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Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits
Journal Pre-proof Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits Yujuan Li, Yingkang Wei, Xiaotao Luo, Changjiu Li, Ninshu Ma
PII:
S1005-0302(19)30410-4
DOI:
https://doi.org/10.1016/j.jmst.2019.09.023
Reference:
JMST 1807
To appear in:
Journal of Materials Science & Technology
Received Date:
14 July 2019
Revised Date:
1 September 2019
Accepted Date:
23 September 2019
Please cite this article as: Li Y, Wei Y, Luo X, Li C, Ma N, Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits, Journal of Materials Science and Technology (2019), doi: https://doi.org/10.1016/j.jmst.2019.09.023
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Research Article Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits
Yujuan Lia, Yingkang Weia, Xiaotao Luoa,*, Changjiu Lia,*, Ninshu Mab
a
State Key Laboratory for Mechanical Behavior of Materials, School of Materials
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Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China b
Joining and Welding Research Institute, Osaka University, Osaka 567-0047, Japan
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*Corresponding author.
E-mail addresses: [email protected] (X.T. Luo), [email protected]
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(C.J. Li)
September 2019]
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[Received 14 July 2019; Received in revised form 1 September 2019; Accepted 23
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Inter-particle bonding is an important factor affecting the property of cold sprayed metallic deposit. Because the interface bonding between particles in deposit is directly
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determined by plastic strain of particles during spraying, Cu deposits were made at series of impact velocities of 578 m s-1 to 745 m s-1 and 807 m s-1 to correlate particle
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impact condition with microstructure and properties of the deposits. Results show that as the average particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1, the deposition efficiency of feedstock powder increases from 58% to 84% and even to 95%. Although all three deposits reveal dense microstructure due to the high ductility of Cu, the deformation degree of the deposited particles remarkably increases with increasing impact velocity. The enhanced plastic deformation of the deposited particles leads to more dispersed oxide scale and thereby stronger inter-particle 1
bonding with the strength of the deposit along the deposition direction increasing from 25.8 MPa to 148.5 MPa. The electrical and thermal conductivities at throughthickness direction of the deposit at particle impact velocity of 807 m s-1 are 78 % IACS, 295 W m-1 K-1, respectively.
Keywords: Cold spray, Particle impact velocity, Deposition efficiency, Electrical
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conductivity, Thermal conductivity, Mechanical property
1. Introduction
Cold spray was being well known as a surface coating technology in the past
decades, and recently it has attracted wide attention because it can also be a promising
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additive manufacturing (AM) process [1-3]. In cold spray, micron-scale particles
(typically 5-50 μm) are fed into supersonic-speed gas flow and are accelerated to 300-
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1200 m s-1, and then impact on the substrate at fully solid state [4-6]. The particle bonding is formed by high velocity impact induced intensive plastic deformation of
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the particle with a deposition temperature well below the melting point of a depositing material [7, 8]. As compared with melting and re-solidification based coating and AM
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processes such as thermal spray, laser cladding and selective laser melting, the low depositing temperature is one of the most remarkable characteristics of cold spray [911]. This makes depositing metallic materials without oxidation realized in open air
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and heating effect to the substrate avoided. Cu [12, 13], Al [14, 15], Ni [16], Ti [17], Cu-based alloys [18], Al-based alloys [19, 20], Ni-based alloys [21], Fe-based alloys
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[8]) and metal-based composites such as Al-Al2O3 [22] and Al-SiC [23] in the forms of coating or parts have been successfully achieved by cold spray so far.
Different from the metallurgical bulk counterparts, cold sprayed deposits are formed through the accumulated deposition of feedstock particles with a diameter of 5-50 μm in completely solid state. Besides the intrinsic characteristic of the depositing material, the inter-particle bonding is a decisive factor influencing penetration of the 2
corrosive media and the transfer of load and electron from one particle to another, and therefore determining the deposit properties, especially strength [13], thermal and electrical conductivities [24], and corrosion resistance [25, 26]. However, the interparticle bonding in cold sprayed deposit may be poor, and pores may appear between particles as the plastic deformation degree of particles is relatively small. Poor interparticle bonding and pores between particles in deposit lead to high electrical resistance, high thermal resistance and low corrosion resistance, thereby resulting in poor performances. In our previous results [24], Cu deposit was prepared by nitrogen
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gas cold spray at a gas pressure 3 MPa and a gas temperature of 400 oC and found that the through-thickness thermal conductivity of the deposit were only 207 W m-1 K-1 (52% of that for annealed bulk Cu). Wei et al. [19] prepared Al6061 deposit on
magnesium alloy plates at 2.5 MPa and 300 oC, and reported that the corrosion
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behavior of the deposit was close to that of the substrate and its corrosion resistance
was far below that of Al6061 bulk. Moreover, many studies [27, 28] have shown that
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the tensile strength of cold sprayed deposit is only 1/3-1/4 of that for corresponding bulk. By further observing the fracture morphology of deposit，it is found that the
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between particles [29].
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crack mainly propagates along the inter-particle interface, indicating the poor bonding
The poor inter-particle bonding can be caused by the oxide scale distributed at the interface of the deposited particle in cold sprayed deposit. As we all known,
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nanometer thick oxide scale is inevitably formed for even all the commercial metallic powders during production and storage. As long as the metallic surface is exposed to
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air, oxide scale generates with a thickness in ranges of 0.3-14.6 nm for Ni [30, 31], 19-134 nm for Fe [32] and 3-76.9 nm for Cu [33]. During the high-velocity particle impact in cold spray, the particle undergoes plastic deformation. As the strain at the interface of the deposited particle reaches a critical value, the oxide scale on the surface of feedstock particle fractures into discontinuous segments, allowing the fresh metallic surface contact and thereby form bonding [34]. Greater interfacial strain will lead to more dispersed oxide scale and more fresh metallic contact, which benefit the 3
bonding formation at the interface and thereby improve the bonding [35, 36].
To improve inter-particle bonding of cold sprayed deposit and further improve deposit performance, it is necessary to increase the degree of particle plastic deformation during particle impact process. In the past decades, many studies have been done to increase the degree of particle plastic deformation during cold spray. Because the degree of particle plastic deformation is directly related to particle impact velocity, the most common method is to improve accelerating behavior of particle. Li
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et al. [37, 38] optimized the nozzle geometry to improve the accelerating behavior of powder particles, and further expect to improve inter-particle bonding and deposit’s performance while the time cost of deposit preparation is increased due to the long cycle of optimizing nozzle geometry. Because of the excellent acceleration
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performance of He gas for powder particle, Huang et al. [27] and Gärtner et al. [29] used He gas as accelerating gas during cold spray, and the obtained deposit showed
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extremely high strength, indicating inter-particle bonding was obviously improved. Above results show that a high particle imapct velocity obtained by adjusting spraying
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parameters is benefit for the inter-particle bonding and thereby the microstructure and propertis of cold sprayed deposit. However, the quantitative relationship between
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particle impact condition and the plastic deformation degree of deposited particle, and the influence of particle impact condition on the interfacial microstructure and property of the deposited particle in corresponding deposit which affect the whole
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property of deposit significantly are not clear.
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In this work, the cold sprayed deposits were prepared at different particle impact
velocities by using different spraying parameters. The particle impact velocity at corresponding spraying parameter was quantitatively characterized. The microstructure and properties, including mechanical property and conductivities, of the deposits prepared at different particle impact velocities were investigated. Meanwhile, the interfacial microstructure of the deposited particle in deposit was also observed. 4
2. Experimental 2.1. Feedstock powder A commercial available water-atomized Cu powder was used as the feedstock powder. The morphology and particle size distribution of Cu powder are shown in Fig. 1. Particle size distribution of the Cu powder was measured by a laser diffractometry (Laser Mastersizer, Malvern Instruments, Malvern, UK) as shown in Fig. 1(b), indicating that the average particle sizes (d50) of the Cu powder was 28
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μm. An oxygen content of 0.044 wt.% was detected for Cu feedstock powder using oxygen/nitrogen analyzer (TC-600, LECO, USA). Commercially pure Cu plates were employed as the substrates. Prior to the coating deposition, the substrates were grit-
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blasted by alumina grits.
2.2. Deposit preparation
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A commercial cold spray system (PCS-1000, Plasma Giken Co., Ltd.) was used for deposit preparation. In this system, a Laval nozzle was used with throat diameter
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of 2 mm, outlet diameter of 4.4 mm, and divergent section length of 170 mm. N2 gas was used as the processing gas. To investigate the influence of the particle impact
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velocity on the inter-particle bonding, the Cu deposits were deposited at three different spraying conditions as listed in Table 1 (named as low, medium and high conditions according to the gas pressure and temperature values). The particle
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velocities of feedstock powders at different spraying parameters were measured by a commercial thermal spray in-flight particle diagnostic system (DPV 2000, Tecnar
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Automation Ltd., St-Bruno, Qc, Canada) equipped with a laser system (CPS-2000). The particle velocity measurement was made at the centerline of the gas flow, 20 mm away from the nozzle outlet which was same as the standoff distance during deposit preparation. For each spraying condition, more than 500 particles were measured.
2.3. Deposit characterization Microstructures of deposit were characterized by scanning electron microscopy 5
(SEM, MIRA 3 LMH, TESCAN, Czech Republic) with an accelerating voltage of 15 kV. Porosity of deposits was measured by image analyzing method from cross section (10 SEM images at 1000× for each specimen). Transmission electron microscopy (TEM) samples were made from the etched cross section using focus ion beam (FIB, JIB-4500 Multi-beam System, JEOL) to examine the interfacial microstructure and track the oxide scale (TEM, JEM-F200, JEOL). Conductivity measurements and uniaxial tensile test were performed on the deposit to quantitatively evaluate the effect of inter-particle bonding quality on properties because there are more inter-particle
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boundaries in unit length along the through-thickness direction. Through-thickness electrical conductivity of deposit was calculated according to the resistance measured by using double bridge method (QJ57, China) with an accuracy of 10-5 Ω and the
samples used were free-standing deposit with dimensions of 40 mm × 1 mm × 1 mm.
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Through-thickness thermal conductivity of deposit was measured by using a laser flash technique (FLASHLINETM SYSTEM 3050, Anter Corporation, US) and the
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sample was also free-standing deposit with a dimension of 𝜙mm×3 mm. The plate tensile specimens with a thickness of 1 mm were sampled along the spray
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direction (Fig. 2). Tensile test was performed by using a material mechanical property test machine (Instron 5848, Micro Tester) according to EN 10002-1 standard. Detailed
3. Results
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dimensions of the plate tensile samples are also shown in Fig. 2.
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3.1. Particle impact velocity
To evaluate the particle impact velocities at different spraying conditions, the
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velocity of the particle was measured by using the DPV 2000 system at the center of flow under the conditions of preparing deposit as listed in Table 1. The measured particle velocity distributions of the particles at 2.5 MPa 500 oC, 3 MPa 800 oC and 5 MPa 800 oC are shown in Fig. 3(a, b and c), respectively. At those spraying conditions, all velocities of the particles range from 200 m s-1 to 1000 m s-1. As expected the average velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1 as the spraying condition varies from 2.5 MPa 500 oC to 3 MPa 800 oC and 5 MPa 6
800 oC.
3.2. Deposition efficiency of feedstock powder Deposition efficiency is one of the most important concerns for thermal/cold spraying of deposits since it partially determines the cost of the deposits. Powder deposition efficiency (DE) is the mass ratio of the prepared deposit to the powder sprayed on the substrate during the spraying process. The mass of deposit is the mass difference of the substrate before and after spraying. To characterize the mass
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difference of substrate, an electrical balance (Fisher Scientific, A-2000DS) with accuracy of 0.1 mg was used to measure the mass of the substrate before (Wsa) and after (Wsb) spraying. The mass of the powder consumed on the substrate can be
expressed by the product of powder feeding rate and the spraying time spent on the
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substrate during spraying. Powder feeding rate is the mass of powder consumed per
unit time. To calculate the powder feeding rate, the mass of the powder in the powder
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feeder before (Wpa) and after (Wpb) spraying was measured by using an electrical balance, and the total spraying time (T) was recorded by a stopwatch. To estimate the
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spraying time consumed on the substrate, the length of substrate (L) was measured, and the transverse movement velocity of nozzle (V), the pass number of spraying (N)
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and the stroke number in one pass spraying (n) were also recorded. The detailed expression to calculate the deposition efficiency of sprayed powder can be seen in Eq. (1):
𝑊sb −𝑊sa 𝑊pa−𝑊 pb 𝐿 × ×𝑁×𝑛 𝑇 𝑉
× 100%
(1)
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DE =
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The deposition efficiency of the Cu powder was measured for more than 3 times at each spraying condition and its average value was taken.
The deposition efficiency of the copper powders as a function of the particle
velocity is shown in Fig. 4. The deposition efficiency increases from 58% to 84% as the average particle velocity increases from 578 m s-1 to 745 m s-1. It is worth noting that a high deposition efficiency of 95% was achieved as the average particle velocity 7
increases to 807 m s-1. As the average particle impact velocity increases obviously, the number of the particles whose velocity is greater than the critical velocity required for particle deposition increase. Therefore, the deposition efficiency of powder particle significantly increases.
3.3. Microstructure of deposits To study the influence of particle impact velocity on microstructure of the cold sprayed Cu, the polished cross sections of the deposits were observed by SEM and are
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shown in Fig. 5(a, c and e). To track the inter-particle boundaries, the polished cross sections of deposit samples were also etched, where inter-particle boundaries can be easily recognized as shown in Fig. 5(b, d and f). Although all three deposits show very dense microstructures with apparent porosity values lower than 0.6%, the
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deformation degrees of particles at different spraying parameters are totally different since different particle impact velocities provide different driving forces for the
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impact induced plastic deformation. As shown in Fig. 5(b, d and f), the deposited particles in the deposits prepared at higher particle velocities are much highly
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flattened. To quantitatively evaluate the difference, flattening ratios were estimated by the image analyzing. The flattening ratio (R) of a particle is defined as: 𝐷
p
(2)
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𝑅 = 𝑑 × 100%
where D is the maximum dimension of the deposited particles perpendicular to the
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depositing direction, and dp is the diameter of the raw particles. The average particle size of the raw Cu particles (28 μm) was taken as dp. D was measured from the SEM
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images of etched cross-section of the deposits. For each deposit, the number of particles counted was more than 200. This results in flattening ratios of 1.2±0.38, 1.5 ±0.41 and 1.9±0.27 for three conditions, respectively.
For cold sprayed deposit, as it is composed of many splats, its performance is highly determined by the inter-particle bonding, including both the number of the inter-particle boundaries in unit length and bonding quality. Either less boundaries or 8
higher bonding quality favors the transfer of current, heat and load across the particles and thereby higher conductivities, and strength and elastic modulus. The interface quantity in a deposit can be characterized by using inter-particle boundary density which is defined as number of the inter-particle boundaries in unit length along a fixed direction. In this work, the inter-particle boundary density is derived from the etched cross-sections of the deposits. To calculate the interface density (dinter) along through-thickness direction, straight lines with distance of 30 m along the depositing direction in the 3000 SEM images of the etched cross-sections of the deposit as
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shown in Fig. 6. The number of interface (N) along each line was counted. dinter can be calculated by the ratio of N to the length of the vertical line as the relation shown in Eq. (3).
N L
(Eq.3)
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dint er
For each deposit, at least 10 measurements were made and the average value was used
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for comparison.
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Fig. 7 shows the through-thickness inter-particle densities of deposits produced at different particle impact velocities. It can be seen that the through-thickness
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interface density increases with the increase of particle impact velocity due to the enhanced plastic deformation. The statistical result indicates interface density in the
m s-1.
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deposit sprayed at 807 m s-1 is 2 times higher than that in the deposit prepared at 578
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Quality of inter-particle bonding is another important concern which determines
the properties of the cold sprayed metals. According to our previous studies [35], the oxide scale arising from the starting particle surfaces will be trapped at inter-particle boundaries, and hinder the metallurgical bonding formation between deposited particles. Therefore, the oxide scale along the inter-particle boundaries was tracked in this work to assess the quality of the inter-particle bonding formed at different particle impact velocities. It is known that during the high velocity particle impact, the ductile 9
particle will be severely plastically deformed. However, the nanometer-thick oxide scale attached on the particle surface would be segmented rather than be deformed due to its brittle nature. Between the oxide-scale segments, clean metallic surface contact and thereby metallurgical bonding can be realized. However, due to the small thickness of oxide scale on the initial surface of powder particles (less than 100 nm), it is difficult to observe the oxide scale at the boundary between particles after deposition by SEM [34, 36]. The fraction of newly created metallic surface is determined by the plastic strain of the interface. It is known that the plastic
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deformation of particle is heterogeneous. The edge area of the interface always undergoes much higher plastic strain than the center area. This means a much clearer difference in dispersion degree of the oxide scale can be observed from the TEM
images taken at the edge area of the interface for deposits made at different particle
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impact velocities. However, in this work, all the properties including strength,
electrical conductivity and thermal conductivity were measured along the depositing
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direction. It is known that the properties along the depositing direction are influenced more by the oxide scale vertical to the depositing direction. Therefore, in this work,
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TEM images were only taken at the center area of the particle interface. TEM samples were made by focus ion beam (FIB) at the center of the interface in the Cu deposits
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prepared at 578 m s-1 (Fig. 8(a)) and 807 m s-1 (Fig. 8(e)). TEM images and the EDS mapping results are also presented in Fig. 8. The continuously-distributed oxide in a white contrast was observed at the center of the boundary in the specimen sprayed at
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578 m s-1 (Fig. 8(b)), while the oxide scale in the white contrast became discontinuous at the boundary in the specimen produced at 807 m s-1 (Fig. 8(f)).
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TEM-EDS mapping results confirm that long oxide scale segments remained at the center of the inter-particle boundary (Fig. 8(d) and (h)). The more dispersed oxide scale at the bottom of the bonded particle resulted from a higher particle impact velocity will benefit for the bonding formation at inter-particle boundary and the corresponding deposit may show higher properties, compared with that sprayed at a lower particle impact velocity. The concentrated oxide scale at the bottom of the bonded particle may also contribute to the anisotropy of cold sprayed metallic 10
deposits since there are more oxide scale layers in unit length along the depositing direction. Yin et al. [39] and Yang et al. [40] reported that the properties anisotropy of cold sprayed deposit was caused by the bonding difference of the inter-particle interface, which is consistent with our results.
3.4. Elastic modulus and tensile strength of the deposits In cold/thermally sprayed metallic deposits, it is known that there are more particle interfaces in unit length along the through-thickness direction than those
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along the in-plain direction. It is rational to infer that the properties along the throughthickness direction much highly depend on the particle bonding. Therefore, in this work, a uniaxial tensile test was performed on the deposits along the through-
thickness direction. Fig. 9 presents the stress-strain curves of the samples produced at
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different particle impact conditions. It can be seen that Cu deposit sprayed at the
highest particle velocity reveals 6 times higher fracture strength than that deposited at
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lowest velocity, suggesting much stronger inter-particle bonding. On the other hand, higher slope of the stress-strain curve was detected with increasing particle impact
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velocities in the elastic deformation range, suggesting higher elastic modulus. Table 2 lists the tensile strength and elastic modulus of all samples. Both strength and elastic
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modulus greatly increase as the particle impact velocity increases and the fracture strength of the deposit prepared at the highest particle velocity is very close to that of the bulk Cu. However, the elastic modulus values of all Cu deposits are lower than
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that of the bulk counterpart (124 GPa), which are 40.4%, 55.8% and 72.9% for the deposit sprayed at 578 m s-1, 745 m s-1 and 807 m s-1, respectively. This fact suggests
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that although high strength can be achieved by depositing at very high spraying parameter resulted in extremely high particle impact velocity (807 m s-1 for Cu in this work), it is hard to achieve the dense Cu deposit with fully bonded particles.
Fig. 10 shows the typical fracture morphologies of the deposits produced at different particle impact conditions. It can be seen from Fig. 10(a, c and e) that the fracture occurs mainly at the inter-particle boundaries, indicating a mainly brittle 11
fracture, which is consistent with the small elongation (<0.2%). As the particle impact velocity increases, the particles are much highly deformed and therefore profiles of the individual particles become more difficult to be recognized. For the deposit prepared at 578 m s-1, no dimple is observed in the fracture morphology and the fracture takes place along the interfaces between deposited particles as shown in Fig. 10(b). However, the fracture morphologies of deposits sprayed at 745 m s-1 and 807 m s-1 show a small amount of dimples, although the fracture mainly occurs along the inter-particle interfaces in Fig. 10(d) and (f). Moreover, the deposited particle in the
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deposit sprayed at 807 m s-1 is more tightly bonded together comparing with that in the deposit produced at 745 m s-1. Therefore, a higher particle impact velocity leads to a better inter-particle bonding and thereby a high tensile strength in Fig. 9. It is worth noting that the ductility of the as-sprayed deposits is very limited as compared with
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the bulk counterpart due to the un-perfect inter-particle bonding and severe work
hardening of the deposited particles during deposition. This drawback could greatly
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hinder the application of cold spray for additively manufacturing parts. As reported in literatures [41, 42], post-spray heat treatment is an effective way to improve the
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especially dislocations.
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ductility by enhancing the inter-particle bonding and eliminating the crystal defects
3.5. Electrical and thermal conductivities of deposits It is known that Cu and its alloys are frequently used as conducting parts due to
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its high intrinsic thermal conductivity and electrical conductivity. Meanwhile, electron transfer is very sensitive to various boundaries due to their diffraction to the electrons.
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Therefore, in the present work, thermal conductivity and electrical conductivity of the deposits sprayed at varies particle impact velocities were experimentally examined. Since the inter-particle boundary density is higher in through-thickness direction, the specimens for thermal and electrical conductivities were also sampled along the depositing direction. Fig. 11 shows thermal conductivity and electrical conductivity of the cold sprayed Cu as a function of the particle impact velocity. Both conductivities increase with increasing particle velocity. The thermal conductivity of the specimen 12
produced from the deposit sprayed at 578 m s-1 is only 204 W m-1 K-1 (51% of that for annealed bulk Cu), and then it increases to 260 W m-1 K-1 for the specimen produced from the deposit sprayed at 745 m s-1 and it further increases to 295 W m-1 K-1 for the specimen produced from the deposit sprayed at 807 m s-1. Similarly, the electrical conductivity of the specimen prepared at 578 m s-1 is only 48% IACS, and then it increases to 68% IACS for that prepared at 745 m/s, and it further increases to 78% IACS for that prepared at 807 m s-1. The obviously improved electrical and thermal conductivities are attributed to high particle impact velocity enhancing the plastic
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deformation of particle, which is benefit for the inter-particle bonding in cold-sprayed deposit.
To quantitatively evaluate the effect of particle impact velocity on inter-particle
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bonding and further on the conductivity of the corresponding deposit, a brick wall structure mode was established to estimate the inter-particle electrical resistance
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values and thereby predict the electrical conductivity of Cu deposits prepared at different impact velocities. In this model, each deposited particle was simplified to be
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a cuboid-shaped “brick” with average dimensions of Ls Ls ts and inter-particle boundary thickness of ti. The oxide scale segments at the interface of the deposited
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particle come from the bottom of the starting Cu particle. Since the oxide scale peeloff is neglected due to its nanometer thickness, the area of the oxide scale (AO) segments at inter-particle boundaries can be calculated by packing them back to the
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bottom of the original particle segments as shown in Fig. 12. Therefore, AO at the interface can be achieved by calculating the surface area the spherical crown of the
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initial powder particle that penetrates into the previously deposited layer as shown in Eq. (4).
𝐴o = 2π𝑟ℎ = 2π𝑟 2 (1 − cos 𝜃)
(4)
where r is the radius of the particle; h is the height of the spherical crown that penetrates into the previously deposited layer; θ is the central angle corresponding to the arc of the initial spray particle which becomes the particle interface after particle deposition. According to the results reported by Assadi et al. [43] and Grujicic et al. 13
[44], the occurrence of adiabatic shear instability (ASI) is an important sign of bonding at the boundary between particles. At the region where ASI occurs, the interfacial material behaves like a viscous material and experiences severe plastic deformation, extruded from the boundary and forming an outward material jet at the rim. This material jet helps to clean up the oxide scale fragments at the boundary, allowing fresh metallic surface to contact and further metallic bonding to occur. Therefore, the oxide scale at the deposited-particle boundary is mainly concentrated at the region without ASI and the location where ASI occurs at the boundary determined
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the amount of oxide scale after impact. Previous literatures [43, 44] indicate that ASI occurs at the bottom of the particle and at an angel to the center line of the particle (θ)
as shown in Fig. 12. In order to simplify the calculation in the brick mode, assuming θ = 45o and the oxide scale is concentrated in a rectangular area with a feature size of C
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C at the center of deposited-particle boundary, the deposit can be expressed to a
“brick-wall” as the schematic diagram shown in Fig. 13(a). The feature size of oxide
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scale at the center of deposited particle boundary (C) can be calculated according to its area and is expressed as:
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𝐶 = √𝐴o
(5)
Assuming that the deposit is composed of n layers of bricks and each layer has m
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bricks. The dimensions of a single brick were derived from the cross sectional microstructure observation results from the etched cross-sections and the boundary
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thickness was obtained from the TEM images. Higher particle velocity leads to a wider and thinner brick due to enhanced flattening of the deposited particles. Detailed
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dimensions of the bricks at various particle velocities are listed in Table 3. It is known that the electrical conductivity of deposits is dominated by the proportion of the well contacted and conductive pathway. Extra contact electrical resistance will be introduced by the inter-particle boundaries due to the oxide scale segmentations. A structure unit marked with the red frame was derived from the brick-wall structure to simplify the calculation. As electrical current transfers along the thickness direction, the movement direction of electron is parallel to the vertical boundary in the deposit, 14
so the vertical boundary has little influence on the electron transmission. Neglecting the interfacial resistance at the vertical direction, the simplified equivalent electrical circuit for the through-thickness current is displayed in Fig. 13(b). According to the equivalent electrical circuit displayed in Fig. 13(b), the total interfacial electrical resistance (RPI) at the parallel direction can be expressed as:
1 1 1 RPI Rr 2 RC RO
(6)
where Rr is the interfacial electrical resistance in the region without oxide scale at the
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rim of the boundary, RC is the electrical contact resistance between brick and oxide scale, RO is the electrical resistance of oxide scale in the through-thickness. Following the results of Holm [45], the electrical contact resistance is expressed as:
S O
(7)
4a
-p
RC
where ρS is the electrical conductivity of the brick, ρO is electrical conductivity of the
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oxide scale (ρO=0.16 ·m [46]), a is the radius of the conducting contact area.
Rsu 2 ( RS RPI )
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Therefore, the total electrical resistance for a structure unit (Rsu) can be expressed as: (8)
where RS is the electrical resistance of the brick, RVI is the electrical resistance
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involved by the parallel inter-particle interface along the through-thickness direction. Further, the relation between the total electrical resistance of the deposit along the
nRsu n 2 ( RS RPI ) 2 2m m
(9)
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R
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through-thickness and the deposit structure can be expressed as:
It is assumed that s (s=0.0172 m) is equal to that of the cold deformed bulk pure copper. RS and RPI were calculated for one single deposited particle. Fig. 14 shows the variation of the electrical conductivity of the corresponding structure model as a function of the particle impact velocity. To verify the validity of the structure model, the experiment results of the electrical conductivities of deposit were also plotted in 15
Fig. 14. It can be found that the electrical conductivity calculated from the corresponding structure model is very close to the experimental results of the deposits, indicating that the hypothesis in the structure model is reasonable and the calculated results are trustable. A comparison of the calculated results for different particle impact conditions are listed in Table 3. It can be found that as the particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807m s-1, a remarkable decrease of 4.7 and 11.6 time was achieved in RPI suggesting a significant increase in inter-particle bonding along through-thickness direction. Although the thickness of the
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inter-particle interface is just ~0.1 m, the electrical resistance increment induced by the series connected inter-particle interface (RPI) can be as high as 28.9%, 14.5% and
13.8% of the total resistance (RPI+RS) for the impact conditions of 578 m s-1, 745m s-1 and 807 m s-1, respectively. This fact suggests that improving the inter-particle
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bonding (decreasing RPI) and reducing the number of inter-particle boundary in unit
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length significantly benefit to the improvement of the eletrical conductivity.
Although only electrical properties of the splats and specimens were calculated
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based on the brick-wall structure to establish the correlation with the microstructure of deposit, thermal conductivity, mechanical properties and other properties dominated
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by inter-particle bonding may be also predicted by the same approach. Because the deposit microstructure is closed related with particle impact condition, the correlation of particle impact condition with deposit properties can be established, which allows
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tailoring the properties of the deposits from the spraying conditions (impact velocity).
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4. Conclusion
In this work, cold sprayed Cu deposits were prepared at different particle impact
conditions. The deposition behavior of feedstock powder, the microstructure and properties of the corresponding deposits, including mechanical properties and conductivities were systematically investigated along the depositing direction. As the average particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1, the deposition efficiency of the Cu powder increases from 58% to 95%. Due to the 16
low hardness and high plastic deformability of Cu, the deposits prepared at all conditions in this study reveal dense microstructure with apparent porosity lower than 0.6%, and however the deformation degree of deposited particle obviously increases as the impact velocity increases. The enhanced plastic deformation of the deposited particles makes the oxide scale arising from the starting powder surface more dispersed along the inter-particle boundary and thereby a stronger inter-particle bonding. Therefore, both mechanical properties and conductivities of the deposits improve significantly. The elastic modulus and tensile strength of the specimen
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produced from the deposit sprayed at 807 m s-1 in the through-thickness reach to 90.4 GPa and 148.5 MPa, respectively. The thermal conductivity of the specimen reaches
to 295 W m-1 K-1 and its electrical conductivity comes up to 78% IACS. These results
indicate that it is not rational to evaluate the quality of cold sprayed metals only by the
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porosity. It can also be suggested that even though the Cu particles are accelerated to a
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high velocity of 807 m s-1, full metallurgy inter-particle bonding cannot be formed.
Acknowledgements
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This work is finacially supported by the National Nature Science Foundation of China (Grant Nos. 51875443 and 51401158), the Shaanxi Co-Innovation Projects (Grant No.
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2015JQ5200).
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2015KTTSGY03-03) and the Shaanxi Natural Science Foundation (Grant No.
17
References [1] W. Li, K. Yang, S. Yin, X. Yang, Y. Xu, Lupoi R, J. Mater. Sci. Technol. 35 (2018) 440-557. [2] N.H. Tariq, L. Gyansah, X. Qiu, C. Jia, H.B. Awais, C. Zheng, H. Du, J.Q. Wang, T.Y. Xiong, J. Mater. Sci. Technol. 35 (2019) 1053-1063. [3] K. Yang, W. Li, X. Guo, X. Yang, Y. Xu, J. Mater. Sci. Technol. 34 (2018) 1570-1579. [4] C. Huang, W. Li, M.P. Planche, H. Liao, G. Montavon, J. Mater. Sci. Technol. 33 (2017) 507-515. [5] C. Huang, W. Li, Y. Xie, M.P. Planche, H. Liao, G. Montavon, J. Mater. Sci. Technol. 33 (2017) 338-346. [6] R. Jenkins, S. Yin, B. Aldwell, M. Meyer, R. Lupoi, J. Mater. Sci. Technol. 35 (2019) 427-431. [7] K. Yang, W. Li, C. Huang, X. Yang, Y. Xu, J. Mater. Sci. Technol. 34 (2018) 2167-2177. [8] S. Yin, W. Li, B. Song, X. Yan, M. Kuang, Y. Xu, K. Wen, R. Lupoi, J. Mater. Sci. Technol. 35
ro of
(2019) 1003-1007.
[9] M. Kamaraj, V.M. Radhakrishnan,, J. Mater. Sci. Technol. 28 (2019) 756-768.
[10] Y. Ichikawa, R. Tokoro, M. Tanno, K. Ogawa, Acta Mater. 164 (2019) 39-49. [11] R. Fernandez, B. Jodoin, J. Therm. Spray Technol. 28 (2019) 737-755.
[12] M.S. Baek, Y.K. Kim, K.A. Lee, J. Therm. Spray Technol. 28 (2019) 917-929.
-p
[13] S. Theimer, M. Graunitz, M. Schulze, F. Gaertner, T. Klassen, J. Therm. Spray Technol. 28 (2019) 124-134.
[14] B.A. Bedard, T.J. Flanagan, A.T. Ernst, A. Nardi, A.M. Dongare, H.D. Brody, J. Therm. Spray
re
Technol. 27 (2018) 1563-1578.
[15] Q. Blochet, F. Delloro, F. N’Guyen, D. Jeulin, F. Borit, M. Jeandin, J. Therm. Spray Technol. 26 (2017) 671-686. 769-779.
lP
[16] Y. Xie, C. Chen, M.P. Planche, S. Deng, R. Huang, Z. Ren, J. Therm. Spray Technol. 27 (2019) [17] D. Cetiner, A.H. Paksoy, O. Tazegul, M. Baydogan, H. Guleryuz, H. Cimenoglu, J. Therm. Spray Technol. 27 (2018) 1414-1427.
na
[18] Y. Li, Y. Hamada, K. Otobe, T. Ando, J. Therm. Spray Technol. 26 (2016) 350-359. [19] Y.K. Wei, X.T. Luo, C.X. Li, C.J. Li, J. Therm. Spray Technol. 26 (2016)173-183. [20] E. Lin, Q. Chen, O.C. Ozdemir, V.K. Champagne, S. Müftü, J. Therm. Spray Technol. 28 (2019)
ur
472-483.
[21] H. Zhao, X. Ning, C. Tan, X. Yu, Z. Nie, X. Sun, J. Therm. Spray Technol. 28 (2019) 780-793. [22] S. J. Michael, C. Desrayaud, R.R. Chromik, S. Descartes, Mater. Sci. Eng., A 684 (2017) 510-
Jo
516.
[23] S. Kumar, S.K. Reddy, S.V. Joshi, Surf. Coat. Technol. 318 (2017) 62-71. [24] Y.J. Li, X.T. Luo, H. Rashid, C.J. Li, J. Alloys Compd. 740 (2018) 406-413. [25] Y.K. Wei, Y.J. Li, Y. Zhang, X.T. Luo, C.J. Li, Corros. Sci. 138 (2018)105-115. [26] M. Yu, M.Y. Ding, H. Ma, H.L. Liao, J. Therm. Spray Technol. 28 (2018) 460-471. [27] R. Huang, H. Fukanuma, J. Therm. Spray Technol. 21 (2011) 541-549. [28] R. Huang, W. Ma, H. Fukanuma, Surf. Coat. Technol. 258 (2014) 832-841. [29] F. Gärtner, T. Stoltenhoff, J. Voyer, H. Kreye, S. Riekehr, M. Koçak, Surf. Coat. Technol. 200 (2006) 6770-6782. [30] C.H. Zhou, H.T. Ma, L. Wang,, Corros. Sci. 52 (2010) 210-215. 18
[31] S. Cabanas-Polo, R. Bermejo, B. Ferrari, A.J Sanchez-Herencia, Corros. Sci. 55 (2012) 172-179. [32] E. A. Gulbransen, in: Ninety-first General Meeting, 1947, pp.573-604. [33] W. E. Campbell, U. B. Thomas, in: Ninety-first General Meeting, 1947, pp623-640. [34] S. Yin, X. Wang, W. Li, H. Liao, H. Jie, Appl. Surf. Sci. 259 (2012) 294-300. [35] W.Y. Li, H. Liao, C.J. Li, H.S. Bang, C. Coddet, Appl. Surf. Sci. 253 (2007) 5084-5091. [36] W.Y. Li, C.J. Li, H. Liao, Appl. Surf. Sci. 256 (2010) 4953-4958. [37] W.Y. Li, H. Liao, H.T. Wang, C.J. Li, G. Zhang, C. Coddet, Appl. Surf. Sci. 253 (2006) 708-713. [38] W.Y. Li, H. Liao, G. Douchy, C. Coddet, Mater. Des. 28 (2007) 2129-2137. [39] S. Yin, R. Jenkins, X. Yan, R. Lupoi. Mater. Sci. Eng. A 734 (2018) 67-76. [40] K. Yang, W. Li, X. Yang, Y. Xu. Surf. Coat. Technol. 335 (2018) 219-227. [41] X.T. Luo, C.J. Li, Mater. Des.140 (2017) 387-399. [42] W.Y. Li, C. Zhang, H. Liao, C. Coddet, J. Coat. Technol. Res. 6 (2008) 401-406. [43] H. Assadi, F. Gärtner, T. Stoltenhoff, H. Kreye, Acta Mater. 51 (2003) 4379-4394.
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[44] M. Grujicic, C.L. Zhao, W.S. DeRosset, D. Helfritch, Mate. Des. 25(2004) 681-688.
[45] R. Holm, Electric Contacts: Theory and Application, Springer Science & Business Media, 2013. [46] L. De Los Santos Valladares, D.H. Salinas, A.B. Dominguez, D.A. Najarro, S.I. Khondaker, T.
Jo
ur
na
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re
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Mitrelias, Thin Solid Films 520 (2012) 6368-6374.
19
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Figure and table lists：
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Fig. 1. Morphology (a) and particle size distribution (b) of Cu powder particles
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Fig. 2. Sampling strategy (a) and dimensions (b) of plate tensile specimen
20
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Fig. 3. Measured velocity distributions of in-flight particles at different spraying parameters by a commercial thermal spray particle diagnostic DPV 2000: (a) 2.5 MPa
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na
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500 oC, (b) 3 MPa 800 oC and (c) 5 MPa 800 oC.
Fig. 4. Deposition efficiency of the feedstock powder at different particle impact velocities
21
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Fig. 5. Microstructure of the deposits prepared at different particle impact velocities showing difference in deformation degree of particles, (a) and (b) 578 m s-1, (c) and
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(d) 745 m s-1, (e) and (f) 807 m s-1.
22
Vertical Line
30 μm
30 μm
30 μm
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Fig. 6. Measurement method of interface density in through-thickness direction
Fig. 7. Through-thickness interface densities of deposits produced at different particle
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impact velocities measured from the etched cross sections of deposits.
23
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Fig. 8. Bright-fielder TEM images and EDS mapping images at center of interparticle interface in specimens sprayed at 578 m s-1 (a, b, c and d) and 807 m s-1 (e, f, g and h); (a) and (e) SEM image, (b) and (f)TEM image, (c) and (g) STEM-EDS
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element maps of Cu, (d) and (h) STEM-EDS element maps of oxygen
24
Fig. 9. Stress-strain curves of the samples produced at different spraying parameters
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na
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re
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along the through-thickness direction.
25
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Fig.10. Fracture surface morphologies of samples produced at different particle
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impact velocities along through-thickness direction; (a) and (b) 578 m s-1, (c) and (d)
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745 m s-1, (e) and (f) 807 m s-1.
26
Fig. 11. Thermal (a) and electrical (b) conductivities of deposit samples prepared at
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different particle impact velocities along through-thickness direction. The blue line represents the thermal conductivity of annealed high purity Cu at 20 °C and red marks
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present the observed data.
Fig. 12. Schematic diagram showing oxide scale fragments at deposited-particle
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interface converted to oxide scale atinterface of initial particle.
27
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28
Fig. 14. Electrical conductivities of specimens calculated according to brick-wall
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na
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model as function with particle velocities
29
Table 1 Cold spray parameters for deposit preparation Main gas Carrier Gas Powder pressure / gas temperature feeding MPa pressure / / oC rate / g MPa min-1 low 2.5 2.6 500 120 Medium 3.0 3.1 800 120 High 5.0 5.1 800 120
Gun traverse speed / mm s-1 300 300 300
Stand-off distance / mm 20 20 20
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Table 2 Tensile strength and elastic modulus ofsamples produced at different particle impact velocities Particle impact velocities / m s-1
Tensile strength
Elastic modulus
/ MPa
/ GPa
578
25.8±5.5
745
86.8±7.9
807
148.5±8.2
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50.1±2.5
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69.2±1.5
m/s
m/s
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m/s
90.4±1.1
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Table 3 The measured dimensions and estimated electrical properties for the splats Ls /m
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Impact velocity
578 m s-1 33.6 745 m s-1 42.0 807 m s-1 53.2
ts /m
ti /m
10.5 6.9 4.8
0.1 0.1 0.1
/ ·m
Rs /
RPI /
R /
0.0172 0.0172 0.0172
0.000155 0.000065 0.000025
0.000063 0.000011 0.000004
0.00129 0.00096 0.00087
30
PII:
S1005-0302(19)30410-4
DOI:
https://doi.org/10.1016/j.jmst.2019.09.023
Reference:
JMST 1807
To appear in:
Journal of Materials Science & Technology
Received Date:
14 July 2019
Revised Date:
1 September 2019
Accepted Date:
23 September 2019
Please cite this article as: Li Y, Wei Y, Luo X, Li C, Ma N, Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits, Journal of Materials Science and Technology (2019), doi: https://doi.org/10.1016/j.jmst.2019.09.023
This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier.
Research Article Correlating particle impact condition with microstructure and properties of the cold-sprayed metallic deposits
Yujuan Lia, Yingkang Weia, Xiaotao Luoa,*, Changjiu Lia,*, Ninshu Mab
a
State Key Laboratory for Mechanical Behavior of Materials, School of Materials
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Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China b
Joining and Welding Research Institute, Osaka University, Osaka 567-0047, Japan
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*Corresponding author.
E-mail addresses: [email protected] (X.T. Luo), [email protected]
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(C.J. Li)
September 2019]
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[Received 14 July 2019; Received in revised form 1 September 2019; Accepted 23
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Inter-particle bonding is an important factor affecting the property of cold sprayed metallic deposit. Because the interface bonding between particles in deposit is directly
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determined by plastic strain of particles during spraying, Cu deposits were made at series of impact velocities of 578 m s-1 to 745 m s-1 and 807 m s-1 to correlate particle
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impact condition with microstructure and properties of the deposits. Results show that as the average particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1, the deposition efficiency of feedstock powder increases from 58% to 84% and even to 95%. Although all three deposits reveal dense microstructure due to the high ductility of Cu, the deformation degree of the deposited particles remarkably increases with increasing impact velocity. The enhanced plastic deformation of the deposited particles leads to more dispersed oxide scale and thereby stronger inter-particle 1
bonding with the strength of the deposit along the deposition direction increasing from 25.8 MPa to 148.5 MPa. The electrical and thermal conductivities at throughthickness direction of the deposit at particle impact velocity of 807 m s-1 are 78 % IACS, 295 W m-1 K-1, respectively.
Keywords: Cold spray, Particle impact velocity, Deposition efficiency, Electrical
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conductivity, Thermal conductivity, Mechanical property
1. Introduction
Cold spray was being well known as a surface coating technology in the past
decades, and recently it has attracted wide attention because it can also be a promising
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additive manufacturing (AM) process [1-3]. In cold spray, micron-scale particles
(typically 5-50 μm) are fed into supersonic-speed gas flow and are accelerated to 300-
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1200 m s-1, and then impact on the substrate at fully solid state [4-6]. The particle bonding is formed by high velocity impact induced intensive plastic deformation of
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the particle with a deposition temperature well below the melting point of a depositing material [7, 8]. As compared with melting and re-solidification based coating and AM
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processes such as thermal spray, laser cladding and selective laser melting, the low depositing temperature is one of the most remarkable characteristics of cold spray [911]. This makes depositing metallic materials without oxidation realized in open air
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and heating effect to the substrate avoided. Cu [12, 13], Al [14, 15], Ni [16], Ti [17], Cu-based alloys [18], Al-based alloys [19, 20], Ni-based alloys [21], Fe-based alloys
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[8]) and metal-based composites such as Al-Al2O3 [22] and Al-SiC [23] in the forms of coating or parts have been successfully achieved by cold spray so far.
Different from the metallurgical bulk counterparts, cold sprayed deposits are formed through the accumulated deposition of feedstock particles with a diameter of 5-50 μm in completely solid state. Besides the intrinsic characteristic of the depositing material, the inter-particle bonding is a decisive factor influencing penetration of the 2
corrosive media and the transfer of load and electron from one particle to another, and therefore determining the deposit properties, especially strength [13], thermal and electrical conductivities [24], and corrosion resistance [25, 26]. However, the interparticle bonding in cold sprayed deposit may be poor, and pores may appear between particles as the plastic deformation degree of particles is relatively small. Poor interparticle bonding and pores between particles in deposit lead to high electrical resistance, high thermal resistance and low corrosion resistance, thereby resulting in poor performances. In our previous results [24], Cu deposit was prepared by nitrogen
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gas cold spray at a gas pressure 3 MPa and a gas temperature of 400 oC and found that the through-thickness thermal conductivity of the deposit were only 207 W m-1 K-1 (52% of that for annealed bulk Cu). Wei et al. [19] prepared Al6061 deposit on
magnesium alloy plates at 2.5 MPa and 300 oC, and reported that the corrosion
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behavior of the deposit was close to that of the substrate and its corrosion resistance
was far below that of Al6061 bulk. Moreover, many studies [27, 28] have shown that
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the tensile strength of cold sprayed deposit is only 1/3-1/4 of that for corresponding bulk. By further observing the fracture morphology of deposit，it is found that the
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between particles [29].
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crack mainly propagates along the inter-particle interface, indicating the poor bonding
The poor inter-particle bonding can be caused by the oxide scale distributed at the interface of the deposited particle in cold sprayed deposit. As we all known,
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nanometer thick oxide scale is inevitably formed for even all the commercial metallic powders during production and storage. As long as the metallic surface is exposed to
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air, oxide scale generates with a thickness in ranges of 0.3-14.6 nm for Ni [30, 31], 19-134 nm for Fe [32] and 3-76.9 nm for Cu [33]. During the high-velocity particle impact in cold spray, the particle undergoes plastic deformation. As the strain at the interface of the deposited particle reaches a critical value, the oxide scale on the surface of feedstock particle fractures into discontinuous segments, allowing the fresh metallic surface contact and thereby form bonding [34]. Greater interfacial strain will lead to more dispersed oxide scale and more fresh metallic contact, which benefit the 3
bonding formation at the interface and thereby improve the bonding [35, 36].
To improve inter-particle bonding of cold sprayed deposit and further improve deposit performance, it is necessary to increase the degree of particle plastic deformation during particle impact process. In the past decades, many studies have been done to increase the degree of particle plastic deformation during cold spray. Because the degree of particle plastic deformation is directly related to particle impact velocity, the most common method is to improve accelerating behavior of particle. Li
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et al. [37, 38] optimized the nozzle geometry to improve the accelerating behavior of powder particles, and further expect to improve inter-particle bonding and deposit’s performance while the time cost of deposit preparation is increased due to the long cycle of optimizing nozzle geometry. Because of the excellent acceleration
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performance of He gas for powder particle, Huang et al. [27] and Gärtner et al. [29] used He gas as accelerating gas during cold spray, and the obtained deposit showed
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extremely high strength, indicating inter-particle bonding was obviously improved. Above results show that a high particle imapct velocity obtained by adjusting spraying
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parameters is benefit for the inter-particle bonding and thereby the microstructure and propertis of cold sprayed deposit. However, the quantitative relationship between
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particle impact condition and the plastic deformation degree of deposited particle, and the influence of particle impact condition on the interfacial microstructure and property of the deposited particle in corresponding deposit which affect the whole
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property of deposit significantly are not clear.
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In this work, the cold sprayed deposits were prepared at different particle impact
velocities by using different spraying parameters. The particle impact velocity at corresponding spraying parameter was quantitatively characterized. The microstructure and properties, including mechanical property and conductivities, of the deposits prepared at different particle impact velocities were investigated. Meanwhile, the interfacial microstructure of the deposited particle in deposit was also observed. 4
2. Experimental 2.1. Feedstock powder A commercial available water-atomized Cu powder was used as the feedstock powder. The morphology and particle size distribution of Cu powder are shown in Fig. 1. Particle size distribution of the Cu powder was measured by a laser diffractometry (Laser Mastersizer, Malvern Instruments, Malvern, UK) as shown in Fig. 1(b), indicating that the average particle sizes (d50) of the Cu powder was 28
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μm. An oxygen content of 0.044 wt.% was detected for Cu feedstock powder using oxygen/nitrogen analyzer (TC-600, LECO, USA). Commercially pure Cu plates were employed as the substrates. Prior to the coating deposition, the substrates were grit-
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blasted by alumina grits.
2.2. Deposit preparation
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A commercial cold spray system (PCS-1000, Plasma Giken Co., Ltd.) was used for deposit preparation. In this system, a Laval nozzle was used with throat diameter
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of 2 mm, outlet diameter of 4.4 mm, and divergent section length of 170 mm. N2 gas was used as the processing gas. To investigate the influence of the particle impact
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velocity on the inter-particle bonding, the Cu deposits were deposited at three different spraying conditions as listed in Table 1 (named as low, medium and high conditions according to the gas pressure and temperature values). The particle
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velocities of feedstock powders at different spraying parameters were measured by a commercial thermal spray in-flight particle diagnostic system (DPV 2000, Tecnar
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Automation Ltd., St-Bruno, Qc, Canada) equipped with a laser system (CPS-2000). The particle velocity measurement was made at the centerline of the gas flow, 20 mm away from the nozzle outlet which was same as the standoff distance during deposit preparation. For each spraying condition, more than 500 particles were measured.
2.3. Deposit characterization Microstructures of deposit were characterized by scanning electron microscopy 5
(SEM, MIRA 3 LMH, TESCAN, Czech Republic) with an accelerating voltage of 15 kV. Porosity of deposits was measured by image analyzing method from cross section (10 SEM images at 1000× for each specimen). Transmission electron microscopy (TEM) samples were made from the etched cross section using focus ion beam (FIB, JIB-4500 Multi-beam System, JEOL) to examine the interfacial microstructure and track the oxide scale (TEM, JEM-F200, JEOL). Conductivity measurements and uniaxial tensile test were performed on the deposit to quantitatively evaluate the effect of inter-particle bonding quality on properties because there are more inter-particle
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boundaries in unit length along the through-thickness direction. Through-thickness electrical conductivity of deposit was calculated according to the resistance measured by using double bridge method (QJ57, China) with an accuracy of 10-5 Ω and the
samples used were free-standing deposit with dimensions of 40 mm × 1 mm × 1 mm.
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Through-thickness thermal conductivity of deposit was measured by using a laser flash technique (FLASHLINETM SYSTEM 3050, Anter Corporation, US) and the
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sample was also free-standing deposit with a dimension of 𝜙mm×3 mm. The plate tensile specimens with a thickness of 1 mm were sampled along the spray
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direction (Fig. 2). Tensile test was performed by using a material mechanical property test machine (Instron 5848, Micro Tester) according to EN 10002-1 standard. Detailed
3. Results
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dimensions of the plate tensile samples are also shown in Fig. 2.
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3.1. Particle impact velocity
To evaluate the particle impact velocities at different spraying conditions, the
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velocity of the particle was measured by using the DPV 2000 system at the center of flow under the conditions of preparing deposit as listed in Table 1. The measured particle velocity distributions of the particles at 2.5 MPa 500 oC, 3 MPa 800 oC and 5 MPa 800 oC are shown in Fig. 3(a, b and c), respectively. At those spraying conditions, all velocities of the particles range from 200 m s-1 to 1000 m s-1. As expected the average velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1 as the spraying condition varies from 2.5 MPa 500 oC to 3 MPa 800 oC and 5 MPa 6
800 oC.
3.2. Deposition efficiency of feedstock powder Deposition efficiency is one of the most important concerns for thermal/cold spraying of deposits since it partially determines the cost of the deposits. Powder deposition efficiency (DE) is the mass ratio of the prepared deposit to the powder sprayed on the substrate during the spraying process. The mass of deposit is the mass difference of the substrate before and after spraying. To characterize the mass
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difference of substrate, an electrical balance (Fisher Scientific, A-2000DS) with accuracy of 0.1 mg was used to measure the mass of the substrate before (Wsa) and after (Wsb) spraying. The mass of the powder consumed on the substrate can be
expressed by the product of powder feeding rate and the spraying time spent on the
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substrate during spraying. Powder feeding rate is the mass of powder consumed per
unit time. To calculate the powder feeding rate, the mass of the powder in the powder
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feeder before (Wpa) and after (Wpb) spraying was measured by using an electrical balance, and the total spraying time (T) was recorded by a stopwatch. To estimate the
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spraying time consumed on the substrate, the length of substrate (L) was measured, and the transverse movement velocity of nozzle (V), the pass number of spraying (N)
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and the stroke number in one pass spraying (n) were also recorded. The detailed expression to calculate the deposition efficiency of sprayed powder can be seen in Eq. (1):
𝑊sb −𝑊sa 𝑊pa−𝑊 pb 𝐿 × ×𝑁×𝑛 𝑇 𝑉
× 100%
(1)
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DE =
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The deposition efficiency of the Cu powder was measured for more than 3 times at each spraying condition and its average value was taken.
The deposition efficiency of the copper powders as a function of the particle
velocity is shown in Fig. 4. The deposition efficiency increases from 58% to 84% as the average particle velocity increases from 578 m s-1 to 745 m s-1. It is worth noting that a high deposition efficiency of 95% was achieved as the average particle velocity 7
increases to 807 m s-1. As the average particle impact velocity increases obviously, the number of the particles whose velocity is greater than the critical velocity required for particle deposition increase. Therefore, the deposition efficiency of powder particle significantly increases.
3.3. Microstructure of deposits To study the influence of particle impact velocity on microstructure of the cold sprayed Cu, the polished cross sections of the deposits were observed by SEM and are
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shown in Fig. 5(a, c and e). To track the inter-particle boundaries, the polished cross sections of deposit samples were also etched, where inter-particle boundaries can be easily recognized as shown in Fig. 5(b, d and f). Although all three deposits show very dense microstructures with apparent porosity values lower than 0.6%, the
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deformation degrees of particles at different spraying parameters are totally different since different particle impact velocities provide different driving forces for the
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impact induced plastic deformation. As shown in Fig. 5(b, d and f), the deposited particles in the deposits prepared at higher particle velocities are much highly
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flattened. To quantitatively evaluate the difference, flattening ratios were estimated by the image analyzing. The flattening ratio (R) of a particle is defined as: 𝐷
p
(2)
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𝑅 = 𝑑 × 100%
where D is the maximum dimension of the deposited particles perpendicular to the
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depositing direction, and dp is the diameter of the raw particles. The average particle size of the raw Cu particles (28 μm) was taken as dp. D was measured from the SEM
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images of etched cross-section of the deposits. For each deposit, the number of particles counted was more than 200. This results in flattening ratios of 1.2±0.38, 1.5 ±0.41 and 1.9±0.27 for three conditions, respectively.
For cold sprayed deposit, as it is composed of many splats, its performance is highly determined by the inter-particle bonding, including both the number of the inter-particle boundaries in unit length and bonding quality. Either less boundaries or 8
higher bonding quality favors the transfer of current, heat and load across the particles and thereby higher conductivities, and strength and elastic modulus. The interface quantity in a deposit can be characterized by using inter-particle boundary density which is defined as number of the inter-particle boundaries in unit length along a fixed direction. In this work, the inter-particle boundary density is derived from the etched cross-sections of the deposits. To calculate the interface density (dinter) along through-thickness direction, straight lines with distance of 30 m along the depositing direction in the 3000 SEM images of the etched cross-sections of the deposit as
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shown in Fig. 6. The number of interface (N) along each line was counted. dinter can be calculated by the ratio of N to the length of the vertical line as the relation shown in Eq. (3).
N L
(Eq.3)
-p
dint er
For each deposit, at least 10 measurements were made and the average value was used
re
for comparison.
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Fig. 7 shows the through-thickness inter-particle densities of deposits produced at different particle impact velocities. It can be seen that the through-thickness
na
interface density increases with the increase of particle impact velocity due to the enhanced plastic deformation. The statistical result indicates interface density in the
m s-1.
ur
deposit sprayed at 807 m s-1 is 2 times higher than that in the deposit prepared at 578
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Quality of inter-particle bonding is another important concern which determines
the properties of the cold sprayed metals. According to our previous studies [35], the oxide scale arising from the starting particle surfaces will be trapped at inter-particle boundaries, and hinder the metallurgical bonding formation between deposited particles. Therefore, the oxide scale along the inter-particle boundaries was tracked in this work to assess the quality of the inter-particle bonding formed at different particle impact velocities. It is known that during the high velocity particle impact, the ductile 9
particle will be severely plastically deformed. However, the nanometer-thick oxide scale attached on the particle surface would be segmented rather than be deformed due to its brittle nature. Between the oxide-scale segments, clean metallic surface contact and thereby metallurgical bonding can be realized. However, due to the small thickness of oxide scale on the initial surface of powder particles (less than 100 nm), it is difficult to observe the oxide scale at the boundary between particles after deposition by SEM [34, 36]. The fraction of newly created metallic surface is determined by the plastic strain of the interface. It is known that the plastic
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deformation of particle is heterogeneous. The edge area of the interface always undergoes much higher plastic strain than the center area. This means a much clearer difference in dispersion degree of the oxide scale can be observed from the TEM
images taken at the edge area of the interface for deposits made at different particle
-p
impact velocities. However, in this work, all the properties including strength,
electrical conductivity and thermal conductivity were measured along the depositing
re
direction. It is known that the properties along the depositing direction are influenced more by the oxide scale vertical to the depositing direction. Therefore, in this work,
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TEM images were only taken at the center area of the particle interface. TEM samples were made by focus ion beam (FIB) at the center of the interface in the Cu deposits
na
prepared at 578 m s-1 (Fig. 8(a)) and 807 m s-1 (Fig. 8(e)). TEM images and the EDS mapping results are also presented in Fig. 8. The continuously-distributed oxide in a white contrast was observed at the center of the boundary in the specimen sprayed at
ur
578 m s-1 (Fig. 8(b)), while the oxide scale in the white contrast became discontinuous at the boundary in the specimen produced at 807 m s-1 (Fig. 8(f)).
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TEM-EDS mapping results confirm that long oxide scale segments remained at the center of the inter-particle boundary (Fig. 8(d) and (h)). The more dispersed oxide scale at the bottom of the bonded particle resulted from a higher particle impact velocity will benefit for the bonding formation at inter-particle boundary and the corresponding deposit may show higher properties, compared with that sprayed at a lower particle impact velocity. The concentrated oxide scale at the bottom of the bonded particle may also contribute to the anisotropy of cold sprayed metallic 10
deposits since there are more oxide scale layers in unit length along the depositing direction. Yin et al. [39] and Yang et al. [40] reported that the properties anisotropy of cold sprayed deposit was caused by the bonding difference of the inter-particle interface, which is consistent with our results.
3.4. Elastic modulus and tensile strength of the deposits In cold/thermally sprayed metallic deposits, it is known that there are more particle interfaces in unit length along the through-thickness direction than those
ro of
along the in-plain direction. It is rational to infer that the properties along the throughthickness direction much highly depend on the particle bonding. Therefore, in this work, a uniaxial tensile test was performed on the deposits along the through-
thickness direction. Fig. 9 presents the stress-strain curves of the samples produced at
-p
different particle impact conditions. It can be seen that Cu deposit sprayed at the
highest particle velocity reveals 6 times higher fracture strength than that deposited at
re
lowest velocity, suggesting much stronger inter-particle bonding. On the other hand, higher slope of the stress-strain curve was detected with increasing particle impact
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velocities in the elastic deformation range, suggesting higher elastic modulus. Table 2 lists the tensile strength and elastic modulus of all samples. Both strength and elastic
na
modulus greatly increase as the particle impact velocity increases and the fracture strength of the deposit prepared at the highest particle velocity is very close to that of the bulk Cu. However, the elastic modulus values of all Cu deposits are lower than
ur
that of the bulk counterpart (124 GPa), which are 40.4%, 55.8% and 72.9% for the deposit sprayed at 578 m s-1, 745 m s-1 and 807 m s-1, respectively. This fact suggests
Jo
that although high strength can be achieved by depositing at very high spraying parameter resulted in extremely high particle impact velocity (807 m s-1 for Cu in this work), it is hard to achieve the dense Cu deposit with fully bonded particles.
Fig. 10 shows the typical fracture morphologies of the deposits produced at different particle impact conditions. It can be seen from Fig. 10(a, c and e) that the fracture occurs mainly at the inter-particle boundaries, indicating a mainly brittle 11
fracture, which is consistent with the small elongation (<0.2%). As the particle impact velocity increases, the particles are much highly deformed and therefore profiles of the individual particles become more difficult to be recognized. For the deposit prepared at 578 m s-1, no dimple is observed in the fracture morphology and the fracture takes place along the interfaces between deposited particles as shown in Fig. 10(b). However, the fracture morphologies of deposits sprayed at 745 m s-1 and 807 m s-1 show a small amount of dimples, although the fracture mainly occurs along the inter-particle interfaces in Fig. 10(d) and (f). Moreover, the deposited particle in the
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deposit sprayed at 807 m s-1 is more tightly bonded together comparing with that in the deposit produced at 745 m s-1. Therefore, a higher particle impact velocity leads to a better inter-particle bonding and thereby a high tensile strength in Fig. 9. It is worth noting that the ductility of the as-sprayed deposits is very limited as compared with
-p
the bulk counterpart due to the un-perfect inter-particle bonding and severe work
hardening of the deposited particles during deposition. This drawback could greatly
re
hinder the application of cold spray for additively manufacturing parts. As reported in literatures [41, 42], post-spray heat treatment is an effective way to improve the
na
especially dislocations.
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ductility by enhancing the inter-particle bonding and eliminating the crystal defects
3.5. Electrical and thermal conductivities of deposits It is known that Cu and its alloys are frequently used as conducting parts due to
ur
its high intrinsic thermal conductivity and electrical conductivity. Meanwhile, electron transfer is very sensitive to various boundaries due to their diffraction to the electrons.
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Therefore, in the present work, thermal conductivity and electrical conductivity of the deposits sprayed at varies particle impact velocities were experimentally examined. Since the inter-particle boundary density is higher in through-thickness direction, the specimens for thermal and electrical conductivities were also sampled along the depositing direction. Fig. 11 shows thermal conductivity and electrical conductivity of the cold sprayed Cu as a function of the particle impact velocity. Both conductivities increase with increasing particle velocity. The thermal conductivity of the specimen 12
produced from the deposit sprayed at 578 m s-1 is only 204 W m-1 K-1 (51% of that for annealed bulk Cu), and then it increases to 260 W m-1 K-1 for the specimen produced from the deposit sprayed at 745 m s-1 and it further increases to 295 W m-1 K-1 for the specimen produced from the deposit sprayed at 807 m s-1. Similarly, the electrical conductivity of the specimen prepared at 578 m s-1 is only 48% IACS, and then it increases to 68% IACS for that prepared at 745 m/s, and it further increases to 78% IACS for that prepared at 807 m s-1. The obviously improved electrical and thermal conductivities are attributed to high particle impact velocity enhancing the plastic
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deformation of particle, which is benefit for the inter-particle bonding in cold-sprayed deposit.
To quantitatively evaluate the effect of particle impact velocity on inter-particle
-p
bonding and further on the conductivity of the corresponding deposit, a brick wall structure mode was established to estimate the inter-particle electrical resistance
re
values and thereby predict the electrical conductivity of Cu deposits prepared at different impact velocities. In this model, each deposited particle was simplified to be
lP
a cuboid-shaped “brick” with average dimensions of Ls Ls ts and inter-particle boundary thickness of ti. The oxide scale segments at the interface of the deposited
na
particle come from the bottom of the starting Cu particle. Since the oxide scale peeloff is neglected due to its nanometer thickness, the area of the oxide scale (AO) segments at inter-particle boundaries can be calculated by packing them back to the
ur
bottom of the original particle segments as shown in Fig. 12. Therefore, AO at the interface can be achieved by calculating the surface area the spherical crown of the
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initial powder particle that penetrates into the previously deposited layer as shown in Eq. (4).
𝐴o = 2π𝑟ℎ = 2π𝑟 2 (1 − cos 𝜃)
(4)
where r is the radius of the particle; h is the height of the spherical crown that penetrates into the previously deposited layer; θ is the central angle corresponding to the arc of the initial spray particle which becomes the particle interface after particle deposition. According to the results reported by Assadi et al. [43] and Grujicic et al. 13
[44], the occurrence of adiabatic shear instability (ASI) is an important sign of bonding at the boundary between particles. At the region where ASI occurs, the interfacial material behaves like a viscous material and experiences severe plastic deformation, extruded from the boundary and forming an outward material jet at the rim. This material jet helps to clean up the oxide scale fragments at the boundary, allowing fresh metallic surface to contact and further metallic bonding to occur. Therefore, the oxide scale at the deposited-particle boundary is mainly concentrated at the region without ASI and the location where ASI occurs at the boundary determined
ro of
the amount of oxide scale after impact. Previous literatures [43, 44] indicate that ASI occurs at the bottom of the particle and at an angel to the center line of the particle (θ)
as shown in Fig. 12. In order to simplify the calculation in the brick mode, assuming θ = 45o and the oxide scale is concentrated in a rectangular area with a feature size of C
-p
C at the center of deposited-particle boundary, the deposit can be expressed to a
“brick-wall” as the schematic diagram shown in Fig. 13(a). The feature size of oxide
re
scale at the center of deposited particle boundary (C) can be calculated according to its area and is expressed as:
lP
𝐶 = √𝐴o
(5)
Assuming that the deposit is composed of n layers of bricks and each layer has m
na
bricks. The dimensions of a single brick were derived from the cross sectional microstructure observation results from the etched cross-sections and the boundary
ur
thickness was obtained from the TEM images. Higher particle velocity leads to a wider and thinner brick due to enhanced flattening of the deposited particles. Detailed
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dimensions of the bricks at various particle velocities are listed in Table 3. It is known that the electrical conductivity of deposits is dominated by the proportion of the well contacted and conductive pathway. Extra contact electrical resistance will be introduced by the inter-particle boundaries due to the oxide scale segmentations. A structure unit marked with the red frame was derived from the brick-wall structure to simplify the calculation. As electrical current transfers along the thickness direction, the movement direction of electron is parallel to the vertical boundary in the deposit, 14
so the vertical boundary has little influence on the electron transmission. Neglecting the interfacial resistance at the vertical direction, the simplified equivalent electrical circuit for the through-thickness current is displayed in Fig. 13(b). According to the equivalent electrical circuit displayed in Fig. 13(b), the total interfacial electrical resistance (RPI) at the parallel direction can be expressed as:
1 1 1 RPI Rr 2 RC RO
(6)
where Rr is the interfacial electrical resistance in the region without oxide scale at the
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rim of the boundary, RC is the electrical contact resistance between brick and oxide scale, RO is the electrical resistance of oxide scale in the through-thickness. Following the results of Holm [45], the electrical contact resistance is expressed as:
S O
(7)
4a
-p
RC
where ρS is the electrical conductivity of the brick, ρO is electrical conductivity of the
re
oxide scale (ρO=0.16 ·m [46]), a is the radius of the conducting contact area.
Rsu 2 ( RS RPI )
lP
Therefore, the total electrical resistance for a structure unit (Rsu) can be expressed as: (8)
where RS is the electrical resistance of the brick, RVI is the electrical resistance
na
involved by the parallel inter-particle interface along the through-thickness direction. Further, the relation between the total electrical resistance of the deposit along the
nRsu n 2 ( RS RPI ) 2 2m m
(9)
Jo
R
ur
through-thickness and the deposit structure can be expressed as:
It is assumed that s (s=0.0172 m) is equal to that of the cold deformed bulk pure copper. RS and RPI were calculated for one single deposited particle. Fig. 14 shows the variation of the electrical conductivity of the corresponding structure model as a function of the particle impact velocity. To verify the validity of the structure model, the experiment results of the electrical conductivities of deposit were also plotted in 15
Fig. 14. It can be found that the electrical conductivity calculated from the corresponding structure model is very close to the experimental results of the deposits, indicating that the hypothesis in the structure model is reasonable and the calculated results are trustable. A comparison of the calculated results for different particle impact conditions are listed in Table 3. It can be found that as the particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807m s-1, a remarkable decrease of 4.7 and 11.6 time was achieved in RPI suggesting a significant increase in inter-particle bonding along through-thickness direction. Although the thickness of the
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inter-particle interface is just ~0.1 m, the electrical resistance increment induced by the series connected inter-particle interface (RPI) can be as high as 28.9%, 14.5% and
13.8% of the total resistance (RPI+RS) for the impact conditions of 578 m s-1, 745m s-1 and 807 m s-1, respectively. This fact suggests that improving the inter-particle
-p
bonding (decreasing RPI) and reducing the number of inter-particle boundary in unit
re
length significantly benefit to the improvement of the eletrical conductivity.
Although only electrical properties of the splats and specimens were calculated
lP
based on the brick-wall structure to establish the correlation with the microstructure of deposit, thermal conductivity, mechanical properties and other properties dominated
na
by inter-particle bonding may be also predicted by the same approach. Because the deposit microstructure is closed related with particle impact condition, the correlation of particle impact condition with deposit properties can be established, which allows
ur
tailoring the properties of the deposits from the spraying conditions (impact velocity).
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4. Conclusion
In this work, cold sprayed Cu deposits were prepared at different particle impact
conditions. The deposition behavior of feedstock powder, the microstructure and properties of the corresponding deposits, including mechanical properties and conductivities were systematically investigated along the depositing direction. As the average particle impact velocity increases from 578 m s-1 to 745 m s-1 and 807 m s-1, the deposition efficiency of the Cu powder increases from 58% to 95%. Due to the 16
low hardness and high plastic deformability of Cu, the deposits prepared at all conditions in this study reveal dense microstructure with apparent porosity lower than 0.6%, and however the deformation degree of deposited particle obviously increases as the impact velocity increases. The enhanced plastic deformation of the deposited particles makes the oxide scale arising from the starting powder surface more dispersed along the inter-particle boundary and thereby a stronger inter-particle bonding. Therefore, both mechanical properties and conductivities of the deposits improve significantly. The elastic modulus and tensile strength of the specimen
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produced from the deposit sprayed at 807 m s-1 in the through-thickness reach to 90.4 GPa and 148.5 MPa, respectively. The thermal conductivity of the specimen reaches
to 295 W m-1 K-1 and its electrical conductivity comes up to 78% IACS. These results
indicate that it is not rational to evaluate the quality of cold sprayed metals only by the
-p
porosity. It can also be suggested that even though the Cu particles are accelerated to a
re
high velocity of 807 m s-1, full metallurgy inter-particle bonding cannot be formed.
Acknowledgements
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This work is finacially supported by the National Nature Science Foundation of China (Grant Nos. 51875443 and 51401158), the Shaanxi Co-Innovation Projects (Grant No.
Jo
ur
2015JQ5200).
na
2015KTTSGY03-03) and the Shaanxi Natural Science Foundation (Grant No.
17
References [1] W. Li, K. Yang, S. Yin, X. Yang, Y. Xu, Lupoi R, J. Mater. Sci. Technol. 35 (2018) 440-557. [2] N.H. Tariq, L. Gyansah, X. Qiu, C. Jia, H.B. Awais, C. Zheng, H. Du, J.Q. Wang, T.Y. Xiong, J. Mater. Sci. Technol. 35 (2019) 1053-1063. [3] K. Yang, W. Li, X. Guo, X. Yang, Y. Xu, J. Mater. Sci. Technol. 34 (2018) 1570-1579. [4] C. Huang, W. Li, M.P. Planche, H. Liao, G. Montavon, J. Mater. Sci. Technol. 33 (2017) 507-515. [5] C. Huang, W. Li, Y. Xie, M.P. Planche, H. Liao, G. Montavon, J. Mater. Sci. Technol. 33 (2017) 338-346. [6] R. Jenkins, S. Yin, B. Aldwell, M. Meyer, R. Lupoi, J. Mater. Sci. Technol. 35 (2019) 427-431. [7] K. Yang, W. Li, C. Huang, X. Yang, Y. Xu, J. Mater. Sci. Technol. 34 (2018) 2167-2177. [8] S. Yin, W. Li, B. Song, X. Yan, M. Kuang, Y. Xu, K. Wen, R. Lupoi, J. Mater. Sci. Technol. 35
ro of
(2019) 1003-1007.
[9] M. Kamaraj, V.M. Radhakrishnan,, J. Mater. Sci. Technol. 28 (2019) 756-768.
[10] Y. Ichikawa, R. Tokoro, M. Tanno, K. Ogawa, Acta Mater. 164 (2019) 39-49. [11] R. Fernandez, B. Jodoin, J. Therm. Spray Technol. 28 (2019) 737-755.
[12] M.S. Baek, Y.K. Kim, K.A. Lee, J. Therm. Spray Technol. 28 (2019) 917-929.
-p
[13] S. Theimer, M. Graunitz, M. Schulze, F. Gaertner, T. Klassen, J. Therm. Spray Technol. 28 (2019) 124-134.
[14] B.A. Bedard, T.J. Flanagan, A.T. Ernst, A. Nardi, A.M. Dongare, H.D. Brody, J. Therm. Spray
re
Technol. 27 (2018) 1563-1578.
[15] Q. Blochet, F. Delloro, F. N’Guyen, D. Jeulin, F. Borit, M. Jeandin, J. Therm. Spray Technol. 26 (2017) 671-686. 769-779.
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[16] Y. Xie, C. Chen, M.P. Planche, S. Deng, R. Huang, Z. Ren, J. Therm. Spray Technol. 27 (2019) [17] D. Cetiner, A.H. Paksoy, O. Tazegul, M. Baydogan, H. Guleryuz, H. Cimenoglu, J. Therm. Spray Technol. 27 (2018) 1414-1427.
na
[18] Y. Li, Y. Hamada, K. Otobe, T. Ando, J. Therm. Spray Technol. 26 (2016) 350-359. [19] Y.K. Wei, X.T. Luo, C.X. Li, C.J. Li, J. Therm. Spray Technol. 26 (2016)173-183. [20] E. Lin, Q. Chen, O.C. Ozdemir, V.K. Champagne, S. Müftü, J. Therm. Spray Technol. 28 (2019)
ur
472-483.
[21] H. Zhao, X. Ning, C. Tan, X. Yu, Z. Nie, X. Sun, J. Therm. Spray Technol. 28 (2019) 780-793. [22] S. J. Michael, C. Desrayaud, R.R. Chromik, S. Descartes, Mater. Sci. Eng., A 684 (2017) 510-
Jo
516.
[23] S. Kumar, S.K. Reddy, S.V. Joshi, Surf. Coat. Technol. 318 (2017) 62-71. [24] Y.J. Li, X.T. Luo, H. Rashid, C.J. Li, J. Alloys Compd. 740 (2018) 406-413. [25] Y.K. Wei, Y.J. Li, Y. Zhang, X.T. Luo, C.J. Li, Corros. Sci. 138 (2018)105-115. [26] M. Yu, M.Y. Ding, H. Ma, H.L. Liao, J. Therm. Spray Technol. 28 (2018) 460-471. [27] R. Huang, H. Fukanuma, J. Therm. Spray Technol. 21 (2011) 541-549. [28] R. Huang, W. Ma, H. Fukanuma, Surf. Coat. Technol. 258 (2014) 832-841. [29] F. Gärtner, T. Stoltenhoff, J. Voyer, H. Kreye, S. Riekehr, M. Koçak, Surf. Coat. Technol. 200 (2006) 6770-6782. [30] C.H. Zhou, H.T. Ma, L. Wang,, Corros. Sci. 52 (2010) 210-215. 18
[31] S. Cabanas-Polo, R. Bermejo, B. Ferrari, A.J Sanchez-Herencia, Corros. Sci. 55 (2012) 172-179. [32] E. A. Gulbransen, in: Ninety-first General Meeting, 1947, pp.573-604. [33] W. E. Campbell, U. B. Thomas, in: Ninety-first General Meeting, 1947, pp623-640. [34] S. Yin, X. Wang, W. Li, H. Liao, H. Jie, Appl. Surf. Sci. 259 (2012) 294-300. [35] W.Y. Li, H. Liao, C.J. Li, H.S. Bang, C. Coddet, Appl. Surf. Sci. 253 (2007) 5084-5091. [36] W.Y. Li, C.J. Li, H. Liao, Appl. Surf. Sci. 256 (2010) 4953-4958. [37] W.Y. Li, H. Liao, H.T. Wang, C.J. Li, G. Zhang, C. Coddet, Appl. Surf. Sci. 253 (2006) 708-713. [38] W.Y. Li, H. Liao, G. Douchy, C. Coddet, Mater. Des. 28 (2007) 2129-2137. [39] S. Yin, R. Jenkins, X. Yan, R. Lupoi. Mater. Sci. Eng. A 734 (2018) 67-76. [40] K. Yang, W. Li, X. Yang, Y. Xu. Surf. Coat. Technol. 335 (2018) 219-227. [41] X.T. Luo, C.J. Li, Mater. Des.140 (2017) 387-399. [42] W.Y. Li, C. Zhang, H. Liao, C. Coddet, J. Coat. Technol. Res. 6 (2008) 401-406. [43] H. Assadi, F. Gärtner, T. Stoltenhoff, H. Kreye, Acta Mater. 51 (2003) 4379-4394.
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[44] M. Grujicic, C.L. Zhao, W.S. DeRosset, D. Helfritch, Mate. Des. 25(2004) 681-688.
[45] R. Holm, Electric Contacts: Theory and Application, Springer Science & Business Media, 2013. [46] L. De Los Santos Valladares, D.H. Salinas, A.B. Dominguez, D.A. Najarro, S.I. Khondaker, T.
Jo
ur
na
lP
re
-p
Mitrelias, Thin Solid Films 520 (2012) 6368-6374.
19
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Figure and table lists：
na
lP
re
-p
Fig. 1. Morphology (a) and particle size distribution (b) of Cu powder particles
Jo
ur
Fig. 2. Sampling strategy (a) and dimensions (b) of plate tensile specimen
20
ro of -p
re
Fig. 3. Measured velocity distributions of in-flight particles at different spraying parameters by a commercial thermal spray particle diagnostic DPV 2000: (a) 2.5 MPa
Jo
ur
na
lP
500 oC, (b) 3 MPa 800 oC and (c) 5 MPa 800 oC.
Fig. 4. Deposition efficiency of the feedstock powder at different particle impact velocities
21
ro of -p re lP na
ur
Fig. 5. Microstructure of the deposits prepared at different particle impact velocities showing difference in deformation degree of particles, (a) and (b) 578 m s-1, (c) and
Jo
(d) 745 m s-1, (e) and (f) 807 m s-1.
22
Vertical Line
30 μm
30 μm
30 μm
lP
re
-p
ro of
Fig. 6. Measurement method of interface density in through-thickness direction
Fig. 7. Through-thickness interface densities of deposits produced at different particle
Jo
ur
na
impact velocities measured from the etched cross sections of deposits.
23
ro of -p re lP
na
Fig. 8. Bright-fielder TEM images and EDS mapping images at center of interparticle interface in specimens sprayed at 578 m s-1 (a, b, c and d) and 807 m s-1 (e, f, g and h); (a) and (e) SEM image, (b) and (f)TEM image, (c) and (g) STEM-EDS
Jo
ur
element maps of Cu, (d) and (h) STEM-EDS element maps of oxygen
24
Fig. 9. Stress-strain curves of the samples produced at different spraying parameters
Jo
ur
na
lP
re
-p
ro of
along the through-thickness direction.
25
ro of -p re lP na
Fig.10. Fracture surface morphologies of samples produced at different particle
ur
impact velocities along through-thickness direction; (a) and (b) 578 m s-1, (c) and (d)
Jo
745 m s-1, (e) and (f) 807 m s-1.
26
Fig. 11. Thermal (a) and electrical (b) conductivities of deposit samples prepared at
ro of
different particle impact velocities along through-thickness direction. The blue line represents the thermal conductivity of annealed high purity Cu at 20 °C and red marks
na
lP
re
-p
present the observed data.
Fig. 12. Schematic diagram showing oxide scale fragments at deposited-particle
Jo
ur
interface converted to oxide scale atinterface of initial particle.
27
ro of -p re lP na ur Jo Fig. 13. A simplified structure of deposit (a) and corresponding equivalent electrical circuit (b) of structure unit in the through-thickness direction
28
Fig. 14. Electrical conductivities of specimens calculated according to brick-wall
Jo
ur
na
lP
re
-p
ro of
model as function with particle velocities
29
Table 1 Cold spray parameters for deposit preparation Main gas Carrier Gas Powder pressure / gas temperature feeding MPa pressure / / oC rate / g MPa min-1 low 2.5 2.6 500 120 Medium 3.0 3.1 800 120 High 5.0 5.1 800 120
Gun traverse speed / mm s-1 300 300 300
Stand-off distance / mm 20 20 20
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Table 2 Tensile strength and elastic modulus ofsamples produced at different particle impact velocities Particle impact velocities / m s-1
Tensile strength
Elastic modulus
/ MPa
/ GPa
578
25.8±5.5
745
86.8±7.9
807
148.5±8.2
-p
50.1±2.5
lP
re
69.2±1.5
m/s
m/s
na
m/s
90.4±1.1
ur
Table 3 The measured dimensions and estimated electrical properties for the splats Ls /m
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Impact velocity
578 m s-1 33.6 745 m s-1 42.0 807 m s-1 53.2
ts /m
ti /m
10.5 6.9 4.8
0.1 0.1 0.1
/ ·m
Rs /
RPI /
R /
0.0172 0.0172 0.0172
0.000155 0.000065 0.000025
0.000063 0.000011 0.000004
0.00129 0.00096 0.00087
30