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Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling
Journal Pre-proof Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling Xiang Wu, Richu Wang, Chaoqun Peng, Xiaofeng Wang PII:
S0925-8388(20)30734-9
DOI:
https://doi.org/10.1016/j.jallcom.2020.154371
Reference:
JALCOM 154371
To appear in:
Journal of Alloys and Compounds
Received Date: 13 October 2019 Revised Date:
7 February 2020
Accepted Date: 13 February 2020
Please cite this article as: X. Wu, R. Wang, C. Peng, X. Wang, Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154371. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
Xiang Wu: Validation, Formal analysis, Investigation, Resources, Writing - Original Draft, Writing - Review & Editing. Richu Wang: Validation, Formal analysis, Resources, Supervision, Project administration, Funding acquisition. Chaoqun Peng: Methodology, Conceptualization, Writing - Review & Editing, Visualization, Supervision. Xiaofeng Wang: Conceptualization, Resources, Writing - Review & Editing, Supervision, Project administration.
Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling Xiang Wu, Richu Wang, Chaoqun Peng, Xiaofeng Wang*
School of Materials Science and Engineering, Central South University, Changsha 410083, China
*Corresponding author. E-mail: [email protected]
Abstract: To obtain a uniform microstructure with fine second-phase particles, the Cu3Ag0.5Zr0.4Cr0.35Nb alloy is fabricated through gas atomization and hot isostatic pressing, followed by cold rolling. The microstructural evolution and tensile properties of the alloy after isothermal annealing at low temperature are investigated. The results show that the low-temperature annealing leads to an obvious increase in elongation but a mild decrease in strength, and the desirable combinations of strength and elongation after annealing for different times are achieved. When annealed at 350 °C for 1 h, the alloy exhibits a high strength of 665 MPa and an acceptable elongation of 10%. Prolonging the duration to 8 h, a decreased strength (610 MPa) and high elongation (15%) are obtained. The stable nano-sized Cr2Nb and micro-sized Cu4AgZr particles provide a strong pinning effect on boundary migration and significantly slow the subgrain growth during recovery, which contributes to the retention of good strength. The improved ductility results from the recovered ability to accumulate dislocations after annealing. Key words: Rapid solidification; Cold rolling; Low-temperature annealing; Microstructure evolution; Tensile properties
1. Introduction Alloys used for combustion chamber liner of liquid rocket engine are subjected to very high gas temperatures on the hot side and cryogenic hydrogen flow on the back side, and these alloys need to have a combination of good thermal conductivity, high temperature strength, ductility, low cycle fatigue life, and creep resistance [1]. The Cu3Ag0.5Zr alloy (NARloy-Z) is one of the ideal candidate materials to meet the requirements above [2-4]. The strength of this alloy is enhanced by the Ag precipitates and the stable Cu4AgZr particles [3, 5-7]. As the increased requirement on the rocket engine performance, it is crucial to improve the strength of the Cu3Ag0.5Zr alloy while retaining a moderate ductility. The introduction of a little amount of nano-sized Cr2Nb particles is one of the most efficient approaches to resolve the above problem. On the one hand, Cr2Nb is quite stable that will not coarsen even exposed at 800 °C for 50 h, and commonly used to prepare the dispersion hardened Cu-8 at.% Cr-4 at.% Nb alloy (GRCop-84) with high temperature strength [8]. On the other hand, the addition of a little amount of nano-sized Cr2Nb particles has a relatively small impact on the ductility of the Cu3Ag0.5Zr alloy, which has been confirmed in our previous work [9].Therefore, it is a feasible way to enhance the strength without dramatic loss of the ductility by adding 0.4 wt.% Cr and 0.35% wt.% Nb (atomic ratio of about 2:1) into the Cu3Ag0.5Zr alloy. To obtain a Cu3Ag0.5Zr0.4Cr0.35Nb alloy with fine and homogeneous particle dispersion, rapid solidification method was used to prepare the initial material, followed by cold rolling. On the one hand, microsegregation phenomenon is virtually inevitable during the conventional solidification process due to the negligible equilibrium solubility of alloying elements (Zr, Cr and Nb) in the Cu matrix [1, 10]. The rapid solidification process helps to secure a high solubility of 2
insoluble elements and obtain a fine and uniform microstructure. On the other hand, the subsequent cold rolling can further refine the microstructure and produce a high density of dislocations to enhance the alloy significantly. However, the inhomogeneous deformation during cold rolling will introduce a high residual stress with inhomogeneous distribution, and a sharp decrease of elongation inevitably occurs simultaneously for most alloys, which limits the practical utility [11-14]. One compromise method is to release residual stress and enhance the ductility without damaging the strength dramatically by giving a low-temperature annealing after cold deformation [12, 15]. Besides, the low-temperature annealing contributes to decrease the high stored energy in the severe deformation alloys. The high energy can cause easy recovery, recrystallization and grain growth at elevated temperature and results in a loss of strength which currently restricts the use of severe deformation alloys in high temperature applications [16]. Therefore, there is necessary to study the low-temperature annealing behavior of the Cu3Ag0.5Zr0.4Cr0.35Nb alloy and explore various combinations of strength and ductility under different conditions. In the present study, the Cu3Ag0.5Zr0.4Cr0.35Nb alloy is fabricated through rapid solidification followed by cold rolling. The microstructural evolution and tensile properties of the alloy after isothermal annealing at low temperature are investigated. 2. Experiments The Cu3Ag0.5Zr0.4Cr0.35Nb (wt.%) pre-alloyed powder was produced by vacuum melting-high pressure gas atomization (Ar). The powder was mechanically sieved passing 200 mesh and then filled in a cylindrical steel-capsule with dimensions of Φ120 mm×150 mm and a wall thickness of 1.5 mm. The capsule was sealed by welding and degassed. The details of powder 3
production and vacuum degassing process were the same as the previous work [17]. The pre-alloyed powder was densified by hot isostatic pressing (HIP) at 830 °C and 150 MPa for 2 h, followed by furnace cooling. As initial material, the HIPed alloy was cut into samples with dimensions of 120 mm × 40 mm × 12 mm, and rolled at room temperature on a S815E2 two-roll mill with a roller diameter of 550 mm. The reduction ratio was below 10% for each rolling pass, and the final thickness of the alloy was approximately 2.4 mm. The rolled samples were annealed at 350 °C for 0.5, 1, 2, 4, and 8 h, respectively. The microstructural characterizations were performed using a transmission electron microscopy (TEM, Tecnai G2 20) with energy dispersive spectroscopy (EDS) detector and a field emission scanning electron microscopy (FESEM, Nova NanoSEM230) equipped with EDS detector and electron backscatter diffraction (EBSD) system. The specimens for microstructural observations were sectioned along the rolling direction. Before microstructural observations, disc-like specimens with a diameter of 3 mm were punched from thin foils ground to 60 µm. The discs were twin-jet electro-polished using an electrolyte of CH3OH:HNO3=2:1, at -30 °C and 12 V. Line intercept method by counting the number of the intersections between concentric circles and dislocations was employed to calculate the dislocation density (ρ), and the following equation was used [18]. ρ = 2 ⁄( )
(1)
Where N is the number of the intersections between concentric circles and dislocations, L is the secant circle length, and t is the foil thickness of TEM specimen. And the average value of the dislocation density was based on five TEM images for each specimen. The volume fraction (fP) of the Ag precipitates in the alloy was also based on the TEM 4
images and was determined by the following equation [19]: =
−2 ln(1 − +4
)
Where r is the radius of the Ag precipitates (replaced by l for the lamellar-shaped precipitates), t is the foil thickness of the TEM specimen, which is estimated as 80 nm according to the literature [20]. Ap is the projected area fraction of the precipitates, which is quantified by the Image J software. The tensile tests were carried out using a mechanical testing machine (MTS LANDMARK) at a loading speed of 2 mm/min. Dog-bone shaped samples for tensile tests were cut along the rolling direction with a gauge length of 25 mm and a cross section of 8 mm×2 mm. Three samples for each condition were tested to obtain reliable results. The morphology of fractures after testing was observed by SEM. 3. Results and discussion 3.1 Microstructure of initial alloy A large number of gray particles (~0.5 µm, indicated by arrows) and white particles (~3 µm) tend to segregate at the grain boundaries in the HIPed alloy, as shown in Fig. 1a. The magnified image of the two kinds of particles is displayed in Fig. 1b, and corresponding EDS results indicate that the composition of gray particles is close to m phase (Cu4AgZr), that has a broad ternary homogeneity range from Cu67.7Ag9.2Zr15.2 to Cu75.2Ag15.4Zr17.2 [21], while the white particles are identified as Ag precipitates. Accordingly, these coarse Ag precipitates at the grain boundaries need to be further crushed by plastic deformation. In addition, a large number of nano-sized precipitates distribute homogeneously in the matrix (Fig. 1b). The TEM images and corresponding selected area diffraction pattern in Fig. 2 demonstrate that these dispersed precipitates with a size 5
of 10-20 nm have the same fcc structure and crystalline orientation as the Cu matrix, which is in agreement with the continuous Ag precipitates [6, 17, 22-24]. For the Cu-Ag and Cu-Ag-Zr alloys, the nano-sized Ag precipitates with a continuous distribution have a more significant effect on the strength of the matrix than larger-sized precipitates with a discontinuous distribution [7, 10, 25-27].
Fig. 1. (a, b) SEM images of the HIPed alloy, and EDS spectrums of (c) Cu4AgZr and (d) Ag.
Fig. 2. TEM images of the continuous Ag precipitates in the HIPed alloy. 3.2 Microstructural evolution Fig. 3 shows the microstructural evolution of the samples annealed at 350 °C for different
6
times. In the as-rolled sample, the white Ag precipitates with a coarse size in the initial alloy have been crushed into smaller-sized particles arranging in bands paralleled to the rolling direction (RD). A more homogeneous and finer structure is obtained after cold rolling. When the sample is annealed at 350 °C for 0.5 h, the size and distribution of the particles have not changed considerably (Fig. 3b). After annealing for 2 h, only a small amount of Cu4AgZr particles have grown slightly coarser, as indicated by the arrows in Fig. 3c. The number of these larger-sized Cu4AgZr particles increases with a longer annealing time. When further annealed for 8 h, some Cu4AgZr particles grow up to approximately 0.5 µm, as shown in Fig. 3e. The high-density dislocations in the rolled alloy will provide fast diffusion paths for the atoms, which accelerates the rate of the atom diffusion during annealing and results in the sluggish coarsening of the Cu4AgZr particles [28]. In addition, a small number of discontinuous Ag precipitates with lamellar structure are found in some grains, as indicated by arrows in Fig. 3f. Similar precipitates are also reported by other researchers [10, 25, 27].
7
Fig. 3. SEM images of the samples annealed for (a) 0 h, (b) 0.5 h, (c) 2 h, (d) 4 h, and (e, f) 8 h. Fig. 4 shows the TEM micrographs of the as-rolled sample. A high density of dislocations (16.2 × 1013 m-2) is observed in the alloy, and the grain boundaries are difficult to identify (Fig. 4a). Besides, some dislocation tangles are in the vicinity of Cu4AgZr particles with the size of 0.2-0.3 µm, as depicted in Fig. 4b. Furthermore, particles with the size of 50-100 nm (indicated by arrows) are visible in the matrix besides the continuous Ag precipitates in Fig. 4 (c). The high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) image and corresponding elemental distribution maps indicate these particles (50-100 nm) contain mainly Cr and Nb (Fig. 5). According to the Cr-Nb phase diagram [29], these particles are reasonable to be identified as Cr2Nb phase, which is a strengthening phase in the Cu-Cr-Nb alloys [8]. The second-phase particles (continuous Ag precipitates, nano-sized Cr2Nb particles, and
8
micro-sized Cu4AgZr particles) in the as-rolled alloy play a significant role in trapping and accumulating dislocations.
Fig. 4. Bright field TEM images of the as-rolled sample.
Fig. 5. STEM-HAADF image of the as-rolled sample and individual elemental distribution maps of Cu, Ag, Zr, Cr and Nb.
The TEM micrographs of the annealed samples are shown in Figs. 6-8. The dislocation density and volume fraction of Ag precipitates counted by TEM images are summarized in Table 1. When the sample is annealed at 350 °C for 0.5 h, the typical deformation microstructure with a high density of dislocations (13.9 × 1013 m-2) is still seen in Fig. 6a, and the high-density
9
dislocation region (H region) shows an irregular shape. Besides, the annihilation and rearrangement of dislocations lead to the formation of low-density dislocation region (L region). Furthermore, equiaxed subgrains with blurred boundaries are observed in Fig. 6b, which suggests that recovery occurs in the alloy. The continuous Ag precipitates with a volume fraction of approximately 2.76% still distribute evenly in the matrix (Fig. 6c). The microstructure of the sample annealed for 2 h is presented in Fig. 6d-f. Deformation band (DB) with well-defined boundaries are observed (Fig. 6d). In addition, some recovered subgrains with few dislocations in the interior is found to occur, as indicated in Fig. 6e. The volume fraction of the continuous Ag precipitates decreases slightly compared with that in the alloy annealed for 0.5 h.
Fig. 6. TEM images of the samples annealed for (a-c) 0.5 h and (d-f) 2 h. As shown in Fig. 7a, many deformation twins (DT) are remained in the deformation band after annealing for 4 h, and these clustered twins have a lenticular shape with an irregular interface to the matrix (this character is more distinct in the EBSD map in Fig. 9c), and the thickness is 10
approximately 20 nm, which are the typical features of the deformation twins [30]. Furthermore, the inserted selected area electron diffraction (SAED) pattern demonstrates the existence of the twins. Besides, more recovered subgrains with well-defined boundaries occur (Fig. 7b), and these subgrains have a relatively larger size compared with those in the sample annealed for 2 h. This means that a subgrain growth happens in the alloy. It is generally accepted that the growth of subgrains is achieved by the low angle boundary migration during recovery, and the rate of migration depends on the triple junction mobility [31]. However, the growth of subgrains is restricted by the fine Cu4AgZr particles, as indicated in Fig. 7c. It is substantial evidence that the fine particles have a strong pinning effect on both grain and subgrain boundaries, thereby contributing to the stabilization of the recovered microstructure [32, 33]. Furthermore, the volume fraction of the continuous Ag precipitates in the matrix decreases to 1.94% and these nano-sized particles present the trend of redissolution (Fig. 7d). When the sample is annealed for 8 h, the density of dislocations decreases significantly (7.2 × 1013 m-2), and the discontinuous Ag precipitates with lamellar structure in some grains are observed in Fig. 7e, which is in agreement with the SEM observation in Fig. 3f. One can also see from Fig. 7f that the continuous Ag precipitates present an inhomogeneous distribution between different grains after annealing for a long time, and both the number density and size of the nano-sized precipitates decrease appreciably in comparison with those in the HIPed and as-rolled alloy. The continuous Ag precipitates tend to redissolve into the matrix and the Ag solutes transfer from finer-sized continuous Ag precipitates to larger-sized discontinuous ones during a long-time annealing, which is attributed to the increased driving force for solute diffusion arising from severe rolling. Therefore, the discontinuous precipitation cannot be inhibited completely by the 11
addition of 0.5 wt.% Zr for the severe-rolled alloy when annealed at 350 °C for 8 h. The precipitation mode of Ag depends on the diffusion conditions, which is confirmed by other groups [5, 34, 35]. The high-density dislocations in the rolled sample contributes to the diffusion of Ag solutes during long-time annealing. Nevertheless, the volume fraction of the discontinuous Ag precipitates is only approximately 0.28%, which is far smaller than that of the continuous ones (1.37%) in the alloy annealed for 8 h, as shown in Table 1. This means that the continuous mode is still the dominating precipitation way because the diffusion process is very slow due to the suppression of the minor Zr [3, 5]. In addition, no significant coarsening of the Cr2Nb particles occurs after annealing for 8 h, as shown in Fig. 8. These nano-sized particles with good thermal stability play the similar role in pinning the subgrain boundaries and retarding the subgrain growth during recovery.
Fig. 7. TEM images of the samples annealed for (a-d) 4 h and (e-f) 8 h.
12
Fig. 8. (a) Cr2Nb particles in the sample annealed for 8 h and (b) corresponding EDS spectrum. The orientation, boundary characteristics and size distribution of grains in the annealed samples were analyzed using EBSD with a step size of 100 nm. The inverse pole figure (IPF) maps of the samples annealed at 350 °C for 0.5, 4, and 8 h, respectively, are shown in Fig. 9a, c and e. The corresponding grain boundary images are shown in Fig. 9b, d and f. The boundaries with a misorientation angle range from 2° to 9° are generally defined as low angle boundaries (LABs), and boundaries with a misorientation angle above 9° are defined as high angle boundaries (HABs), depicted by red and blue lines, respectively. As shown in Fig.9a, the microstructure of the sample annealed for 0.5 h is comprised of fine equiaxed grains, large elongated grains, and deformation band along the rolling direction (RD). The high fraction of LABs introduced by the severe cold rolling are distributed in the interior of the elongated grains and deformation band, as displayed in Fig. 9b. Especially, when the sample is annealed for 4 h, the deformation-induced microstructural features, such as deformation band (DB), deformation twins (DT), are retained (Fig. 9c). After annealing for 8 h, a small amount of recrystallized grains (~6 µm) without LABs inside are observed, which demonstrates that a partial recrystallization starts even before the recovery is completed in the long-time annealed sample. The area fraction of the recrystallized grains is approximately 16%, and the statistical result is based on five boundary maps. Besides,
13
the deformation twins in the deformation band are disappeared. Instead, some annealing twins (AT) occur in the recrystallized grains, as shown in Fig. 9e and f. These annealing twins with the thickness of approximately 0.4 µm penetrate into the whole grains. Several theoretical models have been proposed to explain the formation mechanism of the annealing twins, but there are still no final conclusions [36-40]. Generally, the LABs are composed of dislocations, and it is believed that the disappearance of LABs (dislocation annihilation) in the recrystallized grains plays a vital role in the formation of annealing twins. The partial recrystallization occurs accompanied by the grain boundary migration, which is driven by the annihilation of dislocations on the one side of the grain boundary. The annealing twins originate from the migrating boundary and extend into the dislocation-free recrystallized grain [40]. The resulting histograms of grain size and boundary misorientation angle from the EBSD maps are displayed in Fig. 10. The average grain size, area fraction of the ultrafine grains (UFG), and number fraction of the HABs are presented in Table 1. Due to the limitation of step size, some nano-sized grains are not counted. It is found that approximately 24% grains have an ultrafine size (~600 nm), and over 73% grain boundaries belong to LABs in the sample annealed for 0.5 h. The percentages of UFG and LABs decrease with the increase of the annealing time. It is calculated that the percentage of LABs has decreased to 51%, and only 13% grains have an ultrafine size after annealing for 8 h. The gradual disappearance of UFG and the progressive transformation of grain boundaries from LABs into HABs can be attributed to the extend recovery and partial recrystallization during long time annealing. The LABs are consumed by new-formed HABs, leading to the subgrain boundary migration and subgrain growth.
14
Fig. 9. Orientation maps and corresponding boundary images of the samples annealed at 350 °C for (a, b) 0.5 h, (c, d) 4 h, (e, f) 8 h. The inserted color scheme in the orientation map presents grain orientation (green: [101]; blue: [111]; red [001]). The red lines in the boundary images present LABs, while blue lines present HABs.
15
Fig. 10. Distributions of (a) grain size and (b) boundary misorientation angle for the samples annealed at 350 °C for different time. Table 1 The microstructural characteristics of the as-rolled and the annealed samples. Samples
Average grain
Area fraction
Dislocation
Number fraction
Volume fraction of Ag precipitates
size
of UFG
density
of HABs
(%)
(µm)
(%)
13
(× 10 m )
(%)
Continuous Ag
Discontinuous Ag
As-rolled
1.7 ± 0.3
28
16.2 ± 1.8
18
2.76 ± 0.12
--
Annealed-0.5 h
1.9 ± 0.2
24
13.9 ± 2.1
27
2.55 ± 0.18
--
Annealed-1 h
2.1 ± 0.4
22
12.1 ± 2.3
31
2.36 ± 0.11
--
Annealed-2 h
2.2 ± 0.3
19
10.4 ± 1.7
34
2.19 ± 0.14
--
Annealed-4 h
2.5 ± 0.5
16
8.9 ± 1.9
40
1.94 ± 0.19
--
Annealed-8 h
3.4 ± 0.8
13
7.2 ± 1.6
51
1.37 ± 0.13
0.28 ± 0.16
-2
3.3 Tensile properties Fig. 11a presents the engineering stress-strain curves of the samples annealed at 350 °C for different durations, and the tensile properties are summarized in Fig. 11b. It is obvious that the low-temperature annealing process results in a recovery of ductility and a mild decrease of strength. The as-rolled sample exhibits high ultimate tensile strength (UTS) of 700 MPa, and yield
16
strength (YS) of 670 MPa, respectively, but a low elongation of 5%. After annealing for 1 h, a pretty high strength of 665 MPa and an acceptable elongation of 10% are obtained. The ductility of the alloy gradually recovers with the increase of the annealing time at the expense of the decrease of the strength. When the sample is annealed for 8 h, the values of UTS, YS and elongation are 610 MPa, 581 MPa, and 15%, respectively. Thus, the desirable combinations of strength and ductility have been achieved after low-temperature annealing for different times. The high density of dislocations and second phase particles (continuous Ag precipitates, nano-sized Cr2Nb particles, and micro-sized Cu4AgZr particles) are extensively observed in the as-rolled sample and are mainly responsible for the improvement of the strength of the alloy. However, the high-density dislocations in the rolled alloy also restrict the strain hardening during tensile tests, which results in a low elongation. After annealing at low temperature for a short time, the dislocation density reduces through absorption during recovery. However, the nano-sized Cr2Nb and micro-sized Cu4AgZr particles provide a strong pinning effect on boundary migration and significantly slow the subgrain growth during recovery. Thus, the high-density grain boundaries exert a significant effect on preventing the dislocation motion during the tensile test, which contributes to the relatively high strength. Simultaneously, the reduction of dislocation density in the annealed sample provides new space for dislocation generation and accumulation during the tensile test, which improves the strain hardening of the alloy before necking compared to the as-rolled sample, resulting in the enhancement of the ductility. With prolonging the annealing time, the dislocation annihilation and grain growth due to the extend recovery and partial recrystallization lead to the decrease of strength. Besides, the slight coarsening of the Cu4AgZr particles and the appearance of the discontinuous Ag precipitates during annealing also 17
have an adverse effect on the strength. It is worth mentioning that twin boundaries formed during annealing perform the similar role as the grain boundaries on hindering the dislocation motion, and exert a good ability to accumulate dislocation and improve the ductility during tensile test [41, 42].
Fig. 11. (a) Engineering stress-strain curves and (b) tensile properties of the samples annealed at 350 °C for different time. Compared with the Cu3Ag0.5Zr alloy prepared by various routes [4, 43-45], the Cu3Ag0.5Zr0.4Cr0.35Nb alloy in the present work exhibits superior tensile properties, as demonstrated in Fig. 12, which is more suitable for practical utility. The introduction of the small amount of Cr2Nb particles and the fine microstructure obtained through rapid solidification followed by cold rolling are responsible for the excellent tensile properties.
18
Fig. 12. Comparison in yield strength and elongation of the Cu3Ag0.5Zr0.4Cr0.35Nb alloy and Cu3Ag0.5Zr alloys prepared by various routes. Fig. 13 shows the fracture surfaces of the samples before and after annealing, and the morphologies of the fracture surfaces are dominated by dimple features. The large amounts of shallow dimples with the size 0.2-1.0 µm are observed in the as-rolled sample (Fig. 13a), demonstrating a high strength but a low elongation of this alloy. After annealing at 350 °C for 0.5 h, some bigger and deeper ductile dimples on the fracture surface can be seen in Fig. 13b, indicating a recovered elongation. The dimples in the samples annealed for over 4 h distribute more uniformly, and the size of the dimples increases evidently, which agrees with the lower strength but higher elongation after a long annealing time.
19
Fig. 13. SEM fractographs of the samples annealed at 350 °C for (a) 0 h, (b) 0.5 h, (c) 4 h and (d) 8 h. 4 Conclusions (1) The Cu3Ag0.5Zr0.4Cr0.35Nb alloy fabricated through gas atomization and hot isostatic pressing followed by cold rolling shows a uniform microstructure with continuous Ag precipitates, nano-sized Cr2Nb particles, and micro-sized Cu4AgZr particles, and exhibits a high strength of 700 MPa but a low elongation of 5%. (2) The desirable combinations of strength and ductility have been achieved after low-temperature annealing. When annealed at 350 °C for 1 h, the alloy exhibits a high strength of 665 MPa and an acceptable elongation of 10%. The ductility of the alloy gradually recovers with the increase of annealing time at the expense of the decreased strength. Prolonging the duration to 8 h, the strength decreases to 610 MPa, while a high elongation of 15% is obtained.
20
(3) The retention of a relatively high strength after annealing at a short time is mainly attributed to the pinning effect on grain boundary migration of the nano-sized Cr2Nb and micro-sized Cu4AgZr particles, which significantly slows the subgrain growth during recovery. With prolonging the annealing time, the dislocation annihilation and subgrain growth during the extend recovery and partial recrystallization lead to the decrease of strength. Besides, the slight coarsening of the Cu4AgZr particles and the appearance of the discontinuous Ag precipitates after annealing for 8 h also have an adverse effect on the strength. The recovered ability to accumulate dislocations of the large number of particles and grain boundaries results in the enhancement of the ductility. Besides, the annealing twins also play an important role in accumulating dislocations and improving the ductility. Acknowledgement This work was supported by the Science and Technology Plan Projects of Hunan Province, China [2017GK2261].
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Highlights Cu3Ag0.5Zr0.4Cr0.35Nb alloy is prepared by rapid solidification and cold rolling. Desirable combinations of strength and ductility after annealing are achieved. The stable Cr2Nb and Cu4AgZr particles prevent the grain growth during annealing. Enhanced ductility is attributed to the recovered ability to accumulate dislocation.
No conflict of interest exits in this paper, and the paper is approved by all authors for publication. I would like to declare that the work described in this paper is our original research that has not been published in any form previously, and not under consideration for publication elsewhere, in whole or in part. All the authors listed have approved the paper that is enclosed.
S0925-8388(20)30734-9
DOI:
https://doi.org/10.1016/j.jallcom.2020.154371
Reference:
JALCOM 154371
To appear in:
Journal of Alloys and Compounds
Received Date: 13 October 2019 Revised Date:
7 February 2020
Accepted Date: 13 February 2020
Please cite this article as: X. Wu, R. Wang, C. Peng, X. Wang, Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154371. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
Xiang Wu: Validation, Formal analysis, Investigation, Resources, Writing - Original Draft, Writing - Review & Editing. Richu Wang: Validation, Formal analysis, Resources, Supervision, Project administration, Funding acquisition. Chaoqun Peng: Methodology, Conceptualization, Writing - Review & Editing, Visualization, Supervision. Xiaofeng Wang: Conceptualization, Resources, Writing - Review & Editing, Supervision, Project administration.
Low-temperature annealing behavior and tensile properties of the rapidly solidified Cu3Ag0.5Zr0.4Cr0.35Nb alloy reinforced by cold rolling Xiang Wu, Richu Wang, Chaoqun Peng, Xiaofeng Wang*
School of Materials Science and Engineering, Central South University, Changsha 410083, China
*Corresponding author. E-mail: [email protected]
Abstract: To obtain a uniform microstructure with fine second-phase particles, the Cu3Ag0.5Zr0.4Cr0.35Nb alloy is fabricated through gas atomization and hot isostatic pressing, followed by cold rolling. The microstructural evolution and tensile properties of the alloy after isothermal annealing at low temperature are investigated. The results show that the low-temperature annealing leads to an obvious increase in elongation but a mild decrease in strength, and the desirable combinations of strength and elongation after annealing for different times are achieved. When annealed at 350 °C for 1 h, the alloy exhibits a high strength of 665 MPa and an acceptable elongation of 10%. Prolonging the duration to 8 h, a decreased strength (610 MPa) and high elongation (15%) are obtained. The stable nano-sized Cr2Nb and micro-sized Cu4AgZr particles provide a strong pinning effect on boundary migration and significantly slow the subgrain growth during recovery, which contributes to the retention of good strength. The improved ductility results from the recovered ability to accumulate dislocations after annealing. Key words: Rapid solidification; Cold rolling; Low-temperature annealing; Microstructure evolution; Tensile properties
1. Introduction Alloys used for combustion chamber liner of liquid rocket engine are subjected to very high gas temperatures on the hot side and cryogenic hydrogen flow on the back side, and these alloys need to have a combination of good thermal conductivity, high temperature strength, ductility, low cycle fatigue life, and creep resistance [1]. The Cu3Ag0.5Zr alloy (NARloy-Z) is one of the ideal candidate materials to meet the requirements above [2-4]. The strength of this alloy is enhanced by the Ag precipitates and the stable Cu4AgZr particles [3, 5-7]. As the increased requirement on the rocket engine performance, it is crucial to improve the strength of the Cu3Ag0.5Zr alloy while retaining a moderate ductility. The introduction of a little amount of nano-sized Cr2Nb particles is one of the most efficient approaches to resolve the above problem. On the one hand, Cr2Nb is quite stable that will not coarsen even exposed at 800 °C for 50 h, and commonly used to prepare the dispersion hardened Cu-8 at.% Cr-4 at.% Nb alloy (GRCop-84) with high temperature strength [8]. On the other hand, the addition of a little amount of nano-sized Cr2Nb particles has a relatively small impact on the ductility of the Cu3Ag0.5Zr alloy, which has been confirmed in our previous work [9].Therefore, it is a feasible way to enhance the strength without dramatic loss of the ductility by adding 0.4 wt.% Cr and 0.35% wt.% Nb (atomic ratio of about 2:1) into the Cu3Ag0.5Zr alloy. To obtain a Cu3Ag0.5Zr0.4Cr0.35Nb alloy with fine and homogeneous particle dispersion, rapid solidification method was used to prepare the initial material, followed by cold rolling. On the one hand, microsegregation phenomenon is virtually inevitable during the conventional solidification process due to the negligible equilibrium solubility of alloying elements (Zr, Cr and Nb) in the Cu matrix [1, 10]. The rapid solidification process helps to secure a high solubility of 2
insoluble elements and obtain a fine and uniform microstructure. On the other hand, the subsequent cold rolling can further refine the microstructure and produce a high density of dislocations to enhance the alloy significantly. However, the inhomogeneous deformation during cold rolling will introduce a high residual stress with inhomogeneous distribution, and a sharp decrease of elongation inevitably occurs simultaneously for most alloys, which limits the practical utility [11-14]. One compromise method is to release residual stress and enhance the ductility without damaging the strength dramatically by giving a low-temperature annealing after cold deformation [12, 15]. Besides, the low-temperature annealing contributes to decrease the high stored energy in the severe deformation alloys. The high energy can cause easy recovery, recrystallization and grain growth at elevated temperature and results in a loss of strength which currently restricts the use of severe deformation alloys in high temperature applications [16]. Therefore, there is necessary to study the low-temperature annealing behavior of the Cu3Ag0.5Zr0.4Cr0.35Nb alloy and explore various combinations of strength and ductility under different conditions. In the present study, the Cu3Ag0.5Zr0.4Cr0.35Nb alloy is fabricated through rapid solidification followed by cold rolling. The microstructural evolution and tensile properties of the alloy after isothermal annealing at low temperature are investigated. 2. Experiments The Cu3Ag0.5Zr0.4Cr0.35Nb (wt.%) pre-alloyed powder was produced by vacuum melting-high pressure gas atomization (Ar). The powder was mechanically sieved passing 200 mesh and then filled in a cylindrical steel-capsule with dimensions of Φ120 mm×150 mm and a wall thickness of 1.5 mm. The capsule was sealed by welding and degassed. The details of powder 3
production and vacuum degassing process were the same as the previous work [17]. The pre-alloyed powder was densified by hot isostatic pressing (HIP) at 830 °C and 150 MPa for 2 h, followed by furnace cooling. As initial material, the HIPed alloy was cut into samples with dimensions of 120 mm × 40 mm × 12 mm, and rolled at room temperature on a S815E2 two-roll mill with a roller diameter of 550 mm. The reduction ratio was below 10% for each rolling pass, and the final thickness of the alloy was approximately 2.4 mm. The rolled samples were annealed at 350 °C for 0.5, 1, 2, 4, and 8 h, respectively. The microstructural characterizations were performed using a transmission electron microscopy (TEM, Tecnai G2 20) with energy dispersive spectroscopy (EDS) detector and a field emission scanning electron microscopy (FESEM, Nova NanoSEM230) equipped with EDS detector and electron backscatter diffraction (EBSD) system. The specimens for microstructural observations were sectioned along the rolling direction. Before microstructural observations, disc-like specimens with a diameter of 3 mm were punched from thin foils ground to 60 µm. The discs were twin-jet electro-polished using an electrolyte of CH3OH:HNO3=2:1, at -30 °C and 12 V. Line intercept method by counting the number of the intersections between concentric circles and dislocations was employed to calculate the dislocation density (ρ), and the following equation was used [18]. ρ = 2 ⁄( )
(1)
Where N is the number of the intersections between concentric circles and dislocations, L is the secant circle length, and t is the foil thickness of TEM specimen. And the average value of the dislocation density was based on five TEM images for each specimen. The volume fraction (fP) of the Ag precipitates in the alloy was also based on the TEM 4
images and was determined by the following equation [19]: =
−2 ln(1 − +4
)
Where r is the radius of the Ag precipitates (replaced by l for the lamellar-shaped precipitates), t is the foil thickness of the TEM specimen, which is estimated as 80 nm according to the literature [20]. Ap is the projected area fraction of the precipitates, which is quantified by the Image J software. The tensile tests were carried out using a mechanical testing machine (MTS LANDMARK) at a loading speed of 2 mm/min. Dog-bone shaped samples for tensile tests were cut along the rolling direction with a gauge length of 25 mm and a cross section of 8 mm×2 mm. Three samples for each condition were tested to obtain reliable results. The morphology of fractures after testing was observed by SEM. 3. Results and discussion 3.1 Microstructure of initial alloy A large number of gray particles (~0.5 µm, indicated by arrows) and white particles (~3 µm) tend to segregate at the grain boundaries in the HIPed alloy, as shown in Fig. 1a. The magnified image of the two kinds of particles is displayed in Fig. 1b, and corresponding EDS results indicate that the composition of gray particles is close to m phase (Cu4AgZr), that has a broad ternary homogeneity range from Cu67.7Ag9.2Zr15.2 to Cu75.2Ag15.4Zr17.2 [21], while the white particles are identified as Ag precipitates. Accordingly, these coarse Ag precipitates at the grain boundaries need to be further crushed by plastic deformation. In addition, a large number of nano-sized precipitates distribute homogeneously in the matrix (Fig. 1b). The TEM images and corresponding selected area diffraction pattern in Fig. 2 demonstrate that these dispersed precipitates with a size 5
of 10-20 nm have the same fcc structure and crystalline orientation as the Cu matrix, which is in agreement with the continuous Ag precipitates [6, 17, 22-24]. For the Cu-Ag and Cu-Ag-Zr alloys, the nano-sized Ag precipitates with a continuous distribution have a more significant effect on the strength of the matrix than larger-sized precipitates with a discontinuous distribution [7, 10, 25-27].
Fig. 1. (a, b) SEM images of the HIPed alloy, and EDS spectrums of (c) Cu4AgZr and (d) Ag.
Fig. 2. TEM images of the continuous Ag precipitates in the HIPed alloy. 3.2 Microstructural evolution Fig. 3 shows the microstructural evolution of the samples annealed at 350 °C for different
6
times. In the as-rolled sample, the white Ag precipitates with a coarse size in the initial alloy have been crushed into smaller-sized particles arranging in bands paralleled to the rolling direction (RD). A more homogeneous and finer structure is obtained after cold rolling. When the sample is annealed at 350 °C for 0.5 h, the size and distribution of the particles have not changed considerably (Fig. 3b). After annealing for 2 h, only a small amount of Cu4AgZr particles have grown slightly coarser, as indicated by the arrows in Fig. 3c. The number of these larger-sized Cu4AgZr particles increases with a longer annealing time. When further annealed for 8 h, some Cu4AgZr particles grow up to approximately 0.5 µm, as shown in Fig. 3e. The high-density dislocations in the rolled alloy will provide fast diffusion paths for the atoms, which accelerates the rate of the atom diffusion during annealing and results in the sluggish coarsening of the Cu4AgZr particles [28]. In addition, a small number of discontinuous Ag precipitates with lamellar structure are found in some grains, as indicated by arrows in Fig. 3f. Similar precipitates are also reported by other researchers [10, 25, 27].
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Fig. 3. SEM images of the samples annealed for (a) 0 h, (b) 0.5 h, (c) 2 h, (d) 4 h, and (e, f) 8 h. Fig. 4 shows the TEM micrographs of the as-rolled sample. A high density of dislocations (16.2 × 1013 m-2) is observed in the alloy, and the grain boundaries are difficult to identify (Fig. 4a). Besides, some dislocation tangles are in the vicinity of Cu4AgZr particles with the size of 0.2-0.3 µm, as depicted in Fig. 4b. Furthermore, particles with the size of 50-100 nm (indicated by arrows) are visible in the matrix besides the continuous Ag precipitates in Fig. 4 (c). The high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) image and corresponding elemental distribution maps indicate these particles (50-100 nm) contain mainly Cr and Nb (Fig. 5). According to the Cr-Nb phase diagram [29], these particles are reasonable to be identified as Cr2Nb phase, which is a strengthening phase in the Cu-Cr-Nb alloys [8]. The second-phase particles (continuous Ag precipitates, nano-sized Cr2Nb particles, and
8
micro-sized Cu4AgZr particles) in the as-rolled alloy play a significant role in trapping and accumulating dislocations.
Fig. 4. Bright field TEM images of the as-rolled sample.
Fig. 5. STEM-HAADF image of the as-rolled sample and individual elemental distribution maps of Cu, Ag, Zr, Cr and Nb.
The TEM micrographs of the annealed samples are shown in Figs. 6-8. The dislocation density and volume fraction of Ag precipitates counted by TEM images are summarized in Table 1. When the sample is annealed at 350 °C for 0.5 h, the typical deformation microstructure with a high density of dislocations (13.9 × 1013 m-2) is still seen in Fig. 6a, and the high-density
9
dislocation region (H region) shows an irregular shape. Besides, the annihilation and rearrangement of dislocations lead to the formation of low-density dislocation region (L region). Furthermore, equiaxed subgrains with blurred boundaries are observed in Fig. 6b, which suggests that recovery occurs in the alloy. The continuous Ag precipitates with a volume fraction of approximately 2.76% still distribute evenly in the matrix (Fig. 6c). The microstructure of the sample annealed for 2 h is presented in Fig. 6d-f. Deformation band (DB) with well-defined boundaries are observed (Fig. 6d). In addition, some recovered subgrains with few dislocations in the interior is found to occur, as indicated in Fig. 6e. The volume fraction of the continuous Ag precipitates decreases slightly compared with that in the alloy annealed for 0.5 h.
Fig. 6. TEM images of the samples annealed for (a-c) 0.5 h and (d-f) 2 h. As shown in Fig. 7a, many deformation twins (DT) are remained in the deformation band after annealing for 4 h, and these clustered twins have a lenticular shape with an irregular interface to the matrix (this character is more distinct in the EBSD map in Fig. 9c), and the thickness is 10
approximately 20 nm, which are the typical features of the deformation twins [30]. Furthermore, the inserted selected area electron diffraction (SAED) pattern demonstrates the existence of the twins. Besides, more recovered subgrains with well-defined boundaries occur (Fig. 7b), and these subgrains have a relatively larger size compared with those in the sample annealed for 2 h. This means that a subgrain growth happens in the alloy. It is generally accepted that the growth of subgrains is achieved by the low angle boundary migration during recovery, and the rate of migration depends on the triple junction mobility [31]. However, the growth of subgrains is restricted by the fine Cu4AgZr particles, as indicated in Fig. 7c. It is substantial evidence that the fine particles have a strong pinning effect on both grain and subgrain boundaries, thereby contributing to the stabilization of the recovered microstructure [32, 33]. Furthermore, the volume fraction of the continuous Ag precipitates in the matrix decreases to 1.94% and these nano-sized particles present the trend of redissolution (Fig. 7d). When the sample is annealed for 8 h, the density of dislocations decreases significantly (7.2 × 1013 m-2), and the discontinuous Ag precipitates with lamellar structure in some grains are observed in Fig. 7e, which is in agreement with the SEM observation in Fig. 3f. One can also see from Fig. 7f that the continuous Ag precipitates present an inhomogeneous distribution between different grains after annealing for a long time, and both the number density and size of the nano-sized precipitates decrease appreciably in comparison with those in the HIPed and as-rolled alloy. The continuous Ag precipitates tend to redissolve into the matrix and the Ag solutes transfer from finer-sized continuous Ag precipitates to larger-sized discontinuous ones during a long-time annealing, which is attributed to the increased driving force for solute diffusion arising from severe rolling. Therefore, the discontinuous precipitation cannot be inhibited completely by the 11
addition of 0.5 wt.% Zr for the severe-rolled alloy when annealed at 350 °C for 8 h. The precipitation mode of Ag depends on the diffusion conditions, which is confirmed by other groups [5, 34, 35]. The high-density dislocations in the rolled sample contributes to the diffusion of Ag solutes during long-time annealing. Nevertheless, the volume fraction of the discontinuous Ag precipitates is only approximately 0.28%, which is far smaller than that of the continuous ones (1.37%) in the alloy annealed for 8 h, as shown in Table 1. This means that the continuous mode is still the dominating precipitation way because the diffusion process is very slow due to the suppression of the minor Zr [3, 5]. In addition, no significant coarsening of the Cr2Nb particles occurs after annealing for 8 h, as shown in Fig. 8. These nano-sized particles with good thermal stability play the similar role in pinning the subgrain boundaries and retarding the subgrain growth during recovery.
Fig. 7. TEM images of the samples annealed for (a-d) 4 h and (e-f) 8 h.
12
Fig. 8. (a) Cr2Nb particles in the sample annealed for 8 h and (b) corresponding EDS spectrum. The orientation, boundary characteristics and size distribution of grains in the annealed samples were analyzed using EBSD with a step size of 100 nm. The inverse pole figure (IPF) maps of the samples annealed at 350 °C for 0.5, 4, and 8 h, respectively, are shown in Fig. 9a, c and e. The corresponding grain boundary images are shown in Fig. 9b, d and f. The boundaries with a misorientation angle range from 2° to 9° are generally defined as low angle boundaries (LABs), and boundaries with a misorientation angle above 9° are defined as high angle boundaries (HABs), depicted by red and blue lines, respectively. As shown in Fig.9a, the microstructure of the sample annealed for 0.5 h is comprised of fine equiaxed grains, large elongated grains, and deformation band along the rolling direction (RD). The high fraction of LABs introduced by the severe cold rolling are distributed in the interior of the elongated grains and deformation band, as displayed in Fig. 9b. Especially, when the sample is annealed for 4 h, the deformation-induced microstructural features, such as deformation band (DB), deformation twins (DT), are retained (Fig. 9c). After annealing for 8 h, a small amount of recrystallized grains (~6 µm) without LABs inside are observed, which demonstrates that a partial recrystallization starts even before the recovery is completed in the long-time annealed sample. The area fraction of the recrystallized grains is approximately 16%, and the statistical result is based on five boundary maps. Besides,
13
the deformation twins in the deformation band are disappeared. Instead, some annealing twins (AT) occur in the recrystallized grains, as shown in Fig. 9e and f. These annealing twins with the thickness of approximately 0.4 µm penetrate into the whole grains. Several theoretical models have been proposed to explain the formation mechanism of the annealing twins, but there are still no final conclusions [36-40]. Generally, the LABs are composed of dislocations, and it is believed that the disappearance of LABs (dislocation annihilation) in the recrystallized grains plays a vital role in the formation of annealing twins. The partial recrystallization occurs accompanied by the grain boundary migration, which is driven by the annihilation of dislocations on the one side of the grain boundary. The annealing twins originate from the migrating boundary and extend into the dislocation-free recrystallized grain [40]. The resulting histograms of grain size and boundary misorientation angle from the EBSD maps are displayed in Fig. 10. The average grain size, area fraction of the ultrafine grains (UFG), and number fraction of the HABs are presented in Table 1. Due to the limitation of step size, some nano-sized grains are not counted. It is found that approximately 24% grains have an ultrafine size (~600 nm), and over 73% grain boundaries belong to LABs in the sample annealed for 0.5 h. The percentages of UFG and LABs decrease with the increase of the annealing time. It is calculated that the percentage of LABs has decreased to 51%, and only 13% grains have an ultrafine size after annealing for 8 h. The gradual disappearance of UFG and the progressive transformation of grain boundaries from LABs into HABs can be attributed to the extend recovery and partial recrystallization during long time annealing. The LABs are consumed by new-formed HABs, leading to the subgrain boundary migration and subgrain growth.
14
Fig. 9. Orientation maps and corresponding boundary images of the samples annealed at 350 °C for (a, b) 0.5 h, (c, d) 4 h, (e, f) 8 h. The inserted color scheme in the orientation map presents grain orientation (green: [101]; blue: [111]; red [001]). The red lines in the boundary images present LABs, while blue lines present HABs.
15
Fig. 10. Distributions of (a) grain size and (b) boundary misorientation angle for the samples annealed at 350 °C for different time. Table 1 The microstructural characteristics of the as-rolled and the annealed samples. Samples
Average grain
Area fraction
Dislocation
Number fraction
Volume fraction of Ag precipitates
size
of UFG
density
of HABs
(%)
(µm)
(%)
13
(× 10 m )
(%)
Continuous Ag
Discontinuous Ag
As-rolled
1.7 ± 0.3
28
16.2 ± 1.8
18
2.76 ± 0.12
--
Annealed-0.5 h
1.9 ± 0.2
24
13.9 ± 2.1
27
2.55 ± 0.18
--
Annealed-1 h
2.1 ± 0.4
22
12.1 ± 2.3
31
2.36 ± 0.11
--
Annealed-2 h
2.2 ± 0.3
19
10.4 ± 1.7
34
2.19 ± 0.14
--
Annealed-4 h
2.5 ± 0.5
16
8.9 ± 1.9
40
1.94 ± 0.19
--
Annealed-8 h
3.4 ± 0.8
13
7.2 ± 1.6
51
1.37 ± 0.13
0.28 ± 0.16
-2
3.3 Tensile properties Fig. 11a presents the engineering stress-strain curves of the samples annealed at 350 °C for different durations, and the tensile properties are summarized in Fig. 11b. It is obvious that the low-temperature annealing process results in a recovery of ductility and a mild decrease of strength. The as-rolled sample exhibits high ultimate tensile strength (UTS) of 700 MPa, and yield
16
strength (YS) of 670 MPa, respectively, but a low elongation of 5%. After annealing for 1 h, a pretty high strength of 665 MPa and an acceptable elongation of 10% are obtained. The ductility of the alloy gradually recovers with the increase of the annealing time at the expense of the decrease of the strength. When the sample is annealed for 8 h, the values of UTS, YS and elongation are 610 MPa, 581 MPa, and 15%, respectively. Thus, the desirable combinations of strength and ductility have been achieved after low-temperature annealing for different times. The high density of dislocations and second phase particles (continuous Ag precipitates, nano-sized Cr2Nb particles, and micro-sized Cu4AgZr particles) are extensively observed in the as-rolled sample and are mainly responsible for the improvement of the strength of the alloy. However, the high-density dislocations in the rolled alloy also restrict the strain hardening during tensile tests, which results in a low elongation. After annealing at low temperature for a short time, the dislocation density reduces through absorption during recovery. However, the nano-sized Cr2Nb and micro-sized Cu4AgZr particles provide a strong pinning effect on boundary migration and significantly slow the subgrain growth during recovery. Thus, the high-density grain boundaries exert a significant effect on preventing the dislocation motion during the tensile test, which contributes to the relatively high strength. Simultaneously, the reduction of dislocation density in the annealed sample provides new space for dislocation generation and accumulation during the tensile test, which improves the strain hardening of the alloy before necking compared to the as-rolled sample, resulting in the enhancement of the ductility. With prolonging the annealing time, the dislocation annihilation and grain growth due to the extend recovery and partial recrystallization lead to the decrease of strength. Besides, the slight coarsening of the Cu4AgZr particles and the appearance of the discontinuous Ag precipitates during annealing also 17
have an adverse effect on the strength. It is worth mentioning that twin boundaries formed during annealing perform the similar role as the grain boundaries on hindering the dislocation motion, and exert a good ability to accumulate dislocation and improve the ductility during tensile test [41, 42].
Fig. 11. (a) Engineering stress-strain curves and (b) tensile properties of the samples annealed at 350 °C for different time. Compared with the Cu3Ag0.5Zr alloy prepared by various routes [4, 43-45], the Cu3Ag0.5Zr0.4Cr0.35Nb alloy in the present work exhibits superior tensile properties, as demonstrated in Fig. 12, which is more suitable for practical utility. The introduction of the small amount of Cr2Nb particles and the fine microstructure obtained through rapid solidification followed by cold rolling are responsible for the excellent tensile properties.
18
Fig. 12. Comparison in yield strength and elongation of the Cu3Ag0.5Zr0.4Cr0.35Nb alloy and Cu3Ag0.5Zr alloys prepared by various routes. Fig. 13 shows the fracture surfaces of the samples before and after annealing, and the morphologies of the fracture surfaces are dominated by dimple features. The large amounts of shallow dimples with the size 0.2-1.0 µm are observed in the as-rolled sample (Fig. 13a), demonstrating a high strength but a low elongation of this alloy. After annealing at 350 °C for 0.5 h, some bigger and deeper ductile dimples on the fracture surface can be seen in Fig. 13b, indicating a recovered elongation. The dimples in the samples annealed for over 4 h distribute more uniformly, and the size of the dimples increases evidently, which agrees with the lower strength but higher elongation after a long annealing time.
19
Fig. 13. SEM fractographs of the samples annealed at 350 °C for (a) 0 h, (b) 0.5 h, (c) 4 h and (d) 8 h. 4 Conclusions (1) The Cu3Ag0.5Zr0.4Cr0.35Nb alloy fabricated through gas atomization and hot isostatic pressing followed by cold rolling shows a uniform microstructure with continuous Ag precipitates, nano-sized Cr2Nb particles, and micro-sized Cu4AgZr particles, and exhibits a high strength of 700 MPa but a low elongation of 5%. (2) The desirable combinations of strength and ductility have been achieved after low-temperature annealing. When annealed at 350 °C for 1 h, the alloy exhibits a high strength of 665 MPa and an acceptable elongation of 10%. The ductility of the alloy gradually recovers with the increase of annealing time at the expense of the decreased strength. Prolonging the duration to 8 h, the strength decreases to 610 MPa, while a high elongation of 15% is obtained.
20
(3) The retention of a relatively high strength after annealing at a short time is mainly attributed to the pinning effect on grain boundary migration of the nano-sized Cr2Nb and micro-sized Cu4AgZr particles, which significantly slows the subgrain growth during recovery. With prolonging the annealing time, the dislocation annihilation and subgrain growth during the extend recovery and partial recrystallization lead to the decrease of strength. Besides, the slight coarsening of the Cu4AgZr particles and the appearance of the discontinuous Ag precipitates after annealing for 8 h also have an adverse effect on the strength. The recovered ability to accumulate dislocations of the large number of particles and grain boundaries results in the enhancement of the ductility. Besides, the annealing twins also play an important role in accumulating dislocations and improving the ductility. Acknowledgement This work was supported by the Science and Technology Plan Projects of Hunan Province, China [2017GK2261].
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Highlights Cu3Ag0.5Zr0.4Cr0.35Nb alloy is prepared by rapid solidification and cold rolling. Desirable combinations of strength and ductility after annealing are achieved. The stable Cr2Nb and Cu4AgZr particles prevent the grain growth during annealing. Enhanced ductility is attributed to the recovered ability to accumulate dislocation.
No conflict of interest exits in this paper, and the paper is approved by all authors for publication. I would like to declare that the work described in this paper is our original research that has not been published in any form previously, and not under consideration for publication elsewhere, in whole or in part. All the authors listed have approved the paper that is enclosed.