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Anisotropic response of cold sprayed copper deposits
Accepted Manuscript Anisotropic response of cold sprayed copper deposits
Kang Yang, Wenya Li, Xiawei Yang, Yaxin Xu PII: DOI: Reference:
S0257-8972(17)31266-5 https://doi.org/10.1016/j.surfcoat.2017.12.043 SCT 22961
To appear in:
Surface & Coatings Technology
Received date: Revised date: Accepted date:
21 August 2017 18 December 2017 19 December 2017
Please cite this article as: Kang Yang, Wenya Li, Xiawei Yang, Yaxin Xu , Anisotropic response of cold sprayed copper deposits. The address for the corresponding author was captured as affiliation for all authors. Please check if appropriate. Sct(2017), https://doi.org/10.1016/j.surfcoat.2017.12.043
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ACCEPTED MANUSCRIPT Anisotropic response of cold sprayed copper deposits Kang Yang, Wenya Li *, Xiawei Yang, Yaxin Xu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, P.R. China
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* Corresponding: Email: [email protected];
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Tel.: +86-29-88495226; Fax.: +86-29-88492642
Abstract
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Cold spray (CS) is characterized as a solid-state process of high deposition efficiency
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for metallic coatings as well as additive manufacturing of metals. However, due to high velocity impact and extensive deformation of particles during CS, the as-received
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coatings or deposits may present anisotropic characteristics which could influence the
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performance of deposits. Hence this study aims to investigate the anisotropic behaviors of CSed copper deposits in a systematic way. The microstructure and micromechanical
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properties of the deposits both in the cross-section (v-face) and in the parallel plane to
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the surface (p-face) were characterized. Tensile tests were performed at various loading angles with respect to the nozzle moving direction in the p-face. It is revealed that there exist strong microstructural and mechanical anisotropies in CSed deposits. Different interparticle interaction results in more severe particle impact deformation in v-face than p-face, with larger elastic modulus and microhardness values. The tensile tests show an unexpected anisotropy in both ultimate tensile strength and elongation, with the highest performance occurring at the angle of 20°. The in-plane tensile anisotropy could
ACCEPTED MANUSCRIPT be attributed to the parallel multiple passes. Therefore, a novel weave-spraying method was proposed, which can greatly reduce the tensile anisotropy of CSed deposits. Keywords: Anisotropic response; Cold spray; Tensile property; Weave-spraying
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1. Introduction
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Cold spray (CS) has attracted increasing attention as a versatile solid-state coating
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and additive manufacturing technique [1,2]. In this process, substrates are exposed to a high velocity (300-1200m/s) jet of small particles accelerated by an expanding gas
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stream at temperatures lower than the melting point of spray material [1,2]. Beyond a
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critical velocity, which is governed by material properties and process conditions, a mechanical interlocking or metallurgical bonding is produced. Unlike conventional
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thermal spray (TS), these conditions identify CS as a low-temperature and high-speed
most cases [1-3].
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process with significant advantages such as lack of oxidation or phase transformation in
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Without considering assistant processes, the particle impact velocity and temperature
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play key roles in forming effective bonding between particle/particle and particle/substrate, which are governed by process parameters such as type of gas, gas pressure, gas temperature, nozzle geometry and standoff distance [4,5]. Up to now, a lot of investigations based on particles impact behavior and its effect on bonding have been carried out [1,5]. Because of these unique advantages, a variety of materials including most metals, alloys, composites, ceramics and even nanostructured metallic materials have been
ACCEPTED MANUSCRIPT deposited by CS technique [6-10]. To date, a number of papers have been published on mechanical characteristics of CSed coatings [8,11-13]. After CS, a significant increase of hardness can be obtained compared to feedstock powder, which is due to the strong work hardening effect [13]. Researches of CSed copper, aluminum, titanium and their
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alloys [11-13] show that an excellent tensile strength can be obtained but with a very
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low ductility. These testing specimens are usually limited to one direction parallel or
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vertical to spray gun traversing, and then the testing results of such specimens have been used to characterize the mechanical response of CSed deposits [2,8-13]. However,
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this kind of characterization cannot describe the complex mechanical behaviors of CSed
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deposits, which can result in a significant deviation from the actual mechanical properties. Because of the inhomogeneous particle impact deformation behavior and
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existence of particle interfaces (especially between spraying passes), anisotropies of
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microstructure and mechanical properties may be developed. From previous studies, it is known that strong plastic anisotropies exist in many
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solid-state techniques with severe deformation such as cold rolling [14], extrusion [15]
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and friction stir welding/friction stir processing [16,17]. In the study of friction stir welding [16,17], it is demonstrated that the textures developed in the joint include Cube {001}⟨100⟩ , Cube ND {001}⟨110⟩ , S {123}⟨634⟩ , Brass {011}⟨211⟩ and Goss {011}⟨100⟩. And the strong Cube texture generally resulted in the highest yield stress at the loading angle of 45° because of its pronounced and enduring effect on strain localization. However, only a few attempts have been made on the anisotropic response of CSed
ACCEPTED MANUSCRIPT deposits to date [18-20]. Bolelli et al. [18] employed depth-sensing indentation to assess the mechanical properties of CSed Ta, which reported some differences in micromechanical properties (hardness and elastic modulus) of cross-section and surface. Coddet et al. [19] and Seiner et al. [20] investigated the elastic and electrical
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anisotropies of coatings prepared by CS, and the results revealed that only weak
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deviations from perfect isotropy exhibited in CSed deposits. The above authors have
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carried out comprehensive studies on some anisotropic properties of CS. However, the microstructural and mechanical anisotropies caused by impact behaviors have not been
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well understood.
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As for the optimization design and reliable performance evaluations of CSed coatings (e.g. the formability and quality of CSed deposits), it is of great importance to consider
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the anisotropies of microstructure and mechanical properties in CS because of its
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significant effect on stress-strain distribution. Therefore, the main aim of this work is to characterize the anisotropy between two faces and in-plane tensile anisotropy in CSed
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copper deposits. The contributions of particle interfaces on all-directional tensile
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strengths were evaluated. Finally, a weave-spraying pattern was proposed to modify the in-plane tensile anisotropy of CS.
2. Experimental 2.1 Deposits fabrication The powder used in this study is commercially spherical Cu (Beijing You Xing Lian Technology Co., Ltd, China). Figs. 1a and 1c show the morphology of Cu powder and
ACCEPTED MANUSCRIPT cross-sectional view of Cu particles. The particle size distribution characterized by a laser diffraction sizer (MASTERSIZER 2000, Malvern Instruments Ltd., UK) is shown in Fig. 1b with a median size (D50) of 26.2μm. 3mm thick Cu plates were used as substrates. A custom developed CS system was employed to perform deposition. The
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maximum temperature and pressure of the system are 800°C and 1MPa, respectively.
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The nozzle had an expansion ratio of 6.7 and a divergent section length of 200mm. In
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this experiment, helium was used as the accelerating gas with an inlet pressure of 0.8MPa and temperature of around 500°C. The nozzle standoff distance was 25mm and
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gun traverse speed was 20mm/s. Following 8-passes spraying, 8mm thick Cu deposit
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2.2 Materials characterizations
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was obtained.
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The microstructure of the cross-section (v-face) and the plane parallel to the surface (p-face) was examined with electron back-scattered diffraction (EBSD) performed on a
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scanning electron microscope (SEM, JSM5800LV, JEOL, Japan). The v-face and p-face
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are shown in Fig. 2. The microstructure was also examined with an optical microscope (OM, OLYMPUS GX71, Japan). Depth sensing indentation technique was used to determine micromechanical properties using a Nanoindenter G200 (Agilent Technologies Inc, USA). Indentations with a constant load (10mN) were carried out on splats. Eight indentations were tested in each case, and after ignoring the indentations with extreme values (maximum and minimum values) six lines are shown. The Vickers hardness of both faces was measured with a digital microhardness tester (Struers
ACCEPTED MANUSCRIPT Duramin-A300, Denmark) with a load of 200g and dwell time of 15s. Twelve indentations were taken in each face with ignoring the indentations with extreme values (maximum and minimum values). Both nanoindentation and Vickers hardness tests of two faces were carried out along measurement lines as shown in Fig. 2. The in-plane
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tensile properties in p-face were measured at six directions (0°, 15°, 30°, 45°, 60°, 90°)
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with a tensile test machine (INSTRON-3382, America), where 0° is the direction
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parallel to the spraying gun traverse as shown in Fig. 3b. Three samples were tested to evaluate the tensile properties for each direction. It should be noted that standard
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deviation was applied to get error bars of the Vickers hardness and tensile properties. A
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micro-tensile specimen with 2mm thickness was used [21], where the width in the gauge section and total length were 2mm and 25mm, respectively (Fig. 3a). A specially
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designed fixture (Fig. 3c) was used to ensure correct application of the load and prevent
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secondary bending or other unwanted movement of the specimens. The fracture surfaces of tensile specimens were also observed with SEM. An X-Ray diffractometer system
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(XRD, Siemens D500, Germany) was used to measure microstrain and crystallite size
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with 40kV and 30mA of Cu Kα radiation at a wavelength of 0.1542nm. The scan was conducted in 2θ mode and across a range of 10° to 120° at a step resolution of 0.033°.
3. Results and discussion 3.1 Anisotropy between p-face and v-face 3.1.1
Microstructure
The EBSD maps of v-face and p-face are shown in Figs. 4a and 4b, respectively. It
ACCEPTED MANUSCRIPT can be seen that both particle shapes and grain size distributions are different between v-face and p-face. Particle plastic behavior such as impact and extrusion are more concentrated on v-face than p-face. As a consequence, particle deformation in v-face is more severe to form flat shapes with disc shapes in p-face. This phenomenon can also
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be found in OM images as shown in Fig. 5. Because of the anisotropic impact, grain
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distribution characteristics in the two cases are clearly different. In comparison to the
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p-face, grain distribution in v-face is inhomogeneous, with more refined grains near the particle impact interface. This is due to the recrystallization which has been confirmed
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by previous studies [8,22]. The statistical result of misorientation angle distributions
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(Fig. 4c) shows little difference of grain boundaries fractions between two faces (including High-Angle Grain Boundaries (HAGBs, θ≥15°) and Low-Angle Grain
3.1.2
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Boundaries (LAGBs, 15°>θ≥2°)).
Nanoindentation and Vickers hardness
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The anisotropic particle deformation of v-face and p-face was quantified with depth
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sensing indentation as shown in Figs. 6a and 6b, respectively. Compared to p-face, micromechanical properties of v-face show a remarkable increase of 13.5% for modulus and 11.8% for microhardness. This can be the consequence of the more severe plastic deformation in v-face than that in p-face. The nanoindentation results are consistent with the studies on CSed Ta by Bolelli et al. [18], having some differences (hardness and elastic modulus) between cross-section and surface. But the differences between two faces in our study are greater compared to the above mentioned study [18], which
ACCEPTED MANUSCRIPT could be mainly due to different materials. Fig. 7 shows the Vickers hardness of v-face and p-face. It should be noted that the average microhardness value of pure Cu was measured as 70HV. Both the microhardness values of v-face and p-face are higher than pure Cu, which has been
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confirmed due to the strong work hardening effect during CS [8,11]. It can be seen that
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the Vickers hardness values of two faces are much closer than that of nanoindentation.
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The difference of two test results could be due to the different measuring principles and size effect [23,24]. Compared to nanoindentation, the scope of Vickers hardness test is
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relatively much large, and the Vickers hardness is strongly affected by the material
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porosity and adhesion strength between particles [23].
In-plane tensile anisotropy
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3.2.1
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3.2 In-plane tensile anisotropy and strength contributions
Being a significant mechanical performance, tensile properties in six directions were
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also measured as shown in Figs. 8a and 8b. It can be observed that the stress-strain
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curves of CSed Cu deposits change with loading angles, indicating a strong anisotropy in both tensile strength and elongation. With the increase of loading angle, both the ultimate tensile strength and elongation first rise and then drop with the highest values measured at angles of 15°-30°. For angles over 30°, the tensile performance deteriorates rapidly, reaching a minimum at 90°. Fig. 9 shows typical fracture surfaces of 0°, 30° and 90° tensile samples. No dimples can be found in these fracture surfaces. The fracture takes place along interfaces
ACCEPTED MANUSCRIPT between deposited particles rather than through them, indicating a brittle rupture for the CSed Cu. Consistent with tensile strengths, the morphology of 30° tensile sample is the coarsest (Figs. 9b and 9e), while that of 90° the smoothest (Figs. 9c and 9f). The interlocking zones between Cu particles are marked in the magnified views (Figs. 9d-9f).
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It can be found that the amount of interlocking zones of 30° tensile sample is the most,
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with the least of 90°.
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For the traditional spraying pattern, the spraying passes are usually parallel. The parallel spraying interfaces between passes easily result in anisotropy of CSed coatings.
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It is considered that the different tensile properties in different directions are attributed
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to the following factors: (1) Interparticle interlocking effect between spraying passes has a positive effect; (2) During CS deposition, the coating is formed through bonding
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between impact particles, while bonding quality is dependent on particle deformation
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behavior, which is determined by the surface state of the previously deposited coating. Near the interface between spraying passes as shown in Fig. 10, the normal velocity
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component (𝑉𝑛 ) is lower than the actual velocity (𝑉𝑎 ). The decrease in normal velocity
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between spraying passes indicates a weaker interfacial bonding than that within a pass, which could affect the tensile property in a negative manner. To isolate these two interfacial bonding strengths experimentally, a technique of multiple load indentation and the hardness loss parameter proposed by Goldbaum et al. [23] could be helpful. As a consequence, for angles between 0° and 30°, the positive factor “interlocking effect” is dominant. Hence the tensile properties including ultimate tensile strength and elongation increase with angles. This association reverses for larger angles between 30°
ACCEPTED MANUSCRIPT and 90°, and decreases the tensile property. It should be noted that oxides and pores exist at the interfaces between particles in CSed deposits. In addition, although bonding can happen at the interfaces, the strength is relatively low compared to the particles themselves. Therefore, the interfaces between particles are weak, which lead to the poor
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performance of CSed deposits compared to the bulk material.
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In summary, unlike the anisotropies existed in solid-state techniques like cold rolling
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[14], extrusion [15] and friction stir welding/processing [16,17], which is resulted from grain orientation, the in-plane tensile anisotropy of CS is mainly due to interparticle
3.2.2
Strength contributions
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interfaces especially between spraying passes.
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Because of the main responsibility of interfaces to tensile anisotropy, it is
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indispensable to study and quantify the changes of their contributions to the final ultimate tensile strength. Several additional laws are available for estimating the total
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strength [25,26]. In our case, the final tensile strength (𝜎𝑓−𝐴 ) depends on several
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parameters, which can be associated linearly [25,26]: 𝜎𝑓−𝐴 = 𝜎0 + 𝜎𝐻𝑃 + 𝜎𝑊𝐻 − 𝜎𝑃𝐼−𝐴
(1)
where 𝜎0 is the base strength of pure Cu (128MPa); 𝜎𝐻𝑃 is the grain size strengthening according to the Hall-Petch law; 𝜎𝑊𝐻 is the work hardening contribution; 𝜎𝑃𝐼−𝐴 is the special existence in CS, which means a negative contribution to the final tensile strength due to interfaces between particles [8]. According to literature [27], the strengthening due to grain refinement is described by
ACCEPTED MANUSCRIPT the Hall-Petch relation: 𝜎𝐻𝑃 = 𝑘𝑦 ⁄√𝐷
(2)
where 𝑘𝑦 is a constant depending on the material (37.08MPa√μm for Cu); D is the average grain size after impact, which can be measured with EBSD (1.38μm).
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Accordingly, considering the calculation from Eq. (2), the increment resulted from grain
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refinement is estimated to be 32MPa.
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The general equation to describe the effect of work hardening (dislocation) on the strength is given by [26,28]:
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𝜎𝑊𝐻 = αMGb√𝜌
(3)
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where α is a constant equal to 0.33; M is the Taylor factor (3.06) G is the shear modulus for Cu (3.9 × 104 MPa); b is the Burgers vector (0.3615 × 10−3 μm); 𝜌 is
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the dislocation density, which can be estimated by XRD as shown in Fig. 11, and it can
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be calculated by the Williamson-Hall equation [29]: (4𝑎)
ρ = 14.4𝜀 2 ⁄𝑏 2
(4𝑏)
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𝛽𝑐𝑜𝑠𝜃⁄λ = 0.9⁄𝐷 + 2𝜀sinθ⁄λ
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where 𝛽 is the peak width at half maximum height; 𝜆 is the wave length of X-ray beam (0.1542nm for Cu Kα); 𝐷 is the crystallite size; 𝜀 is the lattice strain; 𝜃 is the Bragg angle. By plotting 𝛽𝑐𝑜𝑠𝜃⁄λ against s𝑖𝑛𝜃, the slope of the graph gives 2𝜀 ⁄λ as shown in Fig. 11b, the Y-intercept gives 0.9⁄𝐷 and ρ is calculated using Eq. (4b) with a value of 4.78 × 1014 m-2. Hence, the increment contributed from dislocation strengthening is 311MPa. Therefore, according to Eq. (1), the particle interfaces contribution 𝜎𝑃𝐼−𝐴 (MPa) of
ACCEPTED MANUSCRIPT different angles can be written: 𝜎𝑃𝐼−𝐴 = 470 − 𝜎𝑓
(5)
In order to get all-angled particle interfaces effects, the 𝜎𝑃𝐼−𝐴 is fitted with Gaussian function (𝐴𝑛𝑔 is the tensile direction, unit: °): 2
𝜎𝑃𝐼−𝐴 = 440 − 60e−2((𝐴𝑛𝑔−19)⁄37)
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(6)
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The fitted function suggests that the lowest 𝜎𝑃𝐼−𝐴 occurs at angle of 20° with the
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highest final ultimate tensile strength value of about 90MPa. It can be concluded here that the negative contribution of the weak existence of interparticle interfaces (𝜎𝑃𝐼 ) is
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the main factor in CS, which hampers the improvement of tensile property. Therefore,
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weakening the existence of interparticle interfaces could further improve the tensile
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3.3 Weave-spraying
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performance in all directions.
The large tensile anisotropy in CSed deposits is not desirable, which is the result of
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multiple parallel spraying passes. Therefore, in order to reduce or eliminate the parallel
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multi-pass effect, an alternative spraying pattern was proposed: weave-spraying as shown in Fig. 12a. The tensile performance achieved with different angles of weave-sprayed Cu is shown in Fig. 12. Obviously, the novel pattern has a great suppression on the tensile anisotropy of CSed deposits. Although the high tensile values for some angles (~15°-30°) are sacrificed, an almost isotropic CSed deposit is produced by weave-spraying, producing a homogeneous material in both tensile strength and elongation. Figs. 13a-13c show the typical fracture surfaces of 0°, 30° and 90° tensile
ACCEPTED MANUSCRIPT samples, respectively, with the respective magnified views showed in Figs. 13d-13f. As expected, the fracture surfaces of the three loading angles show similar morphologies. Weave-spraying well modifies the in-plane tensile anisotropy of CS. For weave-spraying pattern, the spraying passes are orthogonal, which can effectively
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reduce the difference between different directions. In the case of a certain thickness of a
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CSed deposit, it is believed that the multi-pass effect would be further weakened with
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the increase of weave-spraying times. For TS, the effect of spray trajectory has been well investigated in many studies which show that spray trajectory can influence the
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heat and mass transfer significantly, and therefore on the coating microstructure,
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physical and chemical properties and their quality [30-32]. While for CS, only a few investigations on the nozzle trajectory have been reported so far [33,34]. Chen et al. [34]
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proposed a spiral trajectory applicable to CS for repairing damaged workpiece, which
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can improve productivity and reduce cost in machinery work. But they did not study the mechanical properties or anisotropy. In our opinion, this pattern (spiral-trajectory
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spraying) could also be a method to weaken the in-plane tensile anisotropy, because this
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spraying pattern can effectively reduce the difference between different directions. Therefore, weakening the parallel multi-pass effect between different directions is a guiding idea to eliminate the coating anisotropy.
4. Conclusions In this paper, Cu was deposited to investigate the anisotropic behaviors of CS. The existence of strong anisotropies was found in CSed deposits including microstructure
ACCEPTED MANUSCRIPT and mechanical properties, and the responsible mechanisms have been discussed. The negative contribution of the existence of weak interparticle interfaces has also been evaluated. The main conclusions can be drawn as follows: (1) Particle impact deformation and grain size distribution behaviors are different
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between the v-face and p-face, with little difference in misorientation angle
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distributions.
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(2) Nanoindentation tests show that the anisotropic impact deformation results in larger elastic modulus and hardness in the v-face than p-face.
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(3) Strong anisotropy exists in the in-plane tensile properties of CSed deposits. The
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parallel multiple passes have the main responsibility to the in-plane tensile anisotropy. The fitted function suggests that the highest tensile performance
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occurs at the angle of 20°, and the lowest at 90°.
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(4) A novel weave-spraying method was proposed, which can effectively reduce the parallel multi-pass effect between different directions. The weave-spraying has a
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great modification effect on the tensile anisotropy of CSed deposits.
Acknowledgments The authors would like to thank for the financial support from the National Key Research and Development Program of China (2016YFB1100104), the fund of SAST (SAST2016043) and the 111 Project (B08040).
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[27] M.R. Barnett, Z. Keshavarz, A.G. Beer, D. Atwell, Influence of grain size on the
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compressive deformation of wrought Mg-3Al-1Zn, Acta Mater. 52 (2004) 5093-5103.
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[28] L. Zhang, S. Ukai, T. Hoshino, S. Hayashi, X. Qu, Y2O3 evolution and dispersion
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refinement in Co-base ODS alloys, Acta Mater. 57 (2009) 3671-3682. [29] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and
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wolfram, Acta Metall. 1 (1953) 22-31.
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[30] Z.H. Cai, B.C. Qi, C.Y. Tao, J. Luo, Y.P. Chen, C.J. Xie, A robot trajectory optimization approach for thermal barrier coatings used for free-form components, J. Therm. Spray Technol. 26 (2017) 1651-1658. [31] R. Gadow, A. Candel, M. Floristán, Optimized robot trajectory generation for thermal spraying operations and high quality coatings on free-form surfaces, Surf. Coat. Technol. 205 (2010) 1074-1079. [32] S.H. Deng, H. Liang, Z.H. Cai, H.L. Liao, G. Montavon, Kinematic optimization of
ACCEPTED MANUSCRIPT robot trajectories for thermal spray coating application, J. Therm. Spray Technol. 23 (2014) 1382-1389. [33] Z.H. Cai, T.Y. Chen, C.N. Zeng, X.P. Guo, H.J. Lian, Y. Zheng, X.X. Wei, A global approach to the optimal trajectory based on an improved ant colony algorithm for
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cold spray, J. Therm. Spray Technol. 25 (2016) 1631-1637.
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[34] C.Y. Chen, S. Gojon, Y.C. Xie, S. Yin, C. Verdy, Z.M. Ren, H.L. Liao, S.H. Deng, A
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novel spiral trajectory for damage component recovery with cold spray, Surf. Coat.
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Technol. 309 (2017) 719-728.
ACCEPTED MANUSCRIPT Figure captions Fig. 1. Morphology (a), particle size distribution (b) and cross-sectional view (c) of Cu powder. Fig. 2. The sketch map of v-face, p-face and measurement lines. Note: both nanoindentation and Vickers hardness tests were carried out along the lines.
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Fig. 3. (a) Details of micro-tensile specimen, (b) distributions of tensile specimens and (c) specially designed fixture for tensile tests.
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Fig. 4. EBSD maps of (a) v-face, (b) p-face. (c) Misorientation angle distributions of the
Fig. 5. OM images of (a) v-face, (b) p-face.
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two faces.
Fig. 6. Load vs. depth curves obtained from nanoindentation of (a) v-face and (b)
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p-face.
Fig. 7. Vickers hardness values of v-face and p-face.
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Fig. 8. (a) Strain vs. stress curves of CSed Cu deposits at different angles and (b) variation trends of tensile strength and elongation with loading angles.
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Fig. 9. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of CSed Cu
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deposits. Black boxes indicate the selected area for magnifying. Fig. 10. Schematic diagram of particle impact near interface between spraying passes. Fig. 11. (a) XRD spectra of CSed Cu deposits and (b) Williamson-Hall plot for
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dislocation density.
Fig. 12. (a) Strain vs. stress curves of weave-sprayed Cu deposits at different angles and
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(b) variation trends of tensile strength and elongation with loading angles. Note: the schematic diagram of weave-spraying is at the top middle of Fig. 12a. Fig. 13. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of weave-sprayed Cu deposits. Black boxes indicate the selected area for magnifying.
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Fig. 1. Morphology (a), particle size distribution (b) and cross-sectional view (c) of Cu
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powder.
Fig. 2. The sketch map of v-face, p-face and measurement lines. Note: both nanoindentation and Vickers hardness tests were carried out along the lines.
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Fig. 3. (a) Details of micro-tensile specimen, (b) distributions of tensile specimens and
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(c) specially designed fixture for tensile tests.
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Fig. 4. EBSD maps of (a) v-face, (b) p-face. (c) Misorientation angle distributions of the
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Fig. 5. OM images of (a) v-face, (b) p-face.
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Fig. 6. Load vs. depth curves obtained from nanoindentation of (a) v-face and (b)
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p-face.
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Fig. 7. Vickers hardness values of v-face and p-face.
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Fig. 8. (a) Strain vs. stress curves of CSed Cu deposits at different angles and (b)
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variation trends of tensile strength and elongation with loading angles.
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Fig. 9. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of CSed Cu deposits. Black boxes indicate the selected area for magnifying.
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Fig. 10. Schematic diagram of particle impact near interface between spraying passes.
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dislocation density.
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Fig. 11. (a) XRD spectra of CSed Cu deposits and (b) Williamson-Hall plot for
Fig. 12. (a) Strain vs. stress curves of weave-sprayed Cu deposits at different angles and (b) variation trends of tensile strength and elongation with loading angles. Note: the schematic diagram of weave-spraying is at the top middle of Fig. 12a.
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Fig. 13. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of
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Graphical abstract
ACCEPTED MANUSCRIPT Highlights: Strong microstructural and mechanical anisotropies exist in cold sprayed deposits.
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Mechanism of anisotropic behaviors was clarified.
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Tensile strength contributions of interfaces in all directions were evaluated.
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Novel weave-spraying pattern was proposed to modify in-plane tensile anisotropy.
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1.
Kang Yang, Wenya Li, Xiawei Yang, Yaxin Xu PII: DOI: Reference:
S0257-8972(17)31266-5 https://doi.org/10.1016/j.surfcoat.2017.12.043 SCT 22961
To appear in:
Surface & Coatings Technology
Received date: Revised date: Accepted date:
21 August 2017 18 December 2017 19 December 2017
Please cite this article as: Kang Yang, Wenya Li, Xiawei Yang, Yaxin Xu , Anisotropic response of cold sprayed copper deposits. The address for the corresponding author was captured as affiliation for all authors. Please check if appropriate. Sct(2017), https://doi.org/10.1016/j.surfcoat.2017.12.043
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ACCEPTED MANUSCRIPT Anisotropic response of cold sprayed copper deposits Kang Yang, Wenya Li *, Xiawei Yang, Yaxin Xu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, P.R. China
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* Corresponding: Email: [email protected];
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Tel.: +86-29-88495226; Fax.: +86-29-88492642
Abstract
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Cold spray (CS) is characterized as a solid-state process of high deposition efficiency
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for metallic coatings as well as additive manufacturing of metals. However, due to high velocity impact and extensive deformation of particles during CS, the as-received
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coatings or deposits may present anisotropic characteristics which could influence the
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performance of deposits. Hence this study aims to investigate the anisotropic behaviors of CSed copper deposits in a systematic way. The microstructure and micromechanical
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properties of the deposits both in the cross-section (v-face) and in the parallel plane to
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the surface (p-face) were characterized. Tensile tests were performed at various loading angles with respect to the nozzle moving direction in the p-face. It is revealed that there exist strong microstructural and mechanical anisotropies in CSed deposits. Different interparticle interaction results in more severe particle impact deformation in v-face than p-face, with larger elastic modulus and microhardness values. The tensile tests show an unexpected anisotropy in both ultimate tensile strength and elongation, with the highest performance occurring at the angle of 20°. The in-plane tensile anisotropy could
ACCEPTED MANUSCRIPT be attributed to the parallel multiple passes. Therefore, a novel weave-spraying method was proposed, which can greatly reduce the tensile anisotropy of CSed deposits. Keywords: Anisotropic response; Cold spray; Tensile property; Weave-spraying
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1. Introduction
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Cold spray (CS) has attracted increasing attention as a versatile solid-state coating
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and additive manufacturing technique [1,2]. In this process, substrates are exposed to a high velocity (300-1200m/s) jet of small particles accelerated by an expanding gas
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stream at temperatures lower than the melting point of spray material [1,2]. Beyond a
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critical velocity, which is governed by material properties and process conditions, a mechanical interlocking or metallurgical bonding is produced. Unlike conventional
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thermal spray (TS), these conditions identify CS as a low-temperature and high-speed
most cases [1-3].
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process with significant advantages such as lack of oxidation or phase transformation in
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Without considering assistant processes, the particle impact velocity and temperature
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play key roles in forming effective bonding between particle/particle and particle/substrate, which are governed by process parameters such as type of gas, gas pressure, gas temperature, nozzle geometry and standoff distance [4,5]. Up to now, a lot of investigations based on particles impact behavior and its effect on bonding have been carried out [1,5]. Because of these unique advantages, a variety of materials including most metals, alloys, composites, ceramics and even nanostructured metallic materials have been
ACCEPTED MANUSCRIPT deposited by CS technique [6-10]. To date, a number of papers have been published on mechanical characteristics of CSed coatings [8,11-13]. After CS, a significant increase of hardness can be obtained compared to feedstock powder, which is due to the strong work hardening effect [13]. Researches of CSed copper, aluminum, titanium and their
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alloys [11-13] show that an excellent tensile strength can be obtained but with a very
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low ductility. These testing specimens are usually limited to one direction parallel or
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vertical to spray gun traversing, and then the testing results of such specimens have been used to characterize the mechanical response of CSed deposits [2,8-13]. However,
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this kind of characterization cannot describe the complex mechanical behaviors of CSed
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deposits, which can result in a significant deviation from the actual mechanical properties. Because of the inhomogeneous particle impact deformation behavior and
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existence of particle interfaces (especially between spraying passes), anisotropies of
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microstructure and mechanical properties may be developed. From previous studies, it is known that strong plastic anisotropies exist in many
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solid-state techniques with severe deformation such as cold rolling [14], extrusion [15]
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and friction stir welding/friction stir processing [16,17]. In the study of friction stir welding [16,17], it is demonstrated that the textures developed in the joint include Cube {001}⟨100⟩ , Cube ND {001}⟨110⟩ , S {123}⟨634⟩ , Brass {011}⟨211⟩ and Goss {011}⟨100⟩. And the strong Cube texture generally resulted in the highest yield stress at the loading angle of 45° because of its pronounced and enduring effect on strain localization. However, only a few attempts have been made on the anisotropic response of CSed
ACCEPTED MANUSCRIPT deposits to date [18-20]. Bolelli et al. [18] employed depth-sensing indentation to assess the mechanical properties of CSed Ta, which reported some differences in micromechanical properties (hardness and elastic modulus) of cross-section and surface. Coddet et al. [19] and Seiner et al. [20] investigated the elastic and electrical
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anisotropies of coatings prepared by CS, and the results revealed that only weak
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deviations from perfect isotropy exhibited in CSed deposits. The above authors have
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carried out comprehensive studies on some anisotropic properties of CS. However, the microstructural and mechanical anisotropies caused by impact behaviors have not been
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well understood.
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As for the optimization design and reliable performance evaluations of CSed coatings (e.g. the formability and quality of CSed deposits), it is of great importance to consider
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the anisotropies of microstructure and mechanical properties in CS because of its
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significant effect on stress-strain distribution. Therefore, the main aim of this work is to characterize the anisotropy between two faces and in-plane tensile anisotropy in CSed
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copper deposits. The contributions of particle interfaces on all-directional tensile
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strengths were evaluated. Finally, a weave-spraying pattern was proposed to modify the in-plane tensile anisotropy of CS.
2. Experimental 2.1 Deposits fabrication The powder used in this study is commercially spherical Cu (Beijing You Xing Lian Technology Co., Ltd, China). Figs. 1a and 1c show the morphology of Cu powder and
ACCEPTED MANUSCRIPT cross-sectional view of Cu particles. The particle size distribution characterized by a laser diffraction sizer (MASTERSIZER 2000, Malvern Instruments Ltd., UK) is shown in Fig. 1b with a median size (D50) of 26.2μm. 3mm thick Cu plates were used as substrates. A custom developed CS system was employed to perform deposition. The
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maximum temperature and pressure of the system are 800°C and 1MPa, respectively.
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The nozzle had an expansion ratio of 6.7 and a divergent section length of 200mm. In
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this experiment, helium was used as the accelerating gas with an inlet pressure of 0.8MPa and temperature of around 500°C. The nozzle standoff distance was 25mm and
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gun traverse speed was 20mm/s. Following 8-passes spraying, 8mm thick Cu deposit
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2.2 Materials characterizations
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was obtained.
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The microstructure of the cross-section (v-face) and the plane parallel to the surface (p-face) was examined with electron back-scattered diffraction (EBSD) performed on a
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scanning electron microscope (SEM, JSM5800LV, JEOL, Japan). The v-face and p-face
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are shown in Fig. 2. The microstructure was also examined with an optical microscope (OM, OLYMPUS GX71, Japan). Depth sensing indentation technique was used to determine micromechanical properties using a Nanoindenter G200 (Agilent Technologies Inc, USA). Indentations with a constant load (10mN) were carried out on splats. Eight indentations were tested in each case, and after ignoring the indentations with extreme values (maximum and minimum values) six lines are shown. The Vickers hardness of both faces was measured with a digital microhardness tester (Struers
ACCEPTED MANUSCRIPT Duramin-A300, Denmark) with a load of 200g and dwell time of 15s. Twelve indentations were taken in each face with ignoring the indentations with extreme values (maximum and minimum values). Both nanoindentation and Vickers hardness tests of two faces were carried out along measurement lines as shown in Fig. 2. The in-plane
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tensile properties in p-face were measured at six directions (0°, 15°, 30°, 45°, 60°, 90°)
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with a tensile test machine (INSTRON-3382, America), where 0° is the direction
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parallel to the spraying gun traverse as shown in Fig. 3b. Three samples were tested to evaluate the tensile properties for each direction. It should be noted that standard
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deviation was applied to get error bars of the Vickers hardness and tensile properties. A
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micro-tensile specimen with 2mm thickness was used [21], where the width in the gauge section and total length were 2mm and 25mm, respectively (Fig. 3a). A specially
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designed fixture (Fig. 3c) was used to ensure correct application of the load and prevent
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secondary bending or other unwanted movement of the specimens. The fracture surfaces of tensile specimens were also observed with SEM. An X-Ray diffractometer system
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(XRD, Siemens D500, Germany) was used to measure microstrain and crystallite size
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with 40kV and 30mA of Cu Kα radiation at a wavelength of 0.1542nm. The scan was conducted in 2θ mode and across a range of 10° to 120° at a step resolution of 0.033°.
3. Results and discussion 3.1 Anisotropy between p-face and v-face 3.1.1
Microstructure
The EBSD maps of v-face and p-face are shown in Figs. 4a and 4b, respectively. It
ACCEPTED MANUSCRIPT can be seen that both particle shapes and grain size distributions are different between v-face and p-face. Particle plastic behavior such as impact and extrusion are more concentrated on v-face than p-face. As a consequence, particle deformation in v-face is more severe to form flat shapes with disc shapes in p-face. This phenomenon can also
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be found in OM images as shown in Fig. 5. Because of the anisotropic impact, grain
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distribution characteristics in the two cases are clearly different. In comparison to the
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p-face, grain distribution in v-face is inhomogeneous, with more refined grains near the particle impact interface. This is due to the recrystallization which has been confirmed
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by previous studies [8,22]. The statistical result of misorientation angle distributions
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(Fig. 4c) shows little difference of grain boundaries fractions between two faces (including High-Angle Grain Boundaries (HAGBs, θ≥15°) and Low-Angle Grain
3.1.2
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Boundaries (LAGBs, 15°>θ≥2°)).
Nanoindentation and Vickers hardness
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The anisotropic particle deformation of v-face and p-face was quantified with depth
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sensing indentation as shown in Figs. 6a and 6b, respectively. Compared to p-face, micromechanical properties of v-face show a remarkable increase of 13.5% for modulus and 11.8% for microhardness. This can be the consequence of the more severe plastic deformation in v-face than that in p-face. The nanoindentation results are consistent with the studies on CSed Ta by Bolelli et al. [18], having some differences (hardness and elastic modulus) between cross-section and surface. But the differences between two faces in our study are greater compared to the above mentioned study [18], which
ACCEPTED MANUSCRIPT could be mainly due to different materials. Fig. 7 shows the Vickers hardness of v-face and p-face. It should be noted that the average microhardness value of pure Cu was measured as 70HV. Both the microhardness values of v-face and p-face are higher than pure Cu, which has been
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confirmed due to the strong work hardening effect during CS [8,11]. It can be seen that
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the Vickers hardness values of two faces are much closer than that of nanoindentation.
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The difference of two test results could be due to the different measuring principles and size effect [23,24]. Compared to nanoindentation, the scope of Vickers hardness test is
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relatively much large, and the Vickers hardness is strongly affected by the material
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porosity and adhesion strength between particles [23].
In-plane tensile anisotropy
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3.2.1
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3.2 In-plane tensile anisotropy and strength contributions
Being a significant mechanical performance, tensile properties in six directions were
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also measured as shown in Figs. 8a and 8b. It can be observed that the stress-strain
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curves of CSed Cu deposits change with loading angles, indicating a strong anisotropy in both tensile strength and elongation. With the increase of loading angle, both the ultimate tensile strength and elongation first rise and then drop with the highest values measured at angles of 15°-30°. For angles over 30°, the tensile performance deteriorates rapidly, reaching a minimum at 90°. Fig. 9 shows typical fracture surfaces of 0°, 30° and 90° tensile samples. No dimples can be found in these fracture surfaces. The fracture takes place along interfaces
ACCEPTED MANUSCRIPT between deposited particles rather than through them, indicating a brittle rupture for the CSed Cu. Consistent with tensile strengths, the morphology of 30° tensile sample is the coarsest (Figs. 9b and 9e), while that of 90° the smoothest (Figs. 9c and 9f). The interlocking zones between Cu particles are marked in the magnified views (Figs. 9d-9f).
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It can be found that the amount of interlocking zones of 30° tensile sample is the most,
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with the least of 90°.
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For the traditional spraying pattern, the spraying passes are usually parallel. The parallel spraying interfaces between passes easily result in anisotropy of CSed coatings.
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It is considered that the different tensile properties in different directions are attributed
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to the following factors: (1) Interparticle interlocking effect between spraying passes has a positive effect; (2) During CS deposition, the coating is formed through bonding
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between impact particles, while bonding quality is dependent on particle deformation
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behavior, which is determined by the surface state of the previously deposited coating. Near the interface between spraying passes as shown in Fig. 10, the normal velocity
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component (𝑉𝑛 ) is lower than the actual velocity (𝑉𝑎 ). The decrease in normal velocity
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between spraying passes indicates a weaker interfacial bonding than that within a pass, which could affect the tensile property in a negative manner. To isolate these two interfacial bonding strengths experimentally, a technique of multiple load indentation and the hardness loss parameter proposed by Goldbaum et al. [23] could be helpful. As a consequence, for angles between 0° and 30°, the positive factor “interlocking effect” is dominant. Hence the tensile properties including ultimate tensile strength and elongation increase with angles. This association reverses for larger angles between 30°
ACCEPTED MANUSCRIPT and 90°, and decreases the tensile property. It should be noted that oxides and pores exist at the interfaces between particles in CSed deposits. In addition, although bonding can happen at the interfaces, the strength is relatively low compared to the particles themselves. Therefore, the interfaces between particles are weak, which lead to the poor
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performance of CSed deposits compared to the bulk material.
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In summary, unlike the anisotropies existed in solid-state techniques like cold rolling
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[14], extrusion [15] and friction stir welding/processing [16,17], which is resulted from grain orientation, the in-plane tensile anisotropy of CS is mainly due to interparticle
3.2.2
Strength contributions
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interfaces especially between spraying passes.
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Because of the main responsibility of interfaces to tensile anisotropy, it is
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indispensable to study and quantify the changes of their contributions to the final ultimate tensile strength. Several additional laws are available for estimating the total
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strength [25,26]. In our case, the final tensile strength (𝜎𝑓−𝐴 ) depends on several
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parameters, which can be associated linearly [25,26]: 𝜎𝑓−𝐴 = 𝜎0 + 𝜎𝐻𝑃 + 𝜎𝑊𝐻 − 𝜎𝑃𝐼−𝐴
(1)
where 𝜎0 is the base strength of pure Cu (128MPa); 𝜎𝐻𝑃 is the grain size strengthening according to the Hall-Petch law; 𝜎𝑊𝐻 is the work hardening contribution; 𝜎𝑃𝐼−𝐴 is the special existence in CS, which means a negative contribution to the final tensile strength due to interfaces between particles [8]. According to literature [27], the strengthening due to grain refinement is described by
ACCEPTED MANUSCRIPT the Hall-Petch relation: 𝜎𝐻𝑃 = 𝑘𝑦 ⁄√𝐷
(2)
where 𝑘𝑦 is a constant depending on the material (37.08MPa√μm for Cu); D is the average grain size after impact, which can be measured with EBSD (1.38μm).
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Accordingly, considering the calculation from Eq. (2), the increment resulted from grain
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refinement is estimated to be 32MPa.
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The general equation to describe the effect of work hardening (dislocation) on the strength is given by [26,28]:
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𝜎𝑊𝐻 = αMGb√𝜌
(3)
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where α is a constant equal to 0.33; M is the Taylor factor (3.06) G is the shear modulus for Cu (3.9 × 104 MPa); b is the Burgers vector (0.3615 × 10−3 μm); 𝜌 is
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the dislocation density, which can be estimated by XRD as shown in Fig. 11, and it can
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be calculated by the Williamson-Hall equation [29]: (4𝑎)
ρ = 14.4𝜀 2 ⁄𝑏 2
(4𝑏)
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𝛽𝑐𝑜𝑠𝜃⁄λ = 0.9⁄𝐷 + 2𝜀sinθ⁄λ
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where 𝛽 is the peak width at half maximum height; 𝜆 is the wave length of X-ray beam (0.1542nm for Cu Kα); 𝐷 is the crystallite size; 𝜀 is the lattice strain; 𝜃 is the Bragg angle. By plotting 𝛽𝑐𝑜𝑠𝜃⁄λ against s𝑖𝑛𝜃, the slope of the graph gives 2𝜀 ⁄λ as shown in Fig. 11b, the Y-intercept gives 0.9⁄𝐷 and ρ is calculated using Eq. (4b) with a value of 4.78 × 1014 m-2. Hence, the increment contributed from dislocation strengthening is 311MPa. Therefore, according to Eq. (1), the particle interfaces contribution 𝜎𝑃𝐼−𝐴 (MPa) of
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(5)
In order to get all-angled particle interfaces effects, the 𝜎𝑃𝐼−𝐴 is fitted with Gaussian function (𝐴𝑛𝑔 is the tensile direction, unit: °): 2
𝜎𝑃𝐼−𝐴 = 440 − 60e−2((𝐴𝑛𝑔−19)⁄37)
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(6)
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The fitted function suggests that the lowest 𝜎𝑃𝐼−𝐴 occurs at angle of 20° with the
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highest final ultimate tensile strength value of about 90MPa. It can be concluded here that the negative contribution of the weak existence of interparticle interfaces (𝜎𝑃𝐼 ) is
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the main factor in CS, which hampers the improvement of tensile property. Therefore,
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weakening the existence of interparticle interfaces could further improve the tensile
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3.3 Weave-spraying
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performance in all directions.
The large tensile anisotropy in CSed deposits is not desirable, which is the result of
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multiple parallel spraying passes. Therefore, in order to reduce or eliminate the parallel
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multi-pass effect, an alternative spraying pattern was proposed: weave-spraying as shown in Fig. 12a. The tensile performance achieved with different angles of weave-sprayed Cu is shown in Fig. 12. Obviously, the novel pattern has a great suppression on the tensile anisotropy of CSed deposits. Although the high tensile values for some angles (~15°-30°) are sacrificed, an almost isotropic CSed deposit is produced by weave-spraying, producing a homogeneous material in both tensile strength and elongation. Figs. 13a-13c show the typical fracture surfaces of 0°, 30° and 90° tensile
ACCEPTED MANUSCRIPT samples, respectively, with the respective magnified views showed in Figs. 13d-13f. As expected, the fracture surfaces of the three loading angles show similar morphologies. Weave-spraying well modifies the in-plane tensile anisotropy of CS. For weave-spraying pattern, the spraying passes are orthogonal, which can effectively
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reduce the difference between different directions. In the case of a certain thickness of a
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CSed deposit, it is believed that the multi-pass effect would be further weakened with
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the increase of weave-spraying times. For TS, the effect of spray trajectory has been well investigated in many studies which show that spray trajectory can influence the
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heat and mass transfer significantly, and therefore on the coating microstructure,
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physical and chemical properties and their quality [30-32]. While for CS, only a few investigations on the nozzle trajectory have been reported so far [33,34]. Chen et al. [34]
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proposed a spiral trajectory applicable to CS for repairing damaged workpiece, which
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can improve productivity and reduce cost in machinery work. But they did not study the mechanical properties or anisotropy. In our opinion, this pattern (spiral-trajectory
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spraying) could also be a method to weaken the in-plane tensile anisotropy, because this
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spraying pattern can effectively reduce the difference between different directions. Therefore, weakening the parallel multi-pass effect between different directions is a guiding idea to eliminate the coating anisotropy.
4. Conclusions In this paper, Cu was deposited to investigate the anisotropic behaviors of CS. The existence of strong anisotropies was found in CSed deposits including microstructure
ACCEPTED MANUSCRIPT and mechanical properties, and the responsible mechanisms have been discussed. The negative contribution of the existence of weak interparticle interfaces has also been evaluated. The main conclusions can be drawn as follows: (1) Particle impact deformation and grain size distribution behaviors are different
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between the v-face and p-face, with little difference in misorientation angle
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distributions.
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(2) Nanoindentation tests show that the anisotropic impact deformation results in larger elastic modulus and hardness in the v-face than p-face.
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(3) Strong anisotropy exists in the in-plane tensile properties of CSed deposits. The
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parallel multiple passes have the main responsibility to the in-plane tensile anisotropy. The fitted function suggests that the highest tensile performance
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occurs at the angle of 20°, and the lowest at 90°.
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(4) A novel weave-spraying method was proposed, which can effectively reduce the parallel multi-pass effect between different directions. The weave-spraying has a
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great modification effect on the tensile anisotropy of CSed deposits.
Acknowledgments The authors would like to thank for the financial support from the National Key Research and Development Program of China (2016YFB1100104), the fund of SAST (SAST2016043) and the 111 Project (B08040).
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ACCEPTED MANUSCRIPT Figure captions Fig. 1. Morphology (a), particle size distribution (b) and cross-sectional view (c) of Cu powder. Fig. 2. The sketch map of v-face, p-face and measurement lines. Note: both nanoindentation and Vickers hardness tests were carried out along the lines.
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Fig. 3. (a) Details of micro-tensile specimen, (b) distributions of tensile specimens and (c) specially designed fixture for tensile tests.
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Fig. 4. EBSD maps of (a) v-face, (b) p-face. (c) Misorientation angle distributions of the
Fig. 5. OM images of (a) v-face, (b) p-face.
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Fig. 6. Load vs. depth curves obtained from nanoindentation of (a) v-face and (b)
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Fig. 7. Vickers hardness values of v-face and p-face.
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Fig. 8. (a) Strain vs. stress curves of CSed Cu deposits at different angles and (b) variation trends of tensile strength and elongation with loading angles.
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Fig. 9. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of CSed Cu
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deposits. Black boxes indicate the selected area for magnifying. Fig. 10. Schematic diagram of particle impact near interface between spraying passes. Fig. 11. (a) XRD spectra of CSed Cu deposits and (b) Williamson-Hall plot for
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dislocation density.
Fig. 12. (a) Strain vs. stress curves of weave-sprayed Cu deposits at different angles and
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(b) variation trends of tensile strength and elongation with loading angles. Note: the schematic diagram of weave-spraying is at the top middle of Fig. 12a. Fig. 13. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of weave-sprayed Cu deposits. Black boxes indicate the selected area for magnifying.
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Fig. 1. Morphology (a), particle size distribution (b) and cross-sectional view (c) of Cu
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Fig. 2. The sketch map of v-face, p-face and measurement lines. Note: both nanoindentation and Vickers hardness tests were carried out along the lines.
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Fig. 5. OM images of (a) v-face, (b) p-face.
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Fig. 7. Vickers hardness values of v-face and p-face.
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Fig. 8. (a) Strain vs. stress curves of CSed Cu deposits at different angles and (b)
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Fig. 9. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of CSed Cu deposits. Black boxes indicate the selected area for magnifying.
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Fig. 10. Schematic diagram of particle impact near interface between spraying passes.
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Fig. 11. (a) XRD spectra of CSed Cu deposits and (b) Williamson-Hall plot for
Fig. 12. (a) Strain vs. stress curves of weave-sprayed Cu deposits at different angles and (b) variation trends of tensile strength and elongation with loading angles. Note: the schematic diagram of weave-spraying is at the top middle of Fig. 12a.
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Fig. 13. Fracture surfaces of (a, d) 0º, (b, e) 30º, (c, f) 90º tensile samples of
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Graphical abstract
ACCEPTED MANUSCRIPT Highlights: Strong microstructural and mechanical anisotropies exist in cold sprayed deposits.
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Mechanism of anisotropic behaviors was clarified.
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Tensile strength contributions of interfaces in all directions were evaluated.
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Novel weave-spraying pattern was proposed to modify in-plane tensile anisotropy.
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1.