa, respectively, while for 0Si alloy they are 453 MPa and 713 MPa, respectively. Although the elongation decreases slightly, the 0.5Si alloy still possesses good ductility of more than 18% plastic strain. Fig. 2 shows the comparison between the present alloys and other typical Cu-Ni-Sn alloys with the tensile strength versus elongation, indicating that the 0.5Si alloy, a low Sn content Cu-Ni-Sn alloy produced in this work, has a good combination of high tensile strength and good ductility.
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Fig. 3. OM images of the as-extruded 0Si and 0.5 Si alloys: (a and b) the plane perpendicular to the extrusion direction (ED); (c and d) the plane parallel to the extrusion direction (TD). The average diameter of dynamic recrystallized grains is given as d d in Fig. 3(a) and (b).
Fig. 4. TEM micrographs and SADPs of 0Si alloy: (a and b) homogenized at 930 °C for 2.5 h; (c and d) hot extruded at 950 °C. 156
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Fig. 5. TEM micrographs of 0.5Si alloy. (a): homogenized at 930 °C for 2.5 h and (b-e): hot extruded at 950 °C; (c) the HRTEM image of the rectangular area in (b) and the inset is the FFT of the red rectangle area in (c); (d) and (e) is bright-field and central dark-field image, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
Fig. 6. The elemental mapping images of the as-extruded 0.5Si alloy.
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Fig. 7. EBSD results showing IPF maps and local misorientation maps for 0Si alloy (a and b) and 0.5Si alloy (c and d), black lines denote high angle grain boundaries (HAGBs, > 10), red lines denote SAGBs (1°~10°); (e): misorientation profiles measured along the yellow arrow in (d), the black line shows point-to-point misorientation, while the red line shows point-to-origin misorientation respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
radians with a distance δ, and b is the Burges vector [24]. Taking b = 0.2556 nm for pure Cu, with values for θ and δ deriving from misorientation profile in Fig. 7e, the estimated local dislocation densities is 0.5–1 × 1013 m−2. Fig. 8 shows the {111} < 110 > Schmid factor distribution maps of the 0.5Si alloy. It shows that unrecrystallizaiton regions have a darker color indicating a lower Schmid factor (Fig. 8a). The average Schmid factor of the {111} < 110 > is 0.315 and 0.436 for the unrecrystallizaion regions (Fig. 8b) and recrystallization regions (Fig. 8c), respectively, indicating that high strength was achieved in the unrecrystallization regions. Fig. 9 shows the distribution of coincidence site lattice (CSL) boundaries and grain boundaries, in which random high-angle grain boundaries and Σ3 boundaries are denoted by the black lines and red lines, respectively. Σ3 boundaries occurred frequently in the microstructure with numerous long and parallel-side annealing twin boundaries, which are observed in recrystallization regions of both 0Si and 0.5Si alloys. However, no Σ3 boundary is observed in the unrecrystallization regions of the 0.5Si alloy (Fig. 9b). The proportions of Σ3 boundaries and HAGBs (55˚–60˚) are decreased, while SAGBs (≤5˚) are increased obviously with the addition of Si. The average number of twin grains within a parent grain was calculated by determining the number of grains in the region both including and excluding twin boundaries as defining individual grains. The calculated value of the 0Si and 0.5Si alloy is 3.17 and 2.70, respectively. During recrystallization, when growth stagnation occurred, the boundary between the recrystallizing grain and the deformed matrix was often in a region of low dislocation density and/or low misorientation angle [25]. The form of Σ3 twin boundaries had initiated new orientations without the rotation of new grain with respect to the parent grain and apparently offers the additional boundary energy required to facilitate the continued grain growth [26,27]. It means that the density of Σ3 twin boundary may be decreased when the grain growth was retarded. The 0.5Si alloy shows a lower density of Ʃ3 twin boundary than the 0Si alloy owing to that Ni3Si nanoparticles were formed with the addition of Si which can suppress the grain growth.
that a few fine precipitates in the matrix (Fig. 5a), while the hot-extruded sample shows that a high density fine precipitates (20 ± 3 nm) distributed in the Cu matrix and along the grain boundary (Figs. 5b-5e), which were identified as β-Ni3Si from the diffraction pattern (the insert in Fig. 5c). β-Ni3Si was commonly observed in Cu-Ni-Si alloys [21,22], which has a L12 ordering structure with structure parameters of a = 0.351 nm, Pm-3 m(221) [23]. The crystal orientation relationship between Ni3Si and Cu matrix is [100]β//[200]Cu, [010]β//020]Cu and [001]β//[001]Cu. The precipitate-free zones (PFZ) with ~227 nm width were observed on the both sides of the grain boundary according to Figs. 5d and 5e, which is further confirmed by the element mapping (Fig. 6). It indicates that Sn, Al and Si are uniform distributed while Ni is significantly depleted near the grain boundaries. The Ni3Si nanoparticles in the grain boundaries retarded the recrystallization process and kept the coarse grains away from recrystallization. Therefore, the bimodal structure shown in as-extruded 0.5Si alloy is the result of the competition between recrystallization (heating process) and non-recrystallization (recrystallization suppressed by the Ni3Si precipitates in the grain boundaries). Fig. 7 shows the texture in the as extruded 0Si and 0.5Si alloys. Their inverse pole figure (IPF) in the form of contour plot shows the texture intensity. The 0Si alloy shows relatively randomly oriented grains with < 111 > fiber texture which is a typical texture in Cu alloys after extrusion (Fig. 7a), and there are almost no local strains in the grains implying that grains are free of dislocations after recrystallization (Fig. 7b). While the 0.5Si alloy shows a mixture of relatively randomly equiaxed grains and the unrecrystallization grains with a strong < 111 > fiber texture (Fig. 7c). The local misorientation maps suggest that there are almost no local strains in the recrystallization regions of 0.5Si alloy (Fig. 7d). But in the unrecrystallization regions, a large number of small angle grain boundaries (SAGBs, 1°–10°) are observed. Unrecrystallized grains keep high dislocation density, which can pin the dislocation movement and result into the formation of SAGBs. Generally, the larger the long-range misorientation gradient in the same grain, the higher the dislocation density is. The equivalent geometrically necessary dislocation density can roughly be estimated by ρ ≈ θ/(bδ), where θ is the accumulated misorientation angle in 158
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Fig. 8. The {111} < 110 > Schmid factor distribution maps of the 0.5Si alloy (a), and the corresponding regions of unrecrystallized grains (b) and recrystallized grains(c); (d) and (e) is the {111} < 110 > Schmid factor distribution obtained from Fig. b and Fig. c, respectively.
grain size) which shows a good agreement with the experiment data in Cu-Al alloys [33]. In present study, the twin boundaries were considered as grain boundaries and the strengthening effect was roughly calculated by a modified Hall-Petch equation. Therefore, the yield strength increment (σGTBS) due to fine grain and twin boundaries can be estimated by Hall-Petch relationship [35]
3.3. Strengthening mechanisms The low Sn-containing Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy shows a super mechanical property with a high strength of 861 MPa and a large elongation of 18%, which is comparable to the high Sn-containing alloy such as Cu-15Ni-8Sn. The strengthening effect can be explained by the following factors: (i) Grain boundary strengthening (ii) Texture strengthening; (iii) Orowan precipitation strengthening; (iv) Others strengthening, such as strengthening by twins and solid solution. Parameters used in the yield strength calculations were shown in Table 3.
−1/2
df ⎡ σGTBS = σ0 + k ⎢Vf ⎜⎛ ⎟⎞ n ⎣ ⎝ ⎠
⎤ + Vc dc−1/2 ⎥ ⎦
(1)
σ0 is the fractional stress, k is slope, df and dc is the grain size of fine recrystallized grains and coarse unrecrystallized grains, respectively. Vf and Vc is the volume fraction of fine recrystallized grains and coarse unrecrystallized grains, respectively. n is the average number of twin grains within a grain of 0.5Si alloy, which was only considered in recrystallization regions. The value of σGTBS was calculated to be 142 MPa. (ii) Texture strengthening The texture is the result of both of hot extrusion and dynamic recovery. The texture strengthening mechanism is familiar with work
(i) Grain and twin boundary strengthening The microstructure of the as-extruded alloy consisted of a bimodal grain structure of fine recrystallized grains and coarse unrecrystallized grains. The grain boundary strengthening effect was taken into account for the enhanced strength in the fine recrystallized grains. In addition, it has reported that deformation twin boundaries can act as effective barriers to dislocation movements [32–34]. The additional strengthening contribution by deformation twins has been calculated by a modified Hall-Petch equation (using inter spacing between the twins as the “effective” 159
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Fig. 9. EBSD orientation maps showing the microstructural topology in (a) 0Si alloy, (b) 0.5Si alloy. (c) and (d) is CSL boundaries and misorientation angle, respectively. Black lines denote random high angle boundaries; red lines denote Σ3 boundaries.
hardening. The yield strength due to work hardening as given by Eq. (2) [36,37].
ΔσTS = MαGb ρ
τp = 0.81
(2)
ln(dp/ b) Gm b 2π (1 − ν )1/2 λ − dp
(3)
where b is the Burgers Vector, G is the shear modulus of the matrix, v is the Poisson’s radio, dp is the average radius of the precipitates, and λ is the spacing between particles in the glide plane. And λ can be calculated in Eq. (4) [39]:
where M is Taylor factor, α is a geometric constant (about 0.3), G is the shear modulus of the copper matrix, b is the Burgers vector, ρ is the dislocation density. The texture strengthening was only considered in unrecrystallization regions, and the yield strength increase from texture strengthening was 8 MPa. (iii) Orowan precipitation strengthening
λ=
1 dp 2
3π 2fp
(4)
where fp is the volume fraction of second-phase particles. The incremental strength increased by Orowan mechanisms can be calculated as [39]:
Orowan precipitation strengthening mechanism can be described by the following Orowan-Ashby equation [38]: Table 3 Parameters used in the yield strength calculations. Parameter
Description
Value
Units
Refs.
σ0 k Gm b dp fp Vc dc ρ M
Fractional stress A constant for Cu-6Sn Shear modulus of copper Burgers vector of copper The average radius of the Precipitates (0.5Si alloy) Volume fraction of the precipitates Volume fraction of coarse unrecrystallized grains Average diameter of coarse unrecrystallized grains Dislocation density in coarse unrecrystallized grains Taylor factor
60 356 46 0.2556 20 8.85% 26.8% 118 0.75 × 1013 3.06
MPa MPa/μm−1/2 GPa nm nm – – μm m−2 –
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Fig. 10. Yield strength increase calculations from different strengthening mechanisms.
ΔσOPS = Mτp
(5)
where M is Taylor factor. The calculated result showed that the increase strength caused by Orowan mechanisms was 470 MPa for 0.5Si alloy. The total yield strength strengthened by different strengthening mechanisms were shown in Fig. 10, and the Orowan precipitation strengthening plays the main role with 67.3% of total strengthening contribution. 4. Conclusions A low Sn-containing Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy was developed by casting and hot extrusion, showing a super mechanical property with a high strength of 861 MPa and a large elongation of 18%, which is comparable to the high Sn-containing alloy such as Cu-15Ni-8Sn. The addition of 0.5 wt% Si in the Cu-9Ni-2.5Sn-1.5Al alloy can significantly change the microstructure and mechanical properties. A bimodal microstructure was formed in the Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy, which has coarse unrecrystallization regions having a plenty of < 111 > fiber texture and fine equaixed recrystallization grains. The bimodal microstructure in the Cu-9Ni-2.5Sn-1.5Al-0.5Si alloy give rises to super combination of mechanical property such as high strength and good ductility. The enhanced strength is contributed by: (i) grain refinement and twin boundary strengthening; (ii) precipitates strengthening; (iii) texture strengthening. Acknowledgements This work was sponsored by the Guangdong Natural Science Foundation for Research Team (Grant No. 2015A030312003) and Guangdong Special Found Project of Applied Technology Research and Development (Grant No. 2016B090931002) References [1] Y. Zhang, Z. Xiao, Y. Zhao, Effect of thermo-mechanical treatments on corrosion behavior of Cu-15Ni-8Sn alloy in 3.5 wt% NaCl solution, Mater. Chem. Phys. 199 (2017) 54–66, https://doi.org/10.1016/j.matchemphys.2017.06.041. [2] J.B. Singh, W. Cai, P. Bellon, Dry sliding of Cu–15 wt% Ni–8 wt% Sn bronze: wear behaviour and microstructures, Wear 263 (2007) 830–841, https://doi.org/10.1016/j. wear.2007.01.061. [3] J. Caris, D. Li, J.J. Stephens, Microstructural effects on tension behavior of Cu–15Ni–8Sn sheet, Mater. Sci. Eng. A 527 (2010) 769–781, https://doi.org/10.1016/j.msea.2008.01. 061. [4] C. Zhao, W. Zhang, Z. Wang, Improving the mechanical properties of Cu-15Ni-8Sn alloys by addition of titanium, Materials 10 (2017), https://doi.org/10.3390/ma10091038. [5] M.J. Diánez, E. Donoso, M.J. Sayagués, The calorimetric analysis as a tool for studying the aging hardening mechanism of a Cu-10 wt%Ni-5.5 wt%Sn alloy, J. Alloy. Compd. 688 (2016) 288–294, https://doi.org/10.1016/j.jallcom.2016.07.021.
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