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Investigation on the microstructure and mechanical properties of the spray-formed Cu–Cr alloys
Materials Science and Engineering A 460–461 (2007) 69–76
Investigation on the microstructure and mechanical properties of the spray-formed Cu–Cr alloys Xiaofeng Wang ∗ , Jiuzhou Zhao ∗∗ , Jie He Institute of Metal Research, Chinese Academy Science Wenhua Road 72, Shenyang 110016, PR China Received 8 December 2006; received in revised form 2 January 2007; accepted 2 January 2007
Abstract Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys were prepared by spray forming. The as-deposited alloys were subsequently warm rolled at 300 ◦ C with a 40% reduction and then two types of thermo-mechanical treatments were adopted to enhance the mechanical properties of the Cu–Cr alloys. The microstructural features of the Cu–Cr alloys in different thermo-mechanical processing conditions were characterized using optical, scanning electron and transmission electron microscopy techniques. The results show that the spray-formed Cu–Cr alloys exhibit a better response to heat treatment compared to the conventional casting alloys. The chromium content has a great effect on the aging behaviors of the Cu–Cr alloys. Although a higher chromium content leads to a larger volume fraction of chromium precipitates generated during the solidification, it does not cause an increase in the mechanical properties. The chromium content should be lower than 1.2 wt.%. © 2007 Elsevier B.V. All rights reserved. Keywords: Cu–Cr alloys; Spray forming; Strength; Microstructure; Thermo-mechanical processing
1. Introduction Due to having the combination of high strength, good electrical and thermal conductivities, Cu–Cr alloys have been applied in many engineering fields, such as in electric resistance welding electrodes, the liner tube of continuous casting crystallizer, the integrated circuit lead frame, the aerial conductor of electric locomotive and so on [1–4]. Recently, the higher strength is more and more required for copper alloys with the rapid development of the electronic industry. For example, a tensile strength σ b > 600 MPa, a hardness Hv > 180 and a conductivity > 80%IACS are demanded for lead frame materials on a very large scale integrated circuit [5,6]. The Cu–Cr alloys prepared by the conventional casting, however, cannot satisfy these properties indices. Spray forming process is especially useful for the preparation of bulk materials with the characteristics of rapid solidification. Recently, many investigations have been reported on the high strength Cu–Cr alloys prepared by spray forming [7–9]
∗
Corresponding author. Tel.: +86 24 23971905; fax: +86 24 23971918. Corresponding author. E-mail addresses: [email protected] (X. Wang), [email protected] (J. Zhao). ∗∗
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.01.023
and the spray-formed Cu–Cr alloys showed excellent mechanical properties. But the work mainly focused on the Cu–Cr alloys which have the chromium content lower than 1 wt.% [8]. Since the chromium content has a great effect on the microstructural features and the mechanical properties of the spray formed Cu–Cr alloys, in the present investigation, Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys were successfully prepared by spray forming and the effect of chromium content on the microstructure and the mechanical properties of the Cu–Cr alloys was discussed.
2. Experimental details A schematic of the spray atomization and deposition process is shown in Fig. 1. The nominal compositions of the alloys were Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr. The Cu–Cr alloys were melted in the vacuum induction furnace. The melt was homogenized at 1230–1250 ◦ C for 1 h prior to the spray deposition. The detail of the set-up has been described elsewhere [10]. The delivery tube size was 4.5 mm. Nitrogen, at an atomizing pressure of 1.5 MPa, was used as the atomizing gas. The spray jet was deposited onto a substrate at a distance of 450 mm from the nozzle. The rotation and axial movement of the substrate guaranteed a constant deposition distance of 450 mm throughout the depo-
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Fig. 1. Schematic of spray atomization and deposition process.
sition time. The sizes of the spray formed billets were 120 mm in height and 150 mm in diameter. The billets were machined and pieces of 100 mm in diameter were removed from the central region of the deposits for the further investigation. After warm rolled at 300 ◦ C with a 40% cross-reduction, the Cu–Cr alloys were subsequently subjected to different types of thermo-mechanical treatments: (1) aging of the Cu–Cr alloys after cold rolled with 60% reduction (SF + WR + CR + aging); (2) aging of the Cu–Cr alloys after solution treatments and cold rolling with 80% reduction (SF + WR + ST + CR + aging). The solution treatment of the alloy in schedule (2) was carried out at 970 ◦ C for 30 min in a nitrogen atmosphere followed by water quenching. All the aging treatments were carried out at temperatures in the range of 350–500 ◦ C in a nitrogen atmosphere. The microstructure of the alloys in various states was characterized by means of optical microscopy (Leica MEF-4), scanning electron microscopy (SS-550) and transmission electron microscopy (Philips-CM12). The tensile properties were determined using an Instron-testing machine AG-5000A. The cross-head speed is 1 mm/min. 3. Results 3.1. Microstructural features An equiaxed grain morphology is distinctly observed in the spray formed alloys, as shown in Fig. 2a for Cu–1.2%Cr and Fig. 2b for Cu–3.2%Cr alloys. This feature was invariably present in the specimens taken from the various regions of the as-deposited Cu–Cr alloys. A bimodal size distribution of the primary chromium particles generated during solidification was observed in the as-sprayed Cu–3.2Cr alloy, as shown in Fig. 2b. The large particles about 1–3 m in diameter dispersed mainly on the grain boundaries coexisting with smaller chromium par-
Fig. 2. Optical micrographs of as-deposited Cu–Cr alloys: (a) Cu–1.2 wt.%Cr and (b) Cu–3.2 wt.%Cr.
ticles of ∼0.5 m within the grains. Occasionally, some big Cr particles can be found on the grain boundaries. These features were not clearly observed in the as-deposited Cu–1.2 wt.%Cr. Fig. 3 shows the TEM image of as-deposited Cu–1.2 wt.%Cr alloy. It can be found that there are some small precipitates of 10 nm in diameter dispersed in the matrix. It can be assumed that aging process took place due to the slow cooling rate of the as-deposited Cu–Cr alloy at the cooling stage. The average grain size of Cu–1.2 wt.%Cr alloy was measured, using the linear intercept method, to be about 50 ± 3.2 m. While the grain size of as-deposited Cu–3.2 wt.%Cr alloy was ∼30 ± 3.0 m, which is much smaller than that of Cu–1.2 wt.%Cr alloy. The difference of the grain size was attributed to the pinning effect of the primary chromium particles that hinder the grain growth. Cold rolling of the spray-formed alloy generated a large number of deformations. The equiaxed grains were squeezed to lamellar and the pores in the deposit were considerably reduced, see Fig. 4. For Cu–3.2 wt.%Cr alloy (Fig. 4b), undissolved Cr (at most 1 wt.% Cr can be dissolved in the solid state) remained as micron-sized particles, which were elongated as Cr fibers during heavy cold working. It has been reported that such fibers can lead to significant strengthening when present in a large volume faction [11,12]. But in our investigation, such fibers are not
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Fig. 3. TEM image showing the microstructure of as-deposited Cu–1.2 wt.%Cr alloy.
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expected to play a major role here where at most about 2% Cr fibers are present. Fig. 5 shows the microstructure of Cu–1.2 wt.%Cr alloy after the SF + WR + ST + CR treatment. A large number of the dislocations within the grains were observed which arises from the high degree of cold deformation. These dislocations can act as diffusion paths for solute atoms and provide nucleation site for precipitation during aging treatment. Deformation twins were also found in the matrix. The twins seemed to form the bundles of a few micrometers thick and several hundred nanometers long. Within the bundles, twin boundaries were not straight along its length. Most twins’ boundaries curved, which is the typical characteristics of the deformation twins [13]. This phenomenon may originate from the fact that a high density of dislocation exists at the boundaries of the twins, as reported in the severely deformed copper alloys [14,15]. However, there was no evidence of this phenomenon in the cold-rolled alloys with 40% reduction. The age hardening behavior is greatly influenced by the initial microstructural features of the alloy that are developed during various processing stages. The results of the present investigation clearly illustrate that the alloy processed by spray forming exhibits a better response to the heat treatment compared to the conventional casting alloys. Fig. 6a shows the TEM micrograph of the Cu–1.2 wt.%Cr alloy after aging treatment at 450 ◦ C for 1 h. The homogeneously distributed globular precipitates of about 5 nm in diameter can be found. Many investigations showed that for rapidly solidified Cu–Cr alloy the precipitates, corresponding to the peak strength or hardness upon aging, are probably coherent ones, which were proved by the fact that the
Fig. 4. Optical microstructure of as-rolled spray-formed Cu–Cr alloys: (a) Cu–1.2 wt.%Cr; (b) Cu–3.2 wt.%Cr. The cold rolling reduction is 80%.
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Fig. 5. TEM morphology of the dislocations and twins for the spray-formed Cu–1.2 wt.%Cr alloy with SF + WR + ST + CR treatment. The cold rolling reduction is 80%.
two-lobe contrast was clearly observed in the following stages of the aging process. However, in the present work, no obvious twolobe contrast caused by coherent distortion was observed in the microstructure. This phenomenon is probably due to the smaller size of the precipitates in the present work. Usually, the two-lobe contrast can only be observed when the size of the precipitate is larger than 30 nm. For example, Liu et al. [5] have reported that the appearance of the two-lobe contrast corresponding to the precipitates with the size about 40–50 nm in a rapidly solidified Cu–Cr alloy. Furthermore, Mclntype et al. [16] found that the two-lobe type strain field and its shape mainly depended on r3 (r is the radius of the precipitate). With the increase of the aging temperature, the volume fraction of precipitates increase slightly, but the sizes of the precipitates did not increase significantly. The selected area electron diffraction (SAED) analysis of the microstructure corresponding to the peak hardness did not show any sign of the bcc Cr phase. It is obvious that the coherent precipitates are not the bcc Cr phase. According to the previous investigation [17], the coherent precipitate should be the fcc Cr phase. Upon further aging, the fcc coherent precipitate lost its coherency with the matrix and transformed to a bcc structure. Fig. 6b shows the TEM microscopy of the Cu–3.2 wt.%Cr alloy after SF + WR + ST + CR + aging at 300 ◦ C for 1 h. This result indicates that the precipitates of an average size of 50 nm uniformly distributed in the matrix. Compared with that observed in Cu–1.2 wt.%Cr at its peak aging (Fig. 6a), the precipitates in this alloy are much bigger. These precipitates seem to lose the coherency with the matrix. With the aging temperature increased or the aging time prolonged, the particles grow quickly and get coarser.
Fig. 6. TEM morphology of precipitates: (a) Cu–1.2 wt.%Cr (SF + ST + CR + aging at 450 ◦ C for 1 h); (b) Cu–3.2 wt.%Cr (SF + ST + CR + aging at 300 ◦ C for 1 h).
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Fig. 7. TEM image of microstructure of Cu–Cr alloys showing small grains. (a) Warm rolled, solution treated, cold-rolled Cu–1.2 wt.%Cr alloy aged at 450 ◦ C for 1 h and (b) warm rolled, solution treated, cold-rolled Cu–3.2 wt.%Cr alloy aged at 450 ◦ C for 1 h.
It has been observed that the Cu–1.2 wt.%Cr alloy at its peak aging exhibited a larger number of the small grains of 100 nm in size, as shown in Fig. 7a. This indicates that the recrystallization occurs during the aging and coarsening process of these small grains is very slow. However, this phenomenon is more evident for the Cu–3.2 wt.%Cr alloy, as shown in Fig. 7b. Even though this alloy (after warm rolled, solution treated and cold rolled with 80% reduction) was aged at 300 ◦ C for 1 h, the recrystallization took place quickly in this alloy and with the aging temperature increased or the aging time prolonged, abnormal grain growth can be found (see Fig. 7b). 3.2. Mechanical properties The variations of the tensile strengths of the Cu–Cr alloys with the aging temperature and the aging time are given in Fig. 8. It can be found that the strength of the spray-formed Cu–Cr alloy is strongly dependent on the thermo-mechanical treatment. For the Cu–3.2 wt.%Cr alloy, the strength increment resulting from heat treatment was usually small. The highest strength is about 530 MPa, which can be obtained after SF + WR + CR + aging at 300 ◦ C for 1 h. It is concluded that SF + WR + ST + CR aging treatment cannot result in the strength improvement for the spray-formed Cu–3.2 wt.% alloy. In contrast, the strength of the spray-formed Cu–1.2 wt.%Cr alloy is strongly dependent on the thermo-mechanical treatment. With the increase of aging temperature, the tensile strength of
the spray-formed Cu–1.2 wt.%Cr alloy increases quickly from 460 MPa at 350 ◦ C to 585 MPa at 450 ◦ C and then decreases. The maximum tensile strength, 585 MPa, was obtained for the samples of Cu–1.2 wt.%Cr alloy after SF + WR + ST + CR + aging at 450 ◦ C for 1 h. The tensile strength of this alloy after peak aging treatment is much higher than that of the Cu–3.2 wt.%Cr alloy. The tensile strengths of the Cu–Cr alloys were measured with the variations of the aging time at the peak aging temperature. The Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys indicate the peak strengths of 585 and 530 MPa after aged at 450 and 300 ◦ C for 1 h, respectively. Subsequently, the strength value decreases slowly with the aging time prolonged and reaches to the values of 520 and 460 MPa, respectively, when the aging time is 4.5 h. Fig. 9 shows the SEM morphology of the fracture surface. Many equiaxed dimples appeared at the fractured surface. The sizes of the dimple for Cu–3.2 wt.%Cr alloy is bigger than those for Cu–1.2 wt.%Cr alloy. The difference of the fracture surface may be attributed to the effect of the primary chromium particles. The process of crack is readily suggested to be the plastic deformation of the matrix initiating nearby the large particles and then extending when the plastic strain increases. Cracking takes place after the chromium particles are torn out of the matrix under a critical strain. Therefore, the strength of the Cu–Cr alloy can be further improved if the amount and size of the primary chromium particles are reduced with an optimized processing parameters.
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Fig. 8. Variation of the tensile strength of Cu–Cr alloys with aging temperature (a) and aging time (b) at different processing conditions. Superscript 1 represents SF + WR + CR + aging treatment and superscript 2 represents SF + WR + ST + CR + aging treatment. Fig. 9. SEM morphology of the fracture surface: (a) Cu–1.2 wt.%Cr and (b) Cu–3.2 wt.%Cr.
4. Discussions The spray forming process involves the atomization of an alloy melt into a jet of droplets and their subsequent deposition onto a substrate. The droplets experience a high cooling rate of 103 –106 K s−1 [7–9], which therefore renders a rapid solidification features in the spray formed microstructure. The spray formed alloy gives rise to a refined and equiaxed grain morphology in contrast to the dendritic coarse structure, with severe microsegregation, observed in the conventional casting alloy. In addition, the high cooling rate drives the alloy away from the equilibrium condition, and as a result the limit of solid solubility of the alloying elements increases. This induces a better response of the spray formed alloy during aging treatments. However, despite a high cooling rate during atomization, it is difficult to suppress the second phase precipitation as the cooling rate becomes slow after deposition. For the Cu–1.2 wt.%Cr alloy, most of the chromium element remains in solid solution. Only a small part of chromium precipitates on the grain boundaries. On the contrary, for the Cu–3.2 wt.%Cr alloy, more than 2 wt.% chromium precipitate from the matrix during the solidification. A considerable amount of the primary particles has a great effect on the precipitation behavior. Deformation during cold rolling is hindered by these precipitates and therefore
leads to the formation of more geometrically necessary dislocations which aid in a faster rate of precipitation. During aging treatment, Cr solute can easily precipitate from the matrix. Due to the fast diffusion rate, the precipitates grow quickly, coarsen and lose coherent relationship with the matrix, and meanwhile the primary precipitates also coarsen rapidly. All these factors induce the Cu–3.2 wt.%Cr alloy not to have a distinct aging peak. In contrast, for Cu–1.2 wt.%Cr alloy, when most of the chromium remained in solid solution, the absence of precipitates within the grains allows relatively unconstrained flow. During aging, Cr solute can homogeneously precipitate from the matrix and the precipitates keep coherent relationship with the matrix, which results in the enhancement of the strength. In the peak hardness or slight over-aging conditions, the precipitates remain coherent with the matrix. The shear stress increase due to the coherency strengthening σ CS can be calculated by the following equation [18]: rfb 1/2 σCS = MχG(ε)3/2 (1) Γ where ε is the function of the degree of misfit δ, G and b the shear modulus and Burgers vector of the matrix, respectively, r and f the average radius and the volume fraction of the precipitates,
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respectively, Γ the line tension of the dislocations pinned by the precipitates, and χ is a coefficient varying between 2 and 3. Usually, χ is taken as 2.6 [19]. Γ = aGb2 , a varies between 0.089 and 0.5. M is the Taylor factor, with M = 3.06 [20–22]. For the Cu–3.2 wt.%Cr alloys, most of the precipitates are incoherent. The strengthening mechanism usually proposed for the strengthening beyond the peak strength (over-aged alloys having larger and incoherent precipitates) is the Orowan-Ashby model and σ OA can be calculated by the following equation [23]: πr 0.84MGb 1 ln σOA = 2 (2) × 2r0 π (1 − v)1/2 r[8/(3πf )1/2 − 1] where σ OA is the yield strength increase due to the OrowanAshby model and r0 is the inner cut-off radius of the dislocation. The curves corresponding to the variation of σ CS and σ OA with the precipitate volume fraction are presented in Fig. 10. These curves were obtained by substituting the relevant factors and constants (G = 42.1 GPa, v = 0.303, r0 = 3b, b = 0.25 nm for the copper matrix) in Eqs. (1) and (2). Obviously, the strength due to the coherent precipitates is much higher than the strength calculated through the Orowan-Ashby model. With the increase of the coherency precipitates size, the strength increases quickly. On the contrary, the strength calculated through the Orowan-Ashby model decreases with the precipitates increased. Even though the volume fraction of the precipitates in the Cu–1.2 wt.%Cr alloy is about 1.3%, which is much lower than that in the Cu–3.2 wt.%Cr alloy (∼3.3%), the strength of the Cu–1.2 wt.%Cr alloy is much higher than the Cu–3.2 wt.%Cr alloy due to the coherency effect. It appears that the solid solubility of chromium in copper reaches a maximum in the Cu–1.2 wt.%Cr alloy. As the chromium content was lower than 1.2 wt.%, the strength was enhanced with the increase of the Cr content. Beyond this level, although a higher content of chromium lead to a larger volume fraction of chromium precipitates, it does not cause an increase in the mechanical properties. Therefore the selection of the chromium content should be lower than 1.2 wt.%.
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The recrystallization process of the alloys depends on the particle distribution and sizes. At the initial stage of recrystallization, the crystal nuclei, almost free of strain, might grow towards the deformed matrix through migration of large angle boundaries to decrease the strain energy. The driving force for recrystallization FN can be expressed as [24–26]: FN = aGb2 (ρ0 − ρl )
(3)
where a is a constant, G the shear elastic modulus, b Burgers vector, ρ0 and ρl are the dislocation density after deformation and recrystallization, respectively. The above-mentioned results showed that more coarse particles result in higher total strain during cold deformation and consequently the dislocation density in Cu–3.2 wt.%Cr alloys is much higher than that in Cu–1.2 wt.%Cr alloy. The driving force for recrystallization in Cu–3.2 wt.%Cr alloy is about 13.16 MPa and for Cu–1.2 wt.%Cr alloy, this value is about 1.32 MPa. This suggests that the primary particles are responsible for concentrating the strain leading to an increase in driving force for recrystallization. This acceleration is more pronounced during the early stages of aging when the formation of new grains is taking place. For Cu–3.2 wt.%Cr alloy, the recrystallization occurred at 300 ◦ C, which is much lower than that of Cu–1.2 wt.%Cr alloy. The pinning force that the particles produce per unit area of boundary FV can be calculated from the equation [7]: 3fv σgb (4) D where fv is the volume fraction of particles, σ gb is the interface energy per unit area (0.5 J cm−2 for Cu) and D is the average particle diameter. The values obtained when considering the finer particles only are 4.05 and 0.1 MPa for Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloy, respectively. Therefore, the finer particles in Cu–1.2 wt.%Cr produce a pinning force on the boundaries which is nearly forty times larger than that in the Cu–3.2 wt.%Cr alloy. Therefore, the finer particles play an important role in controlling the grain size during recrystallization.
FV =
5. Conclusions The effects of the chromium content and the thermomechanical treatment on the microstructure and mechanical properties of spray-formed Cu–Cr alloys were investigated. The conclusions are as follows:
Fig. 10. Calculated strengthening vs. precipitate volume fraction due to coherency strengthening σ CS for Cu–1.2 wt.%Cr alloy and due to the OrowanAshby model σ OA for Cu–3.2 wt.%Cr alloy.
(1) In the as-deposited state, Cu–Cr alloys exhibit a fine equiaxed grain structure with no macrosegregation. Grain size decreases and the amount of primary precipitates increases as the chromium content increases. Some finer Cr particles precipitated from the matrix due to the slow cooling rate of the billet during the cooling stage. Finer particles play an important role in controlling the grain size during recrystallization. (2) The spray-formed Cu–Cr alloys show excellent deformability. After rolling, aging treatment can be adopted to enhance the mechanical properties of the Cu–Cr alloys. The spray-formed Cu–Cr alloys exhibit a better
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response to heat treatment compared to the conventional casting alloys. The aging process of the Cu–Cr alloys becomes easier with the increase of the chromium content due to the effect of the primary Cr particles. Cu–1.2 wt.%Cr alloy exhibits excellent mechanical properties after proper warm rolling + solution treatment + cold rolling + aging treatment. The thermo-mechanical treatment of warm rolling + solution treatment + cold rolling + aging was not suitable for the spray-formed Cu–3.2 wt.%Cr alloy. (3) Coherency strengthening is the main strengthening mechanism. Attractive combination of strength and ductility can be achieved by the thermo-mechanical processing of the sprayformed Cu–1.2 wt.%Cr alloy. Excess chromium would not lead to the increment of the mechanical properties. Therefore, the selection of the chromium content should be lower than 1.2 wt.%. Acknowledgements The authors gratefully acknowledge the financially support from the National Natural Science Foundation of China (50620130095 and 50671111) and the Natural Science Foundation of Liaoning Province of China (20050047). References
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Investigation on the microstructure and mechanical properties of the spray-formed Cu–Cr alloys Xiaofeng Wang ∗ , Jiuzhou Zhao ∗∗ , Jie He Institute of Metal Research, Chinese Academy Science Wenhua Road 72, Shenyang 110016, PR China Received 8 December 2006; received in revised form 2 January 2007; accepted 2 January 2007
Abstract Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys were prepared by spray forming. The as-deposited alloys were subsequently warm rolled at 300 ◦ C with a 40% reduction and then two types of thermo-mechanical treatments were adopted to enhance the mechanical properties of the Cu–Cr alloys. The microstructural features of the Cu–Cr alloys in different thermo-mechanical processing conditions were characterized using optical, scanning electron and transmission electron microscopy techniques. The results show that the spray-formed Cu–Cr alloys exhibit a better response to heat treatment compared to the conventional casting alloys. The chromium content has a great effect on the aging behaviors of the Cu–Cr alloys. Although a higher chromium content leads to a larger volume fraction of chromium precipitates generated during the solidification, it does not cause an increase in the mechanical properties. The chromium content should be lower than 1.2 wt.%. © 2007 Elsevier B.V. All rights reserved. Keywords: Cu–Cr alloys; Spray forming; Strength; Microstructure; Thermo-mechanical processing
1. Introduction Due to having the combination of high strength, good electrical and thermal conductivities, Cu–Cr alloys have been applied in many engineering fields, such as in electric resistance welding electrodes, the liner tube of continuous casting crystallizer, the integrated circuit lead frame, the aerial conductor of electric locomotive and so on [1–4]. Recently, the higher strength is more and more required for copper alloys with the rapid development of the electronic industry. For example, a tensile strength σ b > 600 MPa, a hardness Hv > 180 and a conductivity > 80%IACS are demanded for lead frame materials on a very large scale integrated circuit [5,6]. The Cu–Cr alloys prepared by the conventional casting, however, cannot satisfy these properties indices. Spray forming process is especially useful for the preparation of bulk materials with the characteristics of rapid solidification. Recently, many investigations have been reported on the high strength Cu–Cr alloys prepared by spray forming [7–9]
∗
Corresponding author. Tel.: +86 24 23971905; fax: +86 24 23971918. Corresponding author. E-mail addresses: [email protected] (X. Wang), [email protected] (J. Zhao). ∗∗
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.01.023
and the spray-formed Cu–Cr alloys showed excellent mechanical properties. But the work mainly focused on the Cu–Cr alloys which have the chromium content lower than 1 wt.% [8]. Since the chromium content has a great effect on the microstructural features and the mechanical properties of the spray formed Cu–Cr alloys, in the present investigation, Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys were successfully prepared by spray forming and the effect of chromium content on the microstructure and the mechanical properties of the Cu–Cr alloys was discussed.
2. Experimental details A schematic of the spray atomization and deposition process is shown in Fig. 1. The nominal compositions of the alloys were Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr. The Cu–Cr alloys were melted in the vacuum induction furnace. The melt was homogenized at 1230–1250 ◦ C for 1 h prior to the spray deposition. The detail of the set-up has been described elsewhere [10]. The delivery tube size was 4.5 mm. Nitrogen, at an atomizing pressure of 1.5 MPa, was used as the atomizing gas. The spray jet was deposited onto a substrate at a distance of 450 mm from the nozzle. The rotation and axial movement of the substrate guaranteed a constant deposition distance of 450 mm throughout the depo-
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Fig. 1. Schematic of spray atomization and deposition process.
sition time. The sizes of the spray formed billets were 120 mm in height and 150 mm in diameter. The billets were machined and pieces of 100 mm in diameter were removed from the central region of the deposits for the further investigation. After warm rolled at 300 ◦ C with a 40% cross-reduction, the Cu–Cr alloys were subsequently subjected to different types of thermo-mechanical treatments: (1) aging of the Cu–Cr alloys after cold rolled with 60% reduction (SF + WR + CR + aging); (2) aging of the Cu–Cr alloys after solution treatments and cold rolling with 80% reduction (SF + WR + ST + CR + aging). The solution treatment of the alloy in schedule (2) was carried out at 970 ◦ C for 30 min in a nitrogen atmosphere followed by water quenching. All the aging treatments were carried out at temperatures in the range of 350–500 ◦ C in a nitrogen atmosphere. The microstructure of the alloys in various states was characterized by means of optical microscopy (Leica MEF-4), scanning electron microscopy (SS-550) and transmission electron microscopy (Philips-CM12). The tensile properties were determined using an Instron-testing machine AG-5000A. The cross-head speed is 1 mm/min. 3. Results 3.1. Microstructural features An equiaxed grain morphology is distinctly observed in the spray formed alloys, as shown in Fig. 2a for Cu–1.2%Cr and Fig. 2b for Cu–3.2%Cr alloys. This feature was invariably present in the specimens taken from the various regions of the as-deposited Cu–Cr alloys. A bimodal size distribution of the primary chromium particles generated during solidification was observed in the as-sprayed Cu–3.2Cr alloy, as shown in Fig. 2b. The large particles about 1–3 m in diameter dispersed mainly on the grain boundaries coexisting with smaller chromium par-
Fig. 2. Optical micrographs of as-deposited Cu–Cr alloys: (a) Cu–1.2 wt.%Cr and (b) Cu–3.2 wt.%Cr.
ticles of ∼0.5 m within the grains. Occasionally, some big Cr particles can be found on the grain boundaries. These features were not clearly observed in the as-deposited Cu–1.2 wt.%Cr. Fig. 3 shows the TEM image of as-deposited Cu–1.2 wt.%Cr alloy. It can be found that there are some small precipitates of 10 nm in diameter dispersed in the matrix. It can be assumed that aging process took place due to the slow cooling rate of the as-deposited Cu–Cr alloy at the cooling stage. The average grain size of Cu–1.2 wt.%Cr alloy was measured, using the linear intercept method, to be about 50 ± 3.2 m. While the grain size of as-deposited Cu–3.2 wt.%Cr alloy was ∼30 ± 3.0 m, which is much smaller than that of Cu–1.2 wt.%Cr alloy. The difference of the grain size was attributed to the pinning effect of the primary chromium particles that hinder the grain growth. Cold rolling of the spray-formed alloy generated a large number of deformations. The equiaxed grains were squeezed to lamellar and the pores in the deposit were considerably reduced, see Fig. 4. For Cu–3.2 wt.%Cr alloy (Fig. 4b), undissolved Cr (at most 1 wt.% Cr can be dissolved in the solid state) remained as micron-sized particles, which were elongated as Cr fibers during heavy cold working. It has been reported that such fibers can lead to significant strengthening when present in a large volume faction [11,12]. But in our investigation, such fibers are not
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Fig. 3. TEM image showing the microstructure of as-deposited Cu–1.2 wt.%Cr alloy.
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expected to play a major role here where at most about 2% Cr fibers are present. Fig. 5 shows the microstructure of Cu–1.2 wt.%Cr alloy after the SF + WR + ST + CR treatment. A large number of the dislocations within the grains were observed which arises from the high degree of cold deformation. These dislocations can act as diffusion paths for solute atoms and provide nucleation site for precipitation during aging treatment. Deformation twins were also found in the matrix. The twins seemed to form the bundles of a few micrometers thick and several hundred nanometers long. Within the bundles, twin boundaries were not straight along its length. Most twins’ boundaries curved, which is the typical characteristics of the deformation twins [13]. This phenomenon may originate from the fact that a high density of dislocation exists at the boundaries of the twins, as reported in the severely deformed copper alloys [14,15]. However, there was no evidence of this phenomenon in the cold-rolled alloys with 40% reduction. The age hardening behavior is greatly influenced by the initial microstructural features of the alloy that are developed during various processing stages. The results of the present investigation clearly illustrate that the alloy processed by spray forming exhibits a better response to the heat treatment compared to the conventional casting alloys. Fig. 6a shows the TEM micrograph of the Cu–1.2 wt.%Cr alloy after aging treatment at 450 ◦ C for 1 h. The homogeneously distributed globular precipitates of about 5 nm in diameter can be found. Many investigations showed that for rapidly solidified Cu–Cr alloy the precipitates, corresponding to the peak strength or hardness upon aging, are probably coherent ones, which were proved by the fact that the
Fig. 4. Optical microstructure of as-rolled spray-formed Cu–Cr alloys: (a) Cu–1.2 wt.%Cr; (b) Cu–3.2 wt.%Cr. The cold rolling reduction is 80%.
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Fig. 5. TEM morphology of the dislocations and twins for the spray-formed Cu–1.2 wt.%Cr alloy with SF + WR + ST + CR treatment. The cold rolling reduction is 80%.
two-lobe contrast was clearly observed in the following stages of the aging process. However, in the present work, no obvious twolobe contrast caused by coherent distortion was observed in the microstructure. This phenomenon is probably due to the smaller size of the precipitates in the present work. Usually, the two-lobe contrast can only be observed when the size of the precipitate is larger than 30 nm. For example, Liu et al. [5] have reported that the appearance of the two-lobe contrast corresponding to the precipitates with the size about 40–50 nm in a rapidly solidified Cu–Cr alloy. Furthermore, Mclntype et al. [16] found that the two-lobe type strain field and its shape mainly depended on r3 (r is the radius of the precipitate). With the increase of the aging temperature, the volume fraction of precipitates increase slightly, but the sizes of the precipitates did not increase significantly. The selected area electron diffraction (SAED) analysis of the microstructure corresponding to the peak hardness did not show any sign of the bcc Cr phase. It is obvious that the coherent precipitates are not the bcc Cr phase. According to the previous investigation [17], the coherent precipitate should be the fcc Cr phase. Upon further aging, the fcc coherent precipitate lost its coherency with the matrix and transformed to a bcc structure. Fig. 6b shows the TEM microscopy of the Cu–3.2 wt.%Cr alloy after SF + WR + ST + CR + aging at 300 ◦ C for 1 h. This result indicates that the precipitates of an average size of 50 nm uniformly distributed in the matrix. Compared with that observed in Cu–1.2 wt.%Cr at its peak aging (Fig. 6a), the precipitates in this alloy are much bigger. These precipitates seem to lose the coherency with the matrix. With the aging temperature increased or the aging time prolonged, the particles grow quickly and get coarser.
Fig. 6. TEM morphology of precipitates: (a) Cu–1.2 wt.%Cr (SF + ST + CR + aging at 450 ◦ C for 1 h); (b) Cu–3.2 wt.%Cr (SF + ST + CR + aging at 300 ◦ C for 1 h).
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Fig. 7. TEM image of microstructure of Cu–Cr alloys showing small grains. (a) Warm rolled, solution treated, cold-rolled Cu–1.2 wt.%Cr alloy aged at 450 ◦ C for 1 h and (b) warm rolled, solution treated, cold-rolled Cu–3.2 wt.%Cr alloy aged at 450 ◦ C for 1 h.
It has been observed that the Cu–1.2 wt.%Cr alloy at its peak aging exhibited a larger number of the small grains of 100 nm in size, as shown in Fig. 7a. This indicates that the recrystallization occurs during the aging and coarsening process of these small grains is very slow. However, this phenomenon is more evident for the Cu–3.2 wt.%Cr alloy, as shown in Fig. 7b. Even though this alloy (after warm rolled, solution treated and cold rolled with 80% reduction) was aged at 300 ◦ C for 1 h, the recrystallization took place quickly in this alloy and with the aging temperature increased or the aging time prolonged, abnormal grain growth can be found (see Fig. 7b). 3.2. Mechanical properties The variations of the tensile strengths of the Cu–Cr alloys with the aging temperature and the aging time are given in Fig. 8. It can be found that the strength of the spray-formed Cu–Cr alloy is strongly dependent on the thermo-mechanical treatment. For the Cu–3.2 wt.%Cr alloy, the strength increment resulting from heat treatment was usually small. The highest strength is about 530 MPa, which can be obtained after SF + WR + CR + aging at 300 ◦ C for 1 h. It is concluded that SF + WR + ST + CR aging treatment cannot result in the strength improvement for the spray-formed Cu–3.2 wt.% alloy. In contrast, the strength of the spray-formed Cu–1.2 wt.%Cr alloy is strongly dependent on the thermo-mechanical treatment. With the increase of aging temperature, the tensile strength of
the spray-formed Cu–1.2 wt.%Cr alloy increases quickly from 460 MPa at 350 ◦ C to 585 MPa at 450 ◦ C and then decreases. The maximum tensile strength, 585 MPa, was obtained for the samples of Cu–1.2 wt.%Cr alloy after SF + WR + ST + CR + aging at 450 ◦ C for 1 h. The tensile strength of this alloy after peak aging treatment is much higher than that of the Cu–3.2 wt.%Cr alloy. The tensile strengths of the Cu–Cr alloys were measured with the variations of the aging time at the peak aging temperature. The Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloys indicate the peak strengths of 585 and 530 MPa after aged at 450 and 300 ◦ C for 1 h, respectively. Subsequently, the strength value decreases slowly with the aging time prolonged and reaches to the values of 520 and 460 MPa, respectively, when the aging time is 4.5 h. Fig. 9 shows the SEM morphology of the fracture surface. Many equiaxed dimples appeared at the fractured surface. The sizes of the dimple for Cu–3.2 wt.%Cr alloy is bigger than those for Cu–1.2 wt.%Cr alloy. The difference of the fracture surface may be attributed to the effect of the primary chromium particles. The process of crack is readily suggested to be the plastic deformation of the matrix initiating nearby the large particles and then extending when the plastic strain increases. Cracking takes place after the chromium particles are torn out of the matrix under a critical strain. Therefore, the strength of the Cu–Cr alloy can be further improved if the amount and size of the primary chromium particles are reduced with an optimized processing parameters.
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Fig. 8. Variation of the tensile strength of Cu–Cr alloys with aging temperature (a) and aging time (b) at different processing conditions. Superscript 1 represents SF + WR + CR + aging treatment and superscript 2 represents SF + WR + ST + CR + aging treatment. Fig. 9. SEM morphology of the fracture surface: (a) Cu–1.2 wt.%Cr and (b) Cu–3.2 wt.%Cr.
4. Discussions The spray forming process involves the atomization of an alloy melt into a jet of droplets and their subsequent deposition onto a substrate. The droplets experience a high cooling rate of 103 –106 K s−1 [7–9], which therefore renders a rapid solidification features in the spray formed microstructure. The spray formed alloy gives rise to a refined and equiaxed grain morphology in contrast to the dendritic coarse structure, with severe microsegregation, observed in the conventional casting alloy. In addition, the high cooling rate drives the alloy away from the equilibrium condition, and as a result the limit of solid solubility of the alloying elements increases. This induces a better response of the spray formed alloy during aging treatments. However, despite a high cooling rate during atomization, it is difficult to suppress the second phase precipitation as the cooling rate becomes slow after deposition. For the Cu–1.2 wt.%Cr alloy, most of the chromium element remains in solid solution. Only a small part of chromium precipitates on the grain boundaries. On the contrary, for the Cu–3.2 wt.%Cr alloy, more than 2 wt.% chromium precipitate from the matrix during the solidification. A considerable amount of the primary particles has a great effect on the precipitation behavior. Deformation during cold rolling is hindered by these precipitates and therefore
leads to the formation of more geometrically necessary dislocations which aid in a faster rate of precipitation. During aging treatment, Cr solute can easily precipitate from the matrix. Due to the fast diffusion rate, the precipitates grow quickly, coarsen and lose coherent relationship with the matrix, and meanwhile the primary precipitates also coarsen rapidly. All these factors induce the Cu–3.2 wt.%Cr alloy not to have a distinct aging peak. In contrast, for Cu–1.2 wt.%Cr alloy, when most of the chromium remained in solid solution, the absence of precipitates within the grains allows relatively unconstrained flow. During aging, Cr solute can homogeneously precipitate from the matrix and the precipitates keep coherent relationship with the matrix, which results in the enhancement of the strength. In the peak hardness or slight over-aging conditions, the precipitates remain coherent with the matrix. The shear stress increase due to the coherency strengthening σ CS can be calculated by the following equation [18]: rfb 1/2 σCS = MχG(ε)3/2 (1) Γ where ε is the function of the degree of misfit δ, G and b the shear modulus and Burgers vector of the matrix, respectively, r and f the average radius and the volume fraction of the precipitates,
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respectively, Γ the line tension of the dislocations pinned by the precipitates, and χ is a coefficient varying between 2 and 3. Usually, χ is taken as 2.6 [19]. Γ = aGb2 , a varies between 0.089 and 0.5. M is the Taylor factor, with M = 3.06 [20–22]. For the Cu–3.2 wt.%Cr alloys, most of the precipitates are incoherent. The strengthening mechanism usually proposed for the strengthening beyond the peak strength (over-aged alloys having larger and incoherent precipitates) is the Orowan-Ashby model and σ OA can be calculated by the following equation [23]: πr 0.84MGb 1 ln σOA = 2 (2) × 2r0 π (1 − v)1/2 r[8/(3πf )1/2 − 1] where σ OA is the yield strength increase due to the OrowanAshby model and r0 is the inner cut-off radius of the dislocation. The curves corresponding to the variation of σ CS and σ OA with the precipitate volume fraction are presented in Fig. 10. These curves were obtained by substituting the relevant factors and constants (G = 42.1 GPa, v = 0.303, r0 = 3b, b = 0.25 nm for the copper matrix) in Eqs. (1) and (2). Obviously, the strength due to the coherent precipitates is much higher than the strength calculated through the Orowan-Ashby model. With the increase of the coherency precipitates size, the strength increases quickly. On the contrary, the strength calculated through the Orowan-Ashby model decreases with the precipitates increased. Even though the volume fraction of the precipitates in the Cu–1.2 wt.%Cr alloy is about 1.3%, which is much lower than that in the Cu–3.2 wt.%Cr alloy (∼3.3%), the strength of the Cu–1.2 wt.%Cr alloy is much higher than the Cu–3.2 wt.%Cr alloy due to the coherency effect. It appears that the solid solubility of chromium in copper reaches a maximum in the Cu–1.2 wt.%Cr alloy. As the chromium content was lower than 1.2 wt.%, the strength was enhanced with the increase of the Cr content. Beyond this level, although a higher content of chromium lead to a larger volume fraction of chromium precipitates, it does not cause an increase in the mechanical properties. Therefore the selection of the chromium content should be lower than 1.2 wt.%.
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The recrystallization process of the alloys depends on the particle distribution and sizes. At the initial stage of recrystallization, the crystal nuclei, almost free of strain, might grow towards the deformed matrix through migration of large angle boundaries to decrease the strain energy. The driving force for recrystallization FN can be expressed as [24–26]: FN = aGb2 (ρ0 − ρl )
(3)
where a is a constant, G the shear elastic modulus, b Burgers vector, ρ0 and ρl are the dislocation density after deformation and recrystallization, respectively. The above-mentioned results showed that more coarse particles result in higher total strain during cold deformation and consequently the dislocation density in Cu–3.2 wt.%Cr alloys is much higher than that in Cu–1.2 wt.%Cr alloy. The driving force for recrystallization in Cu–3.2 wt.%Cr alloy is about 13.16 MPa and for Cu–1.2 wt.%Cr alloy, this value is about 1.32 MPa. This suggests that the primary particles are responsible for concentrating the strain leading to an increase in driving force for recrystallization. This acceleration is more pronounced during the early stages of aging when the formation of new grains is taking place. For Cu–3.2 wt.%Cr alloy, the recrystallization occurred at 300 ◦ C, which is much lower than that of Cu–1.2 wt.%Cr alloy. The pinning force that the particles produce per unit area of boundary FV can be calculated from the equation [7]: 3fv σgb (4) D where fv is the volume fraction of particles, σ gb is the interface energy per unit area (0.5 J cm−2 for Cu) and D is the average particle diameter. The values obtained when considering the finer particles only are 4.05 and 0.1 MPa for Cu–1.2 wt.%Cr and Cu–3.2 wt.%Cr alloy, respectively. Therefore, the finer particles in Cu–1.2 wt.%Cr produce a pinning force on the boundaries which is nearly forty times larger than that in the Cu–3.2 wt.%Cr alloy. Therefore, the finer particles play an important role in controlling the grain size during recrystallization.
FV =
5. Conclusions The effects of the chromium content and the thermomechanical treatment on the microstructure and mechanical properties of spray-formed Cu–Cr alloys were investigated. The conclusions are as follows:
Fig. 10. Calculated strengthening vs. precipitate volume fraction due to coherency strengthening σ CS for Cu–1.2 wt.%Cr alloy and due to the OrowanAshby model σ OA for Cu–3.2 wt.%Cr alloy.
(1) In the as-deposited state, Cu–Cr alloys exhibit a fine equiaxed grain structure with no macrosegregation. Grain size decreases and the amount of primary precipitates increases as the chromium content increases. Some finer Cr particles precipitated from the matrix due to the slow cooling rate of the billet during the cooling stage. Finer particles play an important role in controlling the grain size during recrystallization. (2) The spray-formed Cu–Cr alloys show excellent deformability. After rolling, aging treatment can be adopted to enhance the mechanical properties of the Cu–Cr alloys. The spray-formed Cu–Cr alloys exhibit a better
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response to heat treatment compared to the conventional casting alloys. The aging process of the Cu–Cr alloys becomes easier with the increase of the chromium content due to the effect of the primary Cr particles. Cu–1.2 wt.%Cr alloy exhibits excellent mechanical properties after proper warm rolling + solution treatment + cold rolling + aging treatment. The thermo-mechanical treatment of warm rolling + solution treatment + cold rolling + aging was not suitable for the spray-formed Cu–3.2 wt.%Cr alloy. (3) Coherency strengthening is the main strengthening mechanism. Attractive combination of strength and ductility can be achieved by the thermo-mechanical processing of the sprayformed Cu–1.2 wt.%Cr alloy. Excess chromium would not lead to the increment of the mechanical properties. Therefore, the selection of the chromium content should be lower than 1.2 wt.%. Acknowledgements The authors gratefully acknowledge the financially support from the National Natural Science Foundation of China (50620130095 and 50671111) and the Natural Science Foundation of Liaoning Province of China (20050047). References
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