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Microstructure and texture evolution in the cryorolled CuZr alloy
Accepted Manuscript Microstructure and texture evolution in the cryorolled CuZr alloy Rengeng Li, Shaojian Zhang, Huijun Kang, Zongning Chen, Fenfen Yang, Wei Wang, Cunlei Zou, Tingju Li, Tongmin Wang PII:
S0925-8388(16)32963-2
DOI:
10.1016/j.jallcom.2016.09.209
Reference:
JALCOM 39043
To appear in:
Journal of Alloys and Compounds
Received Date: 17 July 2016 Revised Date:
19 September 2016
Accepted Date: 20 September 2016
Please cite this article as: R. Li, S. Zhang, H. Kang, Z. Chen, F. Yang, W. Wang, C. Zou, T. Li, T. Wang, Microstructure and texture evolution in the cryorolled CuZr alloy, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.09.209. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
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Microstructure and texture evolution in the cryorolled CuZr alloy
Rengeng Li1, Shaojian Zhang1, Huijun Kang2*, Zongning Chen2, 3, Fenfen Yang1,
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Wei Wang1, Cunlei Zou1, Tingju Li2, Tongmin Wang1*
Key Laboratory of Solidification Control and Digital Preparation Technology
(Liaoning Province), School of Materials Science and Engineering, Dalian University
2
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of Technology, Dalian, 116024, China
Laboratory of Special Processing of Raw Materials, School of Materials Science and Engineering, Dalian University of Technology, Dalian, 116024, China School of Chemical Engineering, Dalian University of Technology, Dalian 116024,
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3
China
Corresponding email: [email protected] (Huijun Kang) [email protected]
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(Tongmin Wang)
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Abstract The microstructure and texture evolution of the room temperature and cryogenically rolled Cu-0.3%Zr alloy were investigated at a range of true strains (0.36,
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0.69, 1.20, 2.30) by transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The high angle boundary spacing of lamellar grains along the
normal direction is mainly dependent on the geometrical effect, and cryorolling has a
minor refining effect on the width of the lamellae. However, the deformation twins
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and shear bands within the lamellar structure of the cryorolled CuZr alloy contribute
to the grain refinement and improve the mechanical properties. The cryorolling enhances the inhomogeneous deformation and forms brass-type shear bands. In
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contrast to the copper-type texture in the room temperature CuZr alloy, the cryorolled CuZr alloy exhibits a typical brass-type texture. This texture discrepancy is mainly attributed to the closely spaced twin/matrix lamellae and brass-type shear bands induced by cryorolling. The twin/matrix lamellae stimulate the initial formation of brass-type texture, and the shear bands promote its development. Unlike the room
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temperature rolled CuZr alloy, the Goss component is observed in the remnant grains in the cryorolled CuZr alloy. The occurrence of the Goss component may be attributed to the suppression of cross slip during cryorolling.
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Keywords ： CuZr alloy; cryorolling; deformation twins; shear bands; electron
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backscatter diffraction
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1. Introduction The strength and electrical conductivity are mutually exclusive in nature, and a simultaneous achievement of high strength and high electrical conductivity is still
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facing a tough challenge in practice [1-3]. Among the various methods, cryorolling and subsequent aging is one of the most promising approaches for fabricating excellent balanced high strength and high electrical conductivity copper alloys in large-scale industrial application [2]. For copper alloys, it is well documented that the
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cryorolling can suppress dynamic recovery and stimulate deformation twinning, and thus promote further grain refinement and strengthen the materials [4-6]. Although
nanocrystalline pure copper was not obtained, a refining effect on the grain size was
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obtained during cryorolling compared with the room temperature (RT) rolling, as reported by Lapeire et al. [7]. The limited refining effects may result from the relatively high stacking fault energy (SFE) of pure copper. In fact, lowering the SFE of materials can promote the dissociation of dislocations into partials and thus activate profuse deformation twinning [8]. For the face-centered cubic (fcc) metals with
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medium and low SFE, deformation twinning is an additional source of grain subdivision, thereby facilitating the formation of ultrafine or even nanoscale grains [9-11]. It is also documented that the shear banding plays a key role in grain refinement at a higher strain [12, 13]. Although series of research works have already
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been reported on the cryorolling of solid-solution strengthened copper alloys with low SFE, such as CuZn [8, 14] and CuAl alloys [15], little attention is paid on the
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precipitation strengthened CuZr alloys. Generally, two types of rolling textures are observed in the fcc metals, namely
copper-type (or pure metal type) and brass-type (or alloy-type). The copper-type texture usually contains equally strong S {123} <634>, Brass {110} <112> and Copper {112} <111> components, which exists in the high and medium SFE metals; the brass-type texture is usually featured by a quite pronounced Brass {110} <112> component, which exists in the low SFE metals [16, 17]. The deformation twinning and shear banding are believed to be responsible for the texture transition from copper-type to brass-type [18]. Therefore, logic would seem to have indicated that 3
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lowering the SFE of materials and decreasing the deformation temperature might facilitate the texture transition to brass-type. Konkova et al. [19] and Lapeire et al. [7] observed a typical brass-type texture in cryorolled pure copper, which is distinguished
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from the copper-type texture in the RT rolled pure copper. However, it remains ambiguous that whether the deformation twinning or the suppression of cross-slip determines the texture transition. Contrary to pure copper, the texture evolution of
CuZn alloys during RT and cryogenic rolling was found to be broadly similar due to
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their low SFE [8]. It should be noted that the effects of cryorolling on the texture
evolution mainly depend on the SFE of the materials. Recent works have indicated that the incorporation of Zr element into copper contributes to lower the SFE of the
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alloy [20-22]. Specifically, the SFE of Cu-0.3%Zr alloy falls in between that of pure copper (78 mJ/m2) [14] and Cu-30%Zn alloys (14 mJ/m2) [23]. Therefore, it is necessary to investigate the effects of cryorolling on the texture evolution of CuZr alloy.
In our previous work, while the microstructure of CuZr alloy has been
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characterized by transmission electron microscopy (TEM) [2], its texture evolution during RT and cryogenic rolling remains unclear. Compared with TEM, electron backscatter diffraction (EBSD) enables to provide statistically representative data collecting from larger areas. The main aim of the present work is to investigate the
alloy.
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effects of cryorolling on the hardness, microstructure and texture evolution of CuZr
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2. Methods
The Cu-Zr alloy with a nominal composition of 0.3 wt.% Zr was prepared by
employing a vacuum induction furnace. After homogenizing at 900 °C for 5 hours and removing the oxidation layer, the samples were hot rolled from 30 mm to 20 mm with 33.33% thickness reduction at 850 °C. The hot rolled samples were solution treated at 972 °C for 1 hour prior to water quenching. The solution treated samples were defined as initial material. After removing the oxidation layer and the surface defects, the initial material was cryorolled with 30%, 50%, 70% and 90% overall thickness reductions (true strain: 0.36, 0.69, 1.20, 2.30), respectively. The room temperature 4
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rolled samples were also prepared for comparison. Prior to cryorolling, each sample was dipped in liquid nitrogen (-196 °C) for 10 min. The samples were rolled with 10% reduction for each pass. After each pass, the plates were dipped in liquid nitrogen
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immediately and immersed for 5 min before further deformation. All the EBSD observations were performed on the longitudinal plane containing rolling direction (RD) and normal direction (ND). The microstructure and textural
observations were taken on the mid-thickness section of the rolled samples. Samples
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for EBSD analysis were mechanically ground and polished, followed by the vibration polishing. To eliminate the stress completely, the vibration polishing was conducted about 12 hours. EBSD characterization was conducted employing the SUPRA55 field
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emission gun scanning electron microscope (FEG-SEM) equipped with HKL EBSD system. All the EBSD analyses were conducted using HKL Channel 5 software. Due to the misorientation noise, boundary misorientations were cut-off at less than 2°. A misorientation threshold of 15° was used to differentiate low angle boundaries (LABs) from high angle boundaries (HABs). The twin boundaries were defined by ∑ 3
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coincidence site lattice (CSL) boundaries following the Brandon’s criterion, and the fraction of twin boundaries was defined by the percentage of the total boundary length. In order to acquire reliable texture representations, more than 4000 grains were analyzed in each EBSD map. For simplicity, only three orientation distribution
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function (ODF) sections (φ2=0°, 45°, 65°) are presented, which is sufficient to characterize the typical fcc rolling texture components. Fig. 1 shows ODF sections
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(with φ2 =0°, 45°, 65°) of the main ideal orientations existing in rolled fcc materials [7]. The Miller indices and Euler angles of ideal texture components are listed in Table 1 [24].
Fig. 1 ODF sections (with φ2 =0°, 45°, 65°) of the main ideal orientations existing 5
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in rolled fcc materials. Table 1 Miller indices and Euler angles of ideal texture components. Euler angles φ1, Φ, φ2
Cube
{001} <100>
0°
0°
Copper
{112} <111>
90°
35°
S
{123} <634>
59°
34°
Brass
{110} <112>
35°
45°
Goss
{011} <100>
0°
45°
Dillamore
{4 4 11} <11 11 8>
90°
Y
{111} <112>
30°/90°
Z
{111} <110>
0°/60°
00°
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Miller indices
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Designation
45°
65°
0°/90°
0°/90° 45°
55°
45°
55°
45°
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27°
Discs of 3 mm in diameter punched from the longitudinal section of the plates were ground to 30 µm and then double jet thinned using a 25% nitric acid in methanol solution at - 30 °C. TEM observations were performed using a transmission electron
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microscope operating at 300 kV (Tecnai G2 F30). Vickers micro-hardness (HV) measurements were conducted by applying a load of 200 g for 15 s on the mid-thickness section using a MH-50 type micro-hardness tester. These hardness measurements were carried out ten times at least for each condition and the hardness
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values were calculated by averaging these test results.
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3. Results and discussion
3.1 Microstructure of initial material and micro-hardness variations during rolling
Fig. 2 illustrates the EBSD inverse pole figure map of initial material with the ND
corresponding to the crystal reference system. The HABs are depicted in black lines. The initial microstructure was characterized by nearly equiaxed grains with annealing twins. The fraction of twin boundaries is 38%. The ODF sections (with φ2 =0°, 45°, 65°) of the initial material are shown in Fig. 3. The initial material exhibits a weak 6
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Cube {001} <100> component with a maximum of 4.13× random. For preliminary comparison of microstructure evolution during RT and cryogenic rolling, the hardness values of initial material and rolled samples with various true strains are shown in Fig.
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4. It can be observed that the hardness of both the RT rolled (RTR) and cryogenically rolled (CR) samples increases sharply after a small strain. Compared with the initial material, the hardness values of the CR and RTR sample nearly double at the true
strain of 0.36. With further increase in strain, the increase in hardness becomes
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RTR samples at the whole range of true strain.
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relatively slow. In addition, the hardness values of CR samples are higher than that of
Fig. 2 EBSD inverse pole figure map of initial material. The color code triangle is
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shown in the upper right corner.
Fig. 3 ODF sections (with φ2 =0°, 45°, 65°) of the initial material.
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Fig. 4 Variation of hardness during RT and cryogenic rolling. The error bars represent the standard deviation.
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3.2 Microstructure evolution
The microstructure evolution during RT and cryogenic rolling is shown in Fig. 5. The HABs are depicted in black lines. At the true strain of 0.36, elongated grains along the RD are observed. With increase in strain, the HAB spacing along the ND decreases and the grain boundaries tend to extend parallel to RD. An extremely elongated lamellar structure can be observed at the true strain of 2.3. In a macroscopic
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view, it seems that the RTR and CR CuZr alloy exhibit a similar tendency in grain compression during rolling. For comparison, the average HAB spacing of the lamellar grains is shown in Fig. 6. Theoretically, the reduction in HAB spacing is in proportion to the sample reduction ratio, and thus the HAB spacing of the lamellar grains, λG, can
λG =λ0 exp(-εtrue )
(1)
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be described by [9]
where λ0 is the initial grain width, and εtrue is the true strain. The HAB spacings thus
predicted are also given in Fig. 6 (dashed line). The average HAB spacings of both the RTR and CR CuZr alloy nearly follow the theoretically predicted curve. That is to say, the geometrical effect significantly influences the reduction of the HAB spacing of lamellar grains during the RT and cryogenic rolling. The dynamic grain coarsening is not obvious in the RTR CuZr alloy, which is inconsistent with the previous report [19]. This phenomenon may be attributed to the fact that the final grain size of the RTR CuZr alloy (about 10 µm along the ND) is far from the minimum grain size, 8
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which can be achieved by RT rolling.
Fig. 5 Inverse pole figure maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) RTR, true strain 0.69; (c) RTR, true strain 1.20; (d) RTR, true strain 2.30; (e) CR, true strain 0.36; (f) CR, true strain 0.69;
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(g) CR, true strain 1.20; (h) CR, true strain 2.30. The color code triangle is
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shown in the bottom right corner.
Fig. 6 The average HAB spacing of the lamellar grains as a function of true strain. The error bars represent the standard deviation. Although cryorolling has a minor refining effect on the HAB spacing, the microstructure within the lamellar structure of the CR CuZr alloy differs from that of 9
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the RTR samples. Fig. 7 illuminates the band contrast maps of the RTR and CR CuZr alloy at different true strains. The ∑ 3 type twin boundaries are shown in red line and some shear bands are depicted in dashed lines. Table 2 summarizes the fraction of
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twin boundaries in the RTR and CR CuZr alloy at various true strains. After a small rolling reduction (strain 0.36), the fraction of twin boundaries in both the RTR and CR
CuZr decreases dramatically. It reveals that the annealing twin boundaries are
destroyed significantly after a small rolling reduction. The fraction of twin boundaries
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in the CR CuZr alloy increases from 3.96% to 7.16% at the true strain of 0.69. It is
obvious that deformation twinning has been activated at the true strain of 0.69 in the CR CuZr alloy. With increase in strain, the fraction of twin boundaries in the CR
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CuZr alloy begins to decrease. This phenomenon can be ascribed to the relatively coarse scanning step (600 nm) in this case, which is insufficient to detect the nanoscale deformation twins. In addition, the deformation twins tend to bend and fragment at high strain level [25], which induces the significant deviation of deformation twin boundaries from the ideal ∑ 3 type boundaries. However, an
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increase in the fraction of twin boundaries is observed in the RTR CuZr alloy with increase in strain. The flow stress in the RTR CuZr alloy may be insufficient to induce the fragment of the deformation twins. It should be noted that the fraction of deformation twins in the CR CuZr samples is higher than that in the RTR samples.
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For the high-resolution observations, the TEM technique is employed to characterize the microstructure at high strain level. Fig. 8 shows the bright field TEM images of
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the RTR and CR CuZr alloy at the true strain of 2.30. The inset in Fig. 8 shows the corresponding selected area electron diffraction (SAED) pattern, which confirms the existence of deformation twins. Profuse deformation twins in bundles are observed in the CR CuZr alloy, whereas fewer deformation twins exist in the RTR samples. In addition, the twin/matrix lamellar thickness of CR samples (48 nm) decreases compared with that of RTR samples (87 nm). In contrast to the CR pure copper [19], the annealing twins are not observed in both the RTR and CR CuZr alloy at the true strain of 2.30. Because of the different crystal structure of Zr (hexagonal close packed) and Cu (fcc), the addition of Zr can significantly decrease the SFE of the copper 10
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matrix [21]. The lower SFE of the CuZr alloy can decrease the twin boundary energy and thus decrease the critical twin nucleus thickness [26, 27]. Compared with pure copper, it is easier to form deformation twins in the CuZr alloy. What’s more,
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decreasing the rolling temperature can lead to the increase of driving stress for twin nucleation [26]. Therefore, the addition of Zr and decrease of rolling temperature
stimulate the formation of the closely spaced twin/matrix lamellae. It should be noted that the recrystallized grains are not observed in both the RTR and CR CuZr alloy in
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contrast to the CR pure copper [7]. It can be deduced that the addition of Zr in small
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quantities significantly improves the thermal stability of the CuZr alloy.
Fig. 7 Band contrast maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.69; (b) RTR, true strain 1.20; (c) RTR, true strain 2.30; (d) CR, true strain 0.69; (e) CR, true strain 1.20; (f) CR, true strain 2.30. The ∑ 3
type twin boundaries are shown in red line.
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Table 2 The fraction of twin boundaries in the RTR and CR CuZr alloy Strain 0.36 4.46 3.96
Strain 0.69 Strain 1.20 1.13 1.32 7.16 5.38
Strain 2.30 2.16 2.40
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Rolling condition Room temperature rolled Cryorolled
Fig. 8 Bright field TEM images of the RTR (a) and CR (b) CuZr alloy at the true strain of 2.30. The inset shows the corresponding SAED pattern. Another obvious difference between the RTR and CR CuZr alloy from the band
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contrast maps (Fig. 7) is the occurrence of shear bands. The formation of shear banding is found at the true strain of 1.30 and 0.69 for the RTR and CR CuZr alloy, respectively. Shear banding is usually considered as a consequence of localized
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deformation. Once work hardening cannot be achieved by homogenous plastic deformation, shear banding will dominate the deformation mode [10]. Obviously, cryorolling enhances the inhomogeneous deformation compared with the RT rolling.
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The shear bands observed in the RTR CuZr alloy at the high strain level are inclined at about 35° to the RD. This kind of shear bands usually exists in heavy deformed metals with a high or medium SFE, especially in heavily deformed copper, and can be considered as copper-type shear bands [10]. In contrast to the RTR CuZr alloy, shear bands with different inclination angles corresponding to different strain levels are observed in the CR samples. At the low strain level, the shear bands are mainly inclined at about 25° to the RD (as shown in Fig. 7d). The 35° inclination angle to the RD is found in the CR CuZr alloy at the true strain of 1.20, which is as same as the 12
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RTR samples. At the true strain of 2.30, two sets of shear bands, inclined to the RD at about 30° and 40°, respectively, intersect with each other and form a rhomboidal prism with an axis parallel to the TD (as shown in dashed lines in Fig. 7f and Fig. 8b).
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This kind of shear bands cuts through the deformation twin regimes and divides the deformation twin regimes into several portions (Fig. 8b). These features of the shear bands in the CR CuZr alloy completely match with the brass-type shear bands.
Therefore, decreasing rolling temperature contributes to the transition of shear bands
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from copper-type to brass-type.
The extensive deformation twinning and shear banding in the CR CuZr alloy divide the lamellar grains into finer subgrains or even grains [2], and thus the mechanical
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properties of the CR CuZr alloy are improved compared with the RTR samples (Fig. 4). However, there are still some coarse remnants of original grains in both the RTR and CR CuZr alloy (as marked in Fig. 5). These remnant grains free of deformation twins and shear bands are the consequences of inhomogeneous deformation, and deteriorate the mechanical properties, which will be discussed in section 3.4.
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3.3 Texture evolution
The textures in deformed fcc materials are usually described by several fibers, such as α-fiber, β-fiber, γ-fiber and so on. The α-fiber and γ-fiber are often specified with <110>∥ND and <111>∥ND, respectively [28, 29]. The β-fiber is commonly defined
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as an orientation tube connecting the Copper {112} <111> and Brass {110} <112> orientation via the S {123} <634> components [7]. However, this description fails to
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define the precise position of β-fiber. According to Sidor et al., the β-fiber can be described exactly in the first Euler subspaces by the following expression [16]: {h, 1, h+1} <
hh+1 /-h
,
2hh+1 / -h
,
h2 h-3/4
+
2h h-1/2
>
(2)
The ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains show the crystallographic orientations in details (Fig. 9). Because orthorhombic sample symmetry is observed in both the RTR and CR CuZr alloy at the whole strain range, Euler angle φ1 with a range of 0-90° is shown in the ODF sections. To compare the texture evolution quantitatively during the RT and cryogenic rolling, 13
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the distributions of texture intensity along α-fiber and β-fiber are shown in Fig. 10 and 11, respectively. After a small rolling reduction (strain 0.36), pronounced α-fiber textures are developed in both the RTR and CR CuZr alloy (Fig. 5 and 9). The Goss
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{011} <100> component with a maximum intensity of 9.08 × random is observed in the RTR CuZr alloy, whereas nearly equally strong Goss {011} <100> and Brass {110} <112> components with an intensity of 6 × random appear in the CR CuZr alloy at the
true strain of 0.36 (Fig. 9a and b). With increase in strain, the intensities of overall
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α-fiber and β-fiber textures tend to strengthen. The maximum intensity along the
α-fiber shifts from {011} <211> to {011} <311> (Fig. 10), which lies in the 15° tolerance of Brass orientation. The intensity of Goss orientation tends to fade
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accompanied by the occurrence of Copper and S components along β-fiber in the RTR CuZr alloy. In contrast to RT rolling, the intensity of Goss component tends to strengthen during cryorolling. What’s more, the texture intensities along the β-fiber in the RTR CuZr alloy seem more pronounced than that in the CR CuZr alloy. It is also observed that γ-fiber tends to increase accompanied by the occurrence of shear bands
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(Fig. 5f). The volume fraction of γ-fiber in the CR CuZr alloy increases to 9.44 % at the true strain of 2.30, which is double that in the RTR samples. Equally intensive Copper, S and Brass orientations are the dominant components in the RTR CuZr alloy at the true strain of 2.30, which is the characteristic of copper-type texture [30]. In
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contrast, a pronounced Brass component, together with Goss and S components exists in the CR CuZr alloy at the true strain of 2.30, which can be considered as brass-type
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texture. Therefore, the cryorolling of the CuZr alloy promotes the texture transition from copper-type to brass-type.
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Fig. 9 ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true
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strain 1.20; (g) RTR, true strain 2.30; (h) CR, true strain 2.30.
Fig. 10 Texture intensity distribution along α-fiber of the RTR (a) and CR (b)
CuZr alloy.
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Fig. 11 Texture intensity distribution along β-fiber of the RTR and CR CuZr alloy. (a) RTR, Euler angle φ2=45° ~ 90°; (b) CR, Euler angle φ2=45° ~ 90°.
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Although the texture transition from the copper-type to brass-type was also observed in the CR pure copper [7, 19], the formation mechanism of the brass-type texture during cryorolling is still not clear. Konkova et al. ascribed the texture transition to the suppression of the cross-slip [19], whereas Lapeire et al. argued that the occurrence of deformation twins leads to the texture transition [7]. To analyze the mechanism of texture transition, the distributions of ideal texture components (within
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15° tolerance) in the RTR and CR samples with different true strains are shown in Fig. 12. It can be observed that the grains with Copper orientation are prone to undergo deformation twinning and shear banding, which is consisted with the previous report [18]. The Y component begins to occur in the CR CuZr alloy at the true strain of 0.69
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(as marked by ellipse in Fig. 12f), whereas only a small fraction of Y component is observed in the RTR CuZr alloy at the true strain of 2.30 (Fig. 12d). The formation of
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Y component is mainly related to the closely spaced twin/matrix lamellae. It is difficult to activate deformation twinning or dislocation slipping within the twin/matrix lamellae with increase in strain. A slip system is then activated in the interval between the twin bundles, and the slipping on the plane parallel to the twin/matrix lamellae induces the rotation of twin/matrix lamellae [31]. Once the twin/matrix lamellae rotate to the direction parallel to the RD, Y component forms in the Copper oriented grains. Therefore, the cryorolling of the CuZr alloy can stimulate the deformation twins and form the closely spaced twin/matrix lamellae at the true 16
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strain of 0.69. Although the Y component is observed in the RTR CuZr alloy at the true strain of 2.30, most of the Copper oriented grains are absence of Y component. The thickness and volume fraction of twin/matrix lamellae of RTR samples are not
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sufficient to induce the required degree of latent hardening, which is responsible for this phenomenon [31].
Apart from deformation twins, the type of shear bands also plays a significant role
in the texture evolution. As shown in Fig. 7 and 12, the brass-type shear bands are
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formed in the areas with Y orientation. With the onset of shear banding, the
twin/matrix lamellae within the shear bands rotate toward Goss orientation (as marked by rectangle in Fig. 12g), whereas the twin/matrix lamellae outside the shear bands
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remain Y orientation, which is consisted with the previous report [32]. With increase in strain, new slip systems are activated, and thus the Goss orientation will rotate toward Brass and S orientations (as marked by rectangle in Fig. 12h). Therefore, Brass, Goss and S components dominate the shear banding regions in the CR sample at the true strain of 2.30 (Fig. 12h). In contrast to the CR CuZr alloy, the Copper
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orientation still dominates the shear banding regions in the RTR CuZr alloy (as marked by rectangle in Fig. 12d). It reveals that the copper-type shear bands have a minor effect on the Copper orientation. In terms of the RTR CuZr alloy, the Copper orientation firstly rotates to Dillamore orientation after a small rolling reduction. Once
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the copper-type shear bands form, the Dillamore orientation starts to rotate back to the initial Copper orientation [33]. From the discussion above, the closely spaced
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twin/matrix lamellae induced by cryorolling motivate the formation of brass-type texture, and the brass-type shear bands promote the development of brass-type texture at the later stage.
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Fig. 12 EBSD maps of the ideal texture components (within 15° tolerance). (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20; (g) RTR, true strain 2.30; (h) CR, true strain 2.30. The shear banding regions are marked by
3.4 Remnant grains
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rectangle.
As discussed above, the remnant grains in both the RTR and CR CuZr alloy
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stagnate the grain refinement and thus restrict the improvement of mechanical properties. As shown in Fig. 12d and h, the orientation of the remnant grains in the
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RTR CuZr alloy is mainly Brass component, whereas strong Brass and Goss components dominate the remnant grains in the CR samples. The Brass and Goss orientations are usually known to have a low ratio of the twinning-to-slipping Schmid factors [12], and thus deformation twinning is hard to occur in the grains with Brass and Goss orientations. Therefore, dislocation slip is the dominant deformation mode in the remnant grains. In general, the Brass component is considered as a stable end orientation, whereas the Goss component is believed to be a transient orientation and can rotate to Brass orientation via dislocation slipping [12, 18]. Lowering the SFE and decreasing the deformation temperature can dissociate the dislocations into partials 18
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and thus restrict the dislocation cross slip. Therefore, the orientation transition from Goss to Brass may be hindered by the suppression of cross slip during cryorolling.
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4. Conclusions The evolution of microstructure and texture of the cryogenically and RT rolled
CuZr alloy has been investigated over a range of true strains. The main conclusions can be drawn as follows:
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(1) The HAB spacing of lamellar grains is mainly dependent on the geometrical effect. Cryorolling has a minor refining effect on the HAB spacing. However, the deformation twins and shear bands within the lamellar structure of the CR CuZr alloy
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contribute to the grain refinement and improve the mechanical properties. (2) Cryorolling promotes the transition of shear bands from copper-type to brass-type. The brass-type shear bands are inclined at a variety of angles, with a range between 25° and 40°, to the RD.
(3) The closely spaced twin/matrix lamellae induced by cryorolling stimulate the
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initial formation of brass-type texture, and the brass-type shear bands promote further development to brass-type texture.
(4) Unlike the RTR CuZr alloy, the Goss component is also observed in the remnant grains in the CR CuZr alloy. The occurrence of the Goss component may be attributed
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to the suppression of cross slip during cryorolling.
Acknowledgements
The authors gratefully acknowledge the supports of National Key Research and
Development Program of China (No.2016YFB0701200), National Natural Science Foundation of China (Nos. 51525401, 51274054, U1332115, 51401044, 51601028), the China Postdoctoral Science Foundation (2015M581331), and the Fundamental Research Funds for the Central Universities. 19
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Author contributions
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R. Li, H. Kang and T. Wang conceived the idea and designed the experiment. S. Zhang, Z. Chen, F. Yang performed the experiments. S. Zhang, H. Kang, W. Wang, C. Zou and T. Li contributed to the results analysis and discussion. R. Li wrote the paper.
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Competing financial interests
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The authors declare no competing financial interests.
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3799-3812.
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Figure captions Fig. 1 ODF sections (with φ2 =0°, 45°, 65°) of the main ideal orientations existing in rolled fcc materials.
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Fig. 2 EBSD inverse pole figure map of initial material. The color code triangle is shown in the upper right corner.
Fig. 3 ODF sections (with φ2 =0°, 45°, 65°) of the initial material.
Fig. 4 Variation of hardness during RT and cryogenic rolling. The error bars
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represent the standard deviation.
Fig. 5 Inverse pole figure maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) RTR, true strain 0.69; (c) RTR, true strain 1.20;
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(d) RTR, true strain 2.30; (e) CR, true strain 0.36; (f) CR, true strain 0.69; (g) CR, true strain 1.20; (h) CR, true strain 2.30. The color code triangle is shown in the bottom right corner.
Fig. 6 The average HAB spacing of the lamellar grains as a function of true strain. The error bars represent the standard deviation.
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Fig. 7 Band contrast maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.69; (b) RTR, true strain 1.20; (c) RTR, true strain 2.30; (d) CR, true strain 0.69; (e) CR, true strain 1.20; (f) CR, true strain 2.30. The ∑ 3 type twin boundaries are shown in red line.
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Fig. 8 Bright field TEM images of the RTR (a) and CR (b) CuZr alloy at the true strain of 2.30. The inset shows the corresponding SAED pattern.
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Fig. 9 ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true
strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20;
(g) RTR, true strain 2.30; (h) CR, true strain 2.30. Fig. 10 Texture intensity distribution along α-fiber of the RTR (a) and CR (b) CuZr alloy. Fig. 11 Texture intensity distribution along β-fiber of the RTR and CR CuZr
alloy. (a) RTR, Euler angle φ2=45° ~ 90°; (b) CR, Euler angle φ2=45° ~ 90°. Fig. 12 EBSD maps of the ideal texture components (within 15° tolerance). (a) 23
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RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20; (g) RTR, true strain 2.30;
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(h) CR, true strain 2.30. The shear banding regions are marked by rectangle.
Table captions
Table 1 Miller indices and Euler angles of ideal texture components.
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Table 2 The fraction of twin boundaries in the RTR and CR CuZr alloy
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The texture evolution of CuZr during ambient and cryogenic rolling was studied. The mechanism of the texture transition induced by cryorolling was analyzed. The cryorolling of CuZr alloys promotes the formation of brass-type shear bands. The suppression of cross slip induces the Goss orientation in the remnant grains.
S0925-8388(16)32963-2
DOI:
10.1016/j.jallcom.2016.09.209
Reference:
JALCOM 39043
To appear in:
Journal of Alloys and Compounds
Received Date: 17 July 2016 Revised Date:
19 September 2016
Accepted Date: 20 September 2016
Please cite this article as: R. Li, S. Zhang, H. Kang, Z. Chen, F. Yang, W. Wang, C. Zou, T. Li, T. Wang, Microstructure and texture evolution in the cryorolled CuZr alloy, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.09.209. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
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Microstructure and texture evolution in the cryorolled CuZr alloy
Rengeng Li1, Shaojian Zhang1, Huijun Kang2*, Zongning Chen2, 3, Fenfen Yang1,
1
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Wei Wang1, Cunlei Zou1, Tingju Li2, Tongmin Wang1*
Key Laboratory of Solidification Control and Digital Preparation Technology
(Liaoning Province), School of Materials Science and Engineering, Dalian University
2
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of Technology, Dalian, 116024, China
Laboratory of Special Processing of Raw Materials, School of Materials Science and Engineering, Dalian University of Technology, Dalian, 116024, China School of Chemical Engineering, Dalian University of Technology, Dalian 116024,
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3
China
Corresponding email: [email protected] (Huijun Kang) [email protected]
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(Tongmin Wang)
1
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Abstract The microstructure and texture evolution of the room temperature and cryogenically rolled Cu-0.3%Zr alloy were investigated at a range of true strains (0.36,
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0.69, 1.20, 2.30) by transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The high angle boundary spacing of lamellar grains along the
normal direction is mainly dependent on the geometrical effect, and cryorolling has a
minor refining effect on the width of the lamellae. However, the deformation twins
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and shear bands within the lamellar structure of the cryorolled CuZr alloy contribute
to the grain refinement and improve the mechanical properties. The cryorolling enhances the inhomogeneous deformation and forms brass-type shear bands. In
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contrast to the copper-type texture in the room temperature CuZr alloy, the cryorolled CuZr alloy exhibits a typical brass-type texture. This texture discrepancy is mainly attributed to the closely spaced twin/matrix lamellae and brass-type shear bands induced by cryorolling. The twin/matrix lamellae stimulate the initial formation of brass-type texture, and the shear bands promote its development. Unlike the room
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temperature rolled CuZr alloy, the Goss component is observed in the remnant grains in the cryorolled CuZr alloy. The occurrence of the Goss component may be attributed to the suppression of cross slip during cryorolling.
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Keywords ： CuZr alloy; cryorolling; deformation twins; shear bands; electron
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backscatter diffraction
2
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1. Introduction The strength and electrical conductivity are mutually exclusive in nature, and a simultaneous achievement of high strength and high electrical conductivity is still
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facing a tough challenge in practice [1-3]. Among the various methods, cryorolling and subsequent aging is one of the most promising approaches for fabricating excellent balanced high strength and high electrical conductivity copper alloys in large-scale industrial application [2]. For copper alloys, it is well documented that the
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cryorolling can suppress dynamic recovery and stimulate deformation twinning, and thus promote further grain refinement and strengthen the materials [4-6]. Although
nanocrystalline pure copper was not obtained, a refining effect on the grain size was
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obtained during cryorolling compared with the room temperature (RT) rolling, as reported by Lapeire et al. [7]. The limited refining effects may result from the relatively high stacking fault energy (SFE) of pure copper. In fact, lowering the SFE of materials can promote the dissociation of dislocations into partials and thus activate profuse deformation twinning [8]. For the face-centered cubic (fcc) metals with
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medium and low SFE, deformation twinning is an additional source of grain subdivision, thereby facilitating the formation of ultrafine or even nanoscale grains [9-11]. It is also documented that the shear banding plays a key role in grain refinement at a higher strain [12, 13]. Although series of research works have already
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been reported on the cryorolling of solid-solution strengthened copper alloys with low SFE, such as CuZn [8, 14] and CuAl alloys [15], little attention is paid on the
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precipitation strengthened CuZr alloys. Generally, two types of rolling textures are observed in the fcc metals, namely
copper-type (or pure metal type) and brass-type (or alloy-type). The copper-type texture usually contains equally strong S {123} <634>, Brass {110} <112> and Copper {112} <111> components, which exists in the high and medium SFE metals; the brass-type texture is usually featured by a quite pronounced Brass {110} <112> component, which exists in the low SFE metals [16, 17]. The deformation twinning and shear banding are believed to be responsible for the texture transition from copper-type to brass-type [18]. Therefore, logic would seem to have indicated that 3
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lowering the SFE of materials and decreasing the deformation temperature might facilitate the texture transition to brass-type. Konkova et al. [19] and Lapeire et al. [7] observed a typical brass-type texture in cryorolled pure copper, which is distinguished
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from the copper-type texture in the RT rolled pure copper. However, it remains ambiguous that whether the deformation twinning or the suppression of cross-slip determines the texture transition. Contrary to pure copper, the texture evolution of
CuZn alloys during RT and cryogenic rolling was found to be broadly similar due to
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their low SFE [8]. It should be noted that the effects of cryorolling on the texture
evolution mainly depend on the SFE of the materials. Recent works have indicated that the incorporation of Zr element into copper contributes to lower the SFE of the
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alloy [20-22]. Specifically, the SFE of Cu-0.3%Zr alloy falls in between that of pure copper (78 mJ/m2) [14] and Cu-30%Zn alloys (14 mJ/m2) [23]. Therefore, it is necessary to investigate the effects of cryorolling on the texture evolution of CuZr alloy.
In our previous work, while the microstructure of CuZr alloy has been
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characterized by transmission electron microscopy (TEM) [2], its texture evolution during RT and cryogenic rolling remains unclear. Compared with TEM, electron backscatter diffraction (EBSD) enables to provide statistically representative data collecting from larger areas. The main aim of the present work is to investigate the
alloy.
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effects of cryorolling on the hardness, microstructure and texture evolution of CuZr
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2. Methods
The Cu-Zr alloy with a nominal composition of 0.3 wt.% Zr was prepared by
employing a vacuum induction furnace. After homogenizing at 900 °C for 5 hours and removing the oxidation layer, the samples were hot rolled from 30 mm to 20 mm with 33.33% thickness reduction at 850 °C. The hot rolled samples were solution treated at 972 °C for 1 hour prior to water quenching. The solution treated samples were defined as initial material. After removing the oxidation layer and the surface defects, the initial material was cryorolled with 30%, 50%, 70% and 90% overall thickness reductions (true strain: 0.36, 0.69, 1.20, 2.30), respectively. The room temperature 4
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rolled samples were also prepared for comparison. Prior to cryorolling, each sample was dipped in liquid nitrogen (-196 °C) for 10 min. The samples were rolled with 10% reduction for each pass. After each pass, the plates were dipped in liquid nitrogen
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immediately and immersed for 5 min before further deformation. All the EBSD observations were performed on the longitudinal plane containing rolling direction (RD) and normal direction (ND). The microstructure and textural
observations were taken on the mid-thickness section of the rolled samples. Samples
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for EBSD analysis were mechanically ground and polished, followed by the vibration polishing. To eliminate the stress completely, the vibration polishing was conducted about 12 hours. EBSD characterization was conducted employing the SUPRA55 field
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emission gun scanning electron microscope (FEG-SEM) equipped with HKL EBSD system. All the EBSD analyses were conducted using HKL Channel 5 software. Due to the misorientation noise, boundary misorientations were cut-off at less than 2°. A misorientation threshold of 15° was used to differentiate low angle boundaries (LABs) from high angle boundaries (HABs). The twin boundaries were defined by ∑ 3
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coincidence site lattice (CSL) boundaries following the Brandon’s criterion, and the fraction of twin boundaries was defined by the percentage of the total boundary length. In order to acquire reliable texture representations, more than 4000 grains were analyzed in each EBSD map. For simplicity, only three orientation distribution
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function (ODF) sections (φ2=0°, 45°, 65°) are presented, which is sufficient to characterize the typical fcc rolling texture components. Fig. 1 shows ODF sections
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(with φ2 =0°, 45°, 65°) of the main ideal orientations existing in rolled fcc materials [7]. The Miller indices and Euler angles of ideal texture components are listed in Table 1 [24].
Fig. 1 ODF sections (with φ2 =0°, 45°, 65°) of the main ideal orientations existing 5
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in rolled fcc materials. Table 1 Miller indices and Euler angles of ideal texture components. Euler angles φ1, Φ, φ2
Cube
{001} <100>
0°
0°
Copper
{112} <111>
90°
35°
S
{123} <634>
59°
34°
Brass
{110} <112>
35°
45°
Goss
{011} <100>
0°
45°
Dillamore
{4 4 11} <11 11 8>
90°
Y
{111} <112>
30°/90°
Z
{111} <110>
0°/60°
00°
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Miller indices
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Designation
45°
65°
0°/90°
0°/90° 45°
55°
45°
55°
45°
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27°
Discs of 3 mm in diameter punched from the longitudinal section of the plates were ground to 30 µm and then double jet thinned using a 25% nitric acid in methanol solution at - 30 °C. TEM observations were performed using a transmission electron
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microscope operating at 300 kV (Tecnai G2 F30). Vickers micro-hardness (HV) measurements were conducted by applying a load of 200 g for 15 s on the mid-thickness section using a MH-50 type micro-hardness tester. These hardness measurements were carried out ten times at least for each condition and the hardness
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values were calculated by averaging these test results.
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3. Results and discussion
3.1 Microstructure of initial material and micro-hardness variations during rolling
Fig. 2 illustrates the EBSD inverse pole figure map of initial material with the ND
corresponding to the crystal reference system. The HABs are depicted in black lines. The initial microstructure was characterized by nearly equiaxed grains with annealing twins. The fraction of twin boundaries is 38%. The ODF sections (with φ2 =0°, 45°, 65°) of the initial material are shown in Fig. 3. The initial material exhibits a weak 6
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Cube {001} <100> component with a maximum of 4.13× random. For preliminary comparison of microstructure evolution during RT and cryogenic rolling, the hardness values of initial material and rolled samples with various true strains are shown in Fig.
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4. It can be observed that the hardness of both the RT rolled (RTR) and cryogenically rolled (CR) samples increases sharply after a small strain. Compared with the initial material, the hardness values of the CR and RTR sample nearly double at the true
strain of 0.36. With further increase in strain, the increase in hardness becomes
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RTR samples at the whole range of true strain.
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relatively slow. In addition, the hardness values of CR samples are higher than that of
Fig. 2 EBSD inverse pole figure map of initial material. The color code triangle is
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shown in the upper right corner.
Fig. 3 ODF sections (with φ2 =0°, 45°, 65°) of the initial material.
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Fig. 4 Variation of hardness during RT and cryogenic rolling. The error bars represent the standard deviation.
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3.2 Microstructure evolution
The microstructure evolution during RT and cryogenic rolling is shown in Fig. 5. The HABs are depicted in black lines. At the true strain of 0.36, elongated grains along the RD are observed. With increase in strain, the HAB spacing along the ND decreases and the grain boundaries tend to extend parallel to RD. An extremely elongated lamellar structure can be observed at the true strain of 2.3. In a macroscopic
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view, it seems that the RTR and CR CuZr alloy exhibit a similar tendency in grain compression during rolling. For comparison, the average HAB spacing of the lamellar grains is shown in Fig. 6. Theoretically, the reduction in HAB spacing is in proportion to the sample reduction ratio, and thus the HAB spacing of the lamellar grains, λG, can
λG =λ0 exp(-εtrue )
(1)
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be described by [9]
where λ0 is the initial grain width, and εtrue is the true strain. The HAB spacings thus
predicted are also given in Fig. 6 (dashed line). The average HAB spacings of both the RTR and CR CuZr alloy nearly follow the theoretically predicted curve. That is to say, the geometrical effect significantly influences the reduction of the HAB spacing of lamellar grains during the RT and cryogenic rolling. The dynamic grain coarsening is not obvious in the RTR CuZr alloy, which is inconsistent with the previous report [19]. This phenomenon may be attributed to the fact that the final grain size of the RTR CuZr alloy (about 10 µm along the ND) is far from the minimum grain size, 8
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which can be achieved by RT rolling.
Fig. 5 Inverse pole figure maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) RTR, true strain 0.69; (c) RTR, true strain 1.20; (d) RTR, true strain 2.30; (e) CR, true strain 0.36; (f) CR, true strain 0.69;
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(g) CR, true strain 1.20; (h) CR, true strain 2.30. The color code triangle is
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shown in the bottom right corner.
Fig. 6 The average HAB spacing of the lamellar grains as a function of true strain. The error bars represent the standard deviation. Although cryorolling has a minor refining effect on the HAB spacing, the microstructure within the lamellar structure of the CR CuZr alloy differs from that of 9
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the RTR samples. Fig. 7 illuminates the band contrast maps of the RTR and CR CuZr alloy at different true strains. The ∑ 3 type twin boundaries are shown in red line and some shear bands are depicted in dashed lines. Table 2 summarizes the fraction of
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twin boundaries in the RTR and CR CuZr alloy at various true strains. After a small rolling reduction (strain 0.36), the fraction of twin boundaries in both the RTR and CR
CuZr decreases dramatically. It reveals that the annealing twin boundaries are
destroyed significantly after a small rolling reduction. The fraction of twin boundaries
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in the CR CuZr alloy increases from 3.96% to 7.16% at the true strain of 0.69. It is
obvious that deformation twinning has been activated at the true strain of 0.69 in the CR CuZr alloy. With increase in strain, the fraction of twin boundaries in the CR
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CuZr alloy begins to decrease. This phenomenon can be ascribed to the relatively coarse scanning step (600 nm) in this case, which is insufficient to detect the nanoscale deformation twins. In addition, the deformation twins tend to bend and fragment at high strain level [25], which induces the significant deviation of deformation twin boundaries from the ideal ∑ 3 type boundaries. However, an
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increase in the fraction of twin boundaries is observed in the RTR CuZr alloy with increase in strain. The flow stress in the RTR CuZr alloy may be insufficient to induce the fragment of the deformation twins. It should be noted that the fraction of deformation twins in the CR CuZr samples is higher than that in the RTR samples.
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For the high-resolution observations, the TEM technique is employed to characterize the microstructure at high strain level. Fig. 8 shows the bright field TEM images of
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the RTR and CR CuZr alloy at the true strain of 2.30. The inset in Fig. 8 shows the corresponding selected area electron diffraction (SAED) pattern, which confirms the existence of deformation twins. Profuse deformation twins in bundles are observed in the CR CuZr alloy, whereas fewer deformation twins exist in the RTR samples. In addition, the twin/matrix lamellar thickness of CR samples (48 nm) decreases compared with that of RTR samples (87 nm). In contrast to the CR pure copper [19], the annealing twins are not observed in both the RTR and CR CuZr alloy at the true strain of 2.30. Because of the different crystal structure of Zr (hexagonal close packed) and Cu (fcc), the addition of Zr can significantly decrease the SFE of the copper 10
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matrix [21]. The lower SFE of the CuZr alloy can decrease the twin boundary energy and thus decrease the critical twin nucleus thickness [26, 27]. Compared with pure copper, it is easier to form deformation twins in the CuZr alloy. What’s more,
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decreasing the rolling temperature can lead to the increase of driving stress for twin nucleation [26]. Therefore, the addition of Zr and decrease of rolling temperature
stimulate the formation of the closely spaced twin/matrix lamellae. It should be noted that the recrystallized grains are not observed in both the RTR and CR CuZr alloy in
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contrast to the CR pure copper [7]. It can be deduced that the addition of Zr in small
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quantities significantly improves the thermal stability of the CuZr alloy.
Fig. 7 Band contrast maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.69; (b) RTR, true strain 1.20; (c) RTR, true strain 2.30; (d) CR, true strain 0.69; (e) CR, true strain 1.20; (f) CR, true strain 2.30. The ∑ 3
type twin boundaries are shown in red line.
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Table 2 The fraction of twin boundaries in the RTR and CR CuZr alloy Strain 0.36 4.46 3.96
Strain 0.69 Strain 1.20 1.13 1.32 7.16 5.38
Strain 2.30 2.16 2.40
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Rolling condition Room temperature rolled Cryorolled
Fig. 8 Bright field TEM images of the RTR (a) and CR (b) CuZr alloy at the true strain of 2.30. The inset shows the corresponding SAED pattern. Another obvious difference between the RTR and CR CuZr alloy from the band
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contrast maps (Fig. 7) is the occurrence of shear bands. The formation of shear banding is found at the true strain of 1.30 and 0.69 for the RTR and CR CuZr alloy, respectively. Shear banding is usually considered as a consequence of localized
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deformation. Once work hardening cannot be achieved by homogenous plastic deformation, shear banding will dominate the deformation mode [10]. Obviously, cryorolling enhances the inhomogeneous deformation compared with the RT rolling.
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The shear bands observed in the RTR CuZr alloy at the high strain level are inclined at about 35° to the RD. This kind of shear bands usually exists in heavy deformed metals with a high or medium SFE, especially in heavily deformed copper, and can be considered as copper-type shear bands [10]. In contrast to the RTR CuZr alloy, shear bands with different inclination angles corresponding to different strain levels are observed in the CR samples. At the low strain level, the shear bands are mainly inclined at about 25° to the RD (as shown in Fig. 7d). The 35° inclination angle to the RD is found in the CR CuZr alloy at the true strain of 1.20, which is as same as the 12
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RTR samples. At the true strain of 2.30, two sets of shear bands, inclined to the RD at about 30° and 40°, respectively, intersect with each other and form a rhomboidal prism with an axis parallel to the TD (as shown in dashed lines in Fig. 7f and Fig. 8b).
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This kind of shear bands cuts through the deformation twin regimes and divides the deformation twin regimes into several portions (Fig. 8b). These features of the shear bands in the CR CuZr alloy completely match with the brass-type shear bands.
Therefore, decreasing rolling temperature contributes to the transition of shear bands
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from copper-type to brass-type.
The extensive deformation twinning and shear banding in the CR CuZr alloy divide the lamellar grains into finer subgrains or even grains [2], and thus the mechanical
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properties of the CR CuZr alloy are improved compared with the RTR samples (Fig. 4). However, there are still some coarse remnants of original grains in both the RTR and CR CuZr alloy (as marked in Fig. 5). These remnant grains free of deformation twins and shear bands are the consequences of inhomogeneous deformation, and deteriorate the mechanical properties, which will be discussed in section 3.4.
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3.3 Texture evolution
The textures in deformed fcc materials are usually described by several fibers, such as α-fiber, β-fiber, γ-fiber and so on. The α-fiber and γ-fiber are often specified with <110>∥ND and <111>∥ND, respectively [28, 29]. The β-fiber is commonly defined
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as an orientation tube connecting the Copper {112} <111> and Brass {110} <112> orientation via the S {123} <634> components [7]. However, this description fails to
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define the precise position of β-fiber. According to Sidor et al., the β-fiber can be described exactly in the first Euler subspaces by the following expression [16]: {h, 1, h+1} <
hh+1 /-h
,
2hh+1 / -h
,
h2 h-3/4
+
2h h-1/2
>
(2)
The ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains show the crystallographic orientations in details (Fig. 9). Because orthorhombic sample symmetry is observed in both the RTR and CR CuZr alloy at the whole strain range, Euler angle φ1 with a range of 0-90° is shown in the ODF sections. To compare the texture evolution quantitatively during the RT and cryogenic rolling, 13
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the distributions of texture intensity along α-fiber and β-fiber are shown in Fig. 10 and 11, respectively. After a small rolling reduction (strain 0.36), pronounced α-fiber textures are developed in both the RTR and CR CuZr alloy (Fig. 5 and 9). The Goss
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{011} <100> component with a maximum intensity of 9.08 × random is observed in the RTR CuZr alloy, whereas nearly equally strong Goss {011} <100> and Brass {110} <112> components with an intensity of 6 × random appear in the CR CuZr alloy at the
true strain of 0.36 (Fig. 9a and b). With increase in strain, the intensities of overall
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α-fiber and β-fiber textures tend to strengthen. The maximum intensity along the
α-fiber shifts from {011} <211> to {011} <311> (Fig. 10), which lies in the 15° tolerance of Brass orientation. The intensity of Goss orientation tends to fade
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accompanied by the occurrence of Copper and S components along β-fiber in the RTR CuZr alloy. In contrast to RT rolling, the intensity of Goss component tends to strengthen during cryorolling. What’s more, the texture intensities along the β-fiber in the RTR CuZr alloy seem more pronounced than that in the CR CuZr alloy. It is also observed that γ-fiber tends to increase accompanied by the occurrence of shear bands
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(Fig. 5f). The volume fraction of γ-fiber in the CR CuZr alloy increases to 9.44 % at the true strain of 2.30, which is double that in the RTR samples. Equally intensive Copper, S and Brass orientations are the dominant components in the RTR CuZr alloy at the true strain of 2.30, which is the characteristic of copper-type texture [30]. In
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contrast, a pronounced Brass component, together with Goss and S components exists in the CR CuZr alloy at the true strain of 2.30, which can be considered as brass-type
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texture. Therefore, the cryorolling of the CuZr alloy promotes the texture transition from copper-type to brass-type.
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Fig. 9 ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true
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strain 1.20; (g) RTR, true strain 2.30; (h) CR, true strain 2.30.
Fig. 10 Texture intensity distribution along α-fiber of the RTR (a) and CR (b)
CuZr alloy.
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Fig. 11 Texture intensity distribution along β-fiber of the RTR and CR CuZr alloy. (a) RTR, Euler angle φ2=45° ~ 90°; (b) CR, Euler angle φ2=45° ~ 90°.
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Although the texture transition from the copper-type to brass-type was also observed in the CR pure copper [7, 19], the formation mechanism of the brass-type texture during cryorolling is still not clear. Konkova et al. ascribed the texture transition to the suppression of the cross-slip [19], whereas Lapeire et al. argued that the occurrence of deformation twins leads to the texture transition [7]. To analyze the mechanism of texture transition, the distributions of ideal texture components (within
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15° tolerance) in the RTR and CR samples with different true strains are shown in Fig. 12. It can be observed that the grains with Copper orientation are prone to undergo deformation twinning and shear banding, which is consisted with the previous report [18]. The Y component begins to occur in the CR CuZr alloy at the true strain of 0.69
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(as marked by ellipse in Fig. 12f), whereas only a small fraction of Y component is observed in the RTR CuZr alloy at the true strain of 2.30 (Fig. 12d). The formation of
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Y component is mainly related to the closely spaced twin/matrix lamellae. It is difficult to activate deformation twinning or dislocation slipping within the twin/matrix lamellae with increase in strain. A slip system is then activated in the interval between the twin bundles, and the slipping on the plane parallel to the twin/matrix lamellae induces the rotation of twin/matrix lamellae [31]. Once the twin/matrix lamellae rotate to the direction parallel to the RD, Y component forms in the Copper oriented grains. Therefore, the cryorolling of the CuZr alloy can stimulate the deformation twins and form the closely spaced twin/matrix lamellae at the true 16
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strain of 0.69. Although the Y component is observed in the RTR CuZr alloy at the true strain of 2.30, most of the Copper oriented grains are absence of Y component. The thickness and volume fraction of twin/matrix lamellae of RTR samples are not
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sufficient to induce the required degree of latent hardening, which is responsible for this phenomenon [31].
Apart from deformation twins, the type of shear bands also plays a significant role
in the texture evolution. As shown in Fig. 7 and 12, the brass-type shear bands are
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formed in the areas with Y orientation. With the onset of shear banding, the
twin/matrix lamellae within the shear bands rotate toward Goss orientation (as marked by rectangle in Fig. 12g), whereas the twin/matrix lamellae outside the shear bands
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remain Y orientation, which is consisted with the previous report [32]. With increase in strain, new slip systems are activated, and thus the Goss orientation will rotate toward Brass and S orientations (as marked by rectangle in Fig. 12h). Therefore, Brass, Goss and S components dominate the shear banding regions in the CR sample at the true strain of 2.30 (Fig. 12h). In contrast to the CR CuZr alloy, the Copper
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orientation still dominates the shear banding regions in the RTR CuZr alloy (as marked by rectangle in Fig. 12d). It reveals that the copper-type shear bands have a minor effect on the Copper orientation. In terms of the RTR CuZr alloy, the Copper orientation firstly rotates to Dillamore orientation after a small rolling reduction. Once
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the copper-type shear bands form, the Dillamore orientation starts to rotate back to the initial Copper orientation [33]. From the discussion above, the closely spaced
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twin/matrix lamellae induced by cryorolling motivate the formation of brass-type texture, and the brass-type shear bands promote the development of brass-type texture at the later stage.
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Fig. 12 EBSD maps of the ideal texture components (within 15° tolerance). (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20; (g) RTR, true strain 2.30; (h) CR, true strain 2.30. The shear banding regions are marked by
3.4 Remnant grains
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rectangle.
As discussed above, the remnant grains in both the RTR and CR CuZr alloy
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stagnate the grain refinement and thus restrict the improvement of mechanical properties. As shown in Fig. 12d and h, the orientation of the remnant grains in the
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RTR CuZr alloy is mainly Brass component, whereas strong Brass and Goss components dominate the remnant grains in the CR samples. The Brass and Goss orientations are usually known to have a low ratio of the twinning-to-slipping Schmid factors [12], and thus deformation twinning is hard to occur in the grains with Brass and Goss orientations. Therefore, dislocation slip is the dominant deformation mode in the remnant grains. In general, the Brass component is considered as a stable end orientation, whereas the Goss component is believed to be a transient orientation and can rotate to Brass orientation via dislocation slipping [12, 18]. Lowering the SFE and decreasing the deformation temperature can dissociate the dislocations into partials 18
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and thus restrict the dislocation cross slip. Therefore, the orientation transition from Goss to Brass may be hindered by the suppression of cross slip during cryorolling.
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4. Conclusions The evolution of microstructure and texture of the cryogenically and RT rolled
CuZr alloy has been investigated over a range of true strains. The main conclusions can be drawn as follows:
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(1) The HAB spacing of lamellar grains is mainly dependent on the geometrical effect. Cryorolling has a minor refining effect on the HAB spacing. However, the deformation twins and shear bands within the lamellar structure of the CR CuZr alloy
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contribute to the grain refinement and improve the mechanical properties. (2) Cryorolling promotes the transition of shear bands from copper-type to brass-type. The brass-type shear bands are inclined at a variety of angles, with a range between 25° and 40°, to the RD.
(3) The closely spaced twin/matrix lamellae induced by cryorolling stimulate the
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initial formation of brass-type texture, and the brass-type shear bands promote further development to brass-type texture.
(4) Unlike the RTR CuZr alloy, the Goss component is also observed in the remnant grains in the CR CuZr alloy. The occurrence of the Goss component may be attributed
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to the suppression of cross slip during cryorolling.
Acknowledgements
The authors gratefully acknowledge the supports of National Key Research and
Development Program of China (No.2016YFB0701200), National Natural Science Foundation of China (Nos. 51525401, 51274054, U1332115, 51401044, 51601028), the China Postdoctoral Science Foundation (2015M581331), and the Fundamental Research Funds for the Central Universities. 19
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Author contributions
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R. Li, H. Kang and T. Wang conceived the idea and designed the experiment. S. Zhang, Z. Chen, F. Yang performed the experiments. S. Zhang, H. Kang, W. Wang, C. Zou and T. Li contributed to the results analysis and discussion. R. Li wrote the paper.
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Competing financial interests
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The authors declare no competing financial interests.
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Figure captions Fig. 1 ODF sections (with φ2 =0°, 45°, 65°) of the main ideal orientations existing in rolled fcc materials.
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Fig. 2 EBSD inverse pole figure map of initial material. The color code triangle is shown in the upper right corner.
Fig. 3 ODF sections (with φ2 =0°, 45°, 65°) of the initial material.
Fig. 4 Variation of hardness during RT and cryogenic rolling. The error bars
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represent the standard deviation.
Fig. 5 Inverse pole figure maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) RTR, true strain 0.69; (c) RTR, true strain 1.20;
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(d) RTR, true strain 2.30; (e) CR, true strain 0.36; (f) CR, true strain 0.69; (g) CR, true strain 1.20; (h) CR, true strain 2.30. The color code triangle is shown in the bottom right corner.
Fig. 6 The average HAB spacing of the lamellar grains as a function of true strain. The error bars represent the standard deviation.
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Fig. 7 Band contrast maps of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.69; (b) RTR, true strain 1.20; (c) RTR, true strain 2.30; (d) CR, true strain 0.69; (e) CR, true strain 1.20; (f) CR, true strain 2.30. The ∑ 3 type twin boundaries are shown in red line.
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Fig. 8 Bright field TEM images of the RTR (a) and CR (b) CuZr alloy at the true strain of 2.30. The inset shows the corresponding SAED pattern.
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Fig. 9 ODF sections (with φ2 =0°, 45°, 65°) of the RTR and CR CuZr alloy at various true strains. (a) RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true
strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20;
(g) RTR, true strain 2.30; (h) CR, true strain 2.30. Fig. 10 Texture intensity distribution along α-fiber of the RTR (a) and CR (b) CuZr alloy. Fig. 11 Texture intensity distribution along β-fiber of the RTR and CR CuZr
alloy. (a) RTR, Euler angle φ2=45° ~ 90°; (b) CR, Euler angle φ2=45° ~ 90°. Fig. 12 EBSD maps of the ideal texture components (within 15° tolerance). (a) 23
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RTR, true strain 0.36; (b) CR, true strain 0.36; (c) RTR, true strain 0.69; (d) CR, true strain 0.69; (e) RTR, true strain 1.20; (f) CR, true strain 1.20; (g) RTR, true strain 2.30;
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(h) CR, true strain 2.30. The shear banding regions are marked by rectangle.
Table captions
Table 1 Miller indices and Euler angles of ideal texture components.
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Table 2 The fraction of twin boundaries in the RTR and CR CuZr alloy
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The texture evolution of CuZr during ambient and cryogenic rolling was studied. The mechanism of the texture transition induced by cryorolling was analyzed. The cryorolling of CuZr alloys promotes the formation of brass-type shear bands. The suppression of cross slip induces the Goss orientation in the remnant grains.