Al2O3 metal matrix composites

Journal of Alloys and Compounds 736 (2018) 115e123

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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Cold sprayed AA2024/Al2O3 metal matrix composites improved by friction stir processing: Microstructure characterization, mechanical performance and strengthening mechanisms Kang Yang, Wenya Li*, Pengliang Niu, Xiawei Yang, Yaxin Xu State Key Laboratory of Solidification Processing, Shaanxi Key Laboratory of Friction Welding Technologies, School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, 710072, PR China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 15 August 2017 Received in revised form 29 October 2017 Accepted 10 November 2017 Available online 11 November 2017

In this study, friction stir processing (FSP) was employed as the strengthening method to modify the microstructure and mechanical properties of cold sprayed AA2024/Al2O3 metal matrix composites. Scanning electron microscopy and optical microscopy identify that FSP improves notably the refinement and dispersion of Al2O3 particles. Nanoindentation results show that the hardness and elastic modulus of cold sprayed AA2024 particles decrease with FSP, which are mainly attributed to the dissolution of the Guinier Preston Bagaryatskii (GPB) zone. Higher Vickers hardness values are reached for the FSPed specimens due to the uniform distribution and refinement of reinforcing particles. The tensile properties of cold sprayed AA2024/Al2O3 composites are remarkably enhanced by FSP, with an increase of 25.9% in the ultimate tensile strength and 27.4% in elongation. The improved interparticle bonding and the dispersion strengthening of the refined Al2O3 particles are the main strengthening mechanisms responsible for these enhancements. © 2017 Elsevier B.V. All rights reserved.

Keywords: Cold spray Friction stir processing Metal matrix composites Phase transformation Mechanical properties Strengthening mechanisms

1. Introduction Cold spray (CS), also known as cold gas dynamic spray, kinetic spray or supersonic particle deposition, has been attracting more attention as a versatile solid-state coating and additive manufacturing technique [1e3]. With this process, metallic or dielectric substrates are exposed to high velocity particles accelerated by an expanding gas stream at temperatures lower than the melting point of the spray material [1]. These conditions identify CS as a low-temperature and high-velocity process with significant advantages such as lack of oxidation or phase transformation, which enables feedstock powder to retain its original properties [4,5]. Compared to thermal spray (TS), bonding in CS is achieved at a solid state with severe impact plastic deformations rather than melting at very high temperatures [6]. These advantages allow CS to deposit particle reinforced metal matrix composites (MMCs) such as AA5056/SiC [7], Al/Al2O3 [8] and Cu/CNT/SiC [9]. Compared to CS metals and alloys, the presence of ceramic particles in MMCs can increase plastic deformation of the metal matrix while reducing coating porosity, improving hardness, wear- and corrosionresistance [7e10]. In addition, investigations on particle reinforced * Corresponding author. E-mail address: [email protected] (W. Li). https://doi.org/10.1016/j.jallcom.2017.11.132 0925-8388/© 2017 Elsevier B.V. All rights reserved.

MMCs have shown that the strengthening effect is linked to the pinning of matrix dislocations by dispersed reinforcements, also known as dispersion strengthening [11]. However, this effect is significant only when particle size is very small (below 0.1 mm) [12]. Although ceramic particles can be broken by high-velocity impact, the extent of their refinement is quite limited, which makes the presence of ceramic particles a liability for bonding. So a significant loss of strength and ductility of MMCs follows CS. Heat treatment is the most widely used post-treatment to improve CSed coatings by improving bonding between metallic particles [6,13]. However, this effect is limited on MMCs, as ceramic particles retain their large size during heat treatment. This makes it very difficult to improve dispersion strengthening, leaving weak bonding between ceramic and metallic particles in CSed deposits. Therefore, an alternative post-treatment is necessary to strengthen CSed MMCs. As another eco-friendly solid-state process, friction stir processing (FSP) has been used to modify the microstructural and mechanical properties of metals [14,15]. During this process, a rotating cylindrical tool with a shoulder and pin is first plunged into a metallic plate and then traversed along the surface of the workpiece. The rubbing action of the tool shoulder with the plate's surface generates frictional heat and softens the material under the shoulder which undergoes severe plastic deformation at high strain

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rates by the action of the rotating pin [14,15]. FSP has been employed to co-fabricate MMCs to achieve homogeneous microstructure in reinforcement particle's spreading and dispersion with conventional processes including TS, powder metallurgy (PM) and mechanical alloying (MA) [16]. In the recent few years, FSP has been used for CSed MMCs in order to refine the distribution of reinforcing particles and remove the weak interface between particles [17e20]. Hodder et al. [17] successfully applied FSP to CSed Al/Al2O3 composites and studied the resulting refinement and redistribution of particles. Huang et al. [18] and Peat et al. [19,20] carried out comprehensive studies on the wear and erosion performances of CSed MMCs which were modified by FSP, and confirmed the effectiveness of FSP as a surface engineering technique. In addition to these modifications, it has been shown that significant mechanical property enhancements of CSed Cu-Zn alloys can be achieved by post-spray FSP [10]. Therefore, in this study, FSP was employed as a strengthening method to improve CSed AA2024/Al2O3 MMCs. The microstructural evolution including particles refinement and phase transformation during FSP was characterized. The improvement on micromechanical and tensile properties were measured following FSP. Finally, the multiple strengthening mechanisms were identified and quantified. 2. Material and methods 2.1. MMCs deposition The powders used in this study were a mixture of commercially spherical AA2024 particles (45 þ 15 mm, Beijing Xing Rong Yuan Technology Co., Ltd, China) and irregular Al2O3 particles (45 þ 15 mm, Beijing You Xing Lian Technology Co., Ltd, China), with a volume ratio of 4:1. Fig. 1a and b shows the morphologies of the two powders, respectively. Cu plates of 3 mm thickness were used as substrates. A custom developed CS system was used to perform deposition. The nozzle had an expansion ratio of 6.7 and a divergent section length of 200 mm. Helium was used as the accelerating gas at an inlet pressure of 0.8 MPa and temperature of around 500 C. The nozzle standoff distance was 25 mm and the gun traverse speed was set to 20 mm/s. The thickness of AA2024Al2O3 MMCs deposits were about 5 mm.

MMCs. The processing direction is parallel to the spray gun traverse direction as shown in Fig. 2a. The H13 steel tool had a threaded pin with 3.4 mm in root diameter and 2.9 mm in length, as well as a concave shoulder of 10 mm in diameter. The rotating tool was tilted at 2.5 . A macroscopic image showing a cross-section view of a typical FSPed specimen is shown in Fig. 2b. The pin region denotes an area of stir zone (SZ) through which the tool pin traversed. In this paper, investigations have focused on SZ. 2.3. Materials characterization The microstructure was examined with an optical microscope (OM, OLYMPUS GX71, Japan) and a scanning electron microscope (SEM, JSM5800LV, JEOL, Japan). An X-Ray diffractometer system (XRD, Siemens D500, Germany) was used to measure microstrain and crystallite size with 40 kV and 30 mA of Cu Ka radiation at a wavelength of 0.1542 nm. The scan was conducted in 2q mode and spans across a range of 10 e120 at a step resolution of 0.033 were performed. The precipitation state was quantified by differential scanning calorimetry (DSC, Netzsch STA449C, Germany). The DSC measurements were carried out under argon with a heating at 10 C/ min up to 550 C. The weight of each sample was under 20 mg. The coating hardness was tested with a Vickers hardness indenter (Struers Duramin-A300, Denmark) with a load of 2N for 15s. The depth sensing indentation technique was used to determine micromechanical properties, using a Nanoindenter G200 (Agilent Technologies Inc, America) with a constant load at each point (10 mN). Seven points were tested in each specimen, and after ignoring the points with extreme values (maximum and minimum values) five lines are shown. The tensile properties of CSed and FSPed AA2024/Al2O3 MMCs were measured with a tensile test machine (INSTRON-3382, America). Three samples were tested to evaluate the tensile properties for each group. The tensile sample position is also shown in Fig. 2a. A micro-tensile specimen with 2 mm-thickness was used, where the width in the gauge section and total length were 2 and 25 mm, respectively (Fig. 3b). A specially designed fixture (Fig. 3a) was used for these specimens to ensure correct application of the load and prevent secondary bending or any other unwanted movement of the specimens. The fracture surfaces of tensile specimens were also observed with SEM and OM. 3. Results and discussion

2.2. Post-spray FSP treatment 3.1. Microstructural evolution The FSP was carried out on a commercial friction stir welding (FSW) machine (FSW-RL31-010, Beijing FSW Technology Co., Ltd., PR China). The processing speed of 100 mm/min and a rotation speed of 1500 rpm were employed to modify CSed AA2024/Al2O3

3.1.1. Al2O3 particles refinement Fig. 4 shows the OM and SEM images of the CSed and FSPed AA2024/Al2O3 MMCs. The absence of defects in CSed MMCs is

Fig. 1. Morphologies of (a) AA2024 powder, (b) Al2O3 powder.


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Fig. 2. (a) Schematic diagram of FSP and (b) macroscopic cross-section of the FSPed AA2024/Al2O3 MMCs.

leads to a volume fraction of nanoscale Al2O3 particles in the FSPed state of 7.1%. In literature [21,22], the level of refinement has been shown to increase with FSP passes. Therefore, it is expected that additional FSP passes through the MMCs would redistribute the asdeposited reinforcing ceramic particles further even in nanoscale.

Fig. 3. (a) Specially designed fixture for tensile tests and (b) details of micro-tensile specimen.

attributed to the ceramic Al2O3 particles in a ductile metallic binder as shown in Fig. 4a. Following impact during CS, a few Al2O3 particles are broken into smaller sized particles, and this refinement is very limited. After FSP, ceramic particles are refined markedly and the AA2024 interfaces disappear as marked in Fig. 4b. When compared to the CSed state, the level of separation and dispersion of Al2O3 in the AA2024 matrix increases after FSP and few ceramic particles retain their initial large size. This refinement is due to the shear forces that exerted by the tool as it stirs the materials [20]. Although a homogeneous MMCs deposit with broken ceramic particles is fabricated with FSP, a large number of Al2O3 particles can still be seen (~>1 mm). This indicates an insufficient refinement action which limits the dispersion strengthening effect [12]. By taking five OM images of each case, the volume fractions of Al2O3 in CSed and FSPed MMCs were counted. It is assumed that there are no nanoscale Al2O3 particles in the CSed state and the volume fraction of ceramics in the two states remains the same. Therefore, the volume fraction of nanoscale Al2O3 particles in the FSPed state, which contributes to dispersion strengthening, can be obtained by subtracting the total volume from the visible micro-sized fraction. Statistical analysis shows that the volume fraction of Al2O3 particles in the CSed state is 28.4%, while it is 21.3% in the FSPed state, which

3.1.2. XRD analysis XRD profiles obtained from the CSed and FSPed MMCs specimens are shown in Fig. 5a. Compared to the CSed state, some peaks corresponding to Al2O3 particles in the FSPed state become a little shorter. This could be due to a reduction in size of Al2O3 and a higher level of separation and distribution of Al2O3 particles [23]. In addition, this could be attributed to an insufficient refinement action. There is a significant reduction in the width of peak of the Al phase indicating the presence of relatively large crystallites and the development of microstrain in the FSPed state. The WilliamsonHall method was used to extract the average crystallite size, amount of microstrain and dislocation density of Al phase [24]. The Williamson-Hall equation is written in the form of

bcosq 2εsinq 0:9 ¼ þ D l l

(1)

ε2 b2

(2)

r ¼ 14:4

where b is the full width at half maximum in rad, l is the wavelength of X-ray beam (0.1542 nm for CuKa), D is the crystallite size, ε is the lattice strain, q is the Bragg angle, r is the dislocation density q and b is the Burgers vector (0.286 nm for Al). By plotting bcos l 2ε 0:9 against sinq, the slope of the graph gives l , the Y-intercept gives D and r is calculated using Eq. (2). The estimated values based on this methodology are shown in Fig. 5b. The large values of microstrain and dislocation density in the CSed MMCs are due to the very high plastic strain rates developing by high-velocity impact during CS, which can be up to 109s1 [25]. The investigation of Huang et al. [26] shows that the dislocation density in metals would increase exponentially for plastic strain rate over 103s1. For FSPed MMCs, it can be shown that the microstrain and dislocation density decrease when compared to the CSed state. This is due to the dynamic recrystallization caused

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Fig. 4. OM (a, b) and SEM (c, d) micrographs of CSed (a, c) and FSPed (b, d) AA2024/Al2O3 MMCs.

Fig. 5. (a) XRD spectra of the CSed and FSPed AA2024/Al2O3 MMCs and (b) Williamson-Hall plot for dislocation density, microstrain and crystallite size.

by stirring during FSP. During CS of Al alloy, the only recovery that is taking place is due to the high stacking fault energy without recrystallization taking place [27]. While in the FSP case, dynamic recrystallization does develop because of the severe plastic deformation and heat production caused by stirring, which has been confirmed by numerous studies [28e30]. The crystallite size (D) calculated corresponds to the dislocation cell size (defect free regions bounded by dislocation walls) [31]. The little larger crystallite size in the FSPed state is due to the lower dislocation density, which is caused by dynamic recrystallization. As crystallite size often matches grain size, there are exceptions especially for the severely deformed materials [31], which is the

case for both processes studied here. This case means that the calculated crystallite size of AA2024 is not equal to grain size here. 3.1.3. DSC analysis The reactions such as exo- and endo-thermic were assessed using DSC to examine dissolution and precipitation behaviors of their precipitation as a function of heating rate. As AA2024 consists of a Guinier Preston Bagaryatskii (GPB) zone and a S(S0 ) phase [28,32,33]. The metastable S0 phase, which forms at higher temperatures, is described as a slightly strained version of the S phase and on a DSC thermogram it is not possible to separate the two phases. As a consequence, phase S(S0 ) is reported [33]. Fig. 6 shows


2 1 n2 H Ecot a arctanh Y¼ Ecot a 2 1 n2 Cq

119

(7)

h i where Cq ¼ p2ffiffiffi 2:845 0:023757 p2 a ; n is Poisson's ratio of 3

material (0.33); a can be assumed equal to 70.3 for a Berkovich indenter; E, H are the elastic modulus and hardness, respectively. The calculated stresses of the CSed and FSPed AA2024 particles are about 847 MPa and 736 MPa, respectively. It can be seen that yield stress also decreases after FSP. In order to explain this phenomenon, the yield stress (Y) can be expressed as [33,36]:

Y ¼ s0 þ sHP þ sss þ

Fig. 6. DSC thermograms of the CSed and FSPed AA2024/Al2O3 MMCs.

the regions of dissolution of GPB zone (Peak A) and formation/ growth of S(S0 ) phase (Peak B). We assume that the peak area A of the CSed state is A1 with a volume fraction of fGPBA1 , and the peak area B1 of the CSed state is B1 with a volume fraction of fSðS0 ÞB1 . According to Ref. [33], the relative fraction of GPB zone for the FSPed state can be calculated as a ratio between the GPB zone peak area A2 of the FSPed to that of the CSed:

fGPBA2 ¼

A2 f A1 GPBA1

(3)

Because the CSed state is different from T351 state which consists of solely of the GPB zone, the calculation of relative fraction of S(S0 ) precipitates in the FSPed is changed accordingly [33]:

fSðS0 ÞB2 ¼ 1

B2 B2 þ f 0 B1 B1 SðS ÞB1

(4)

Following these equations:

fGPBA2 ¼ 0:65fGPBA1

(5)

fSðS0 ÞB2 ¼ 4:21fSðS0 ÞB1 3:21

(6)

When compared to the CSed state, it is apparent that the GPB zone becomes reduced. During FSP, temperatures reached are sufficiently high to cause the dissolution of the GPB zone and the formation and coarsening of the S(S0 ) phase [28,33]. Investigations by Zhang et al. [28] shows that the GPB zone of AA2024 plays an important role in improving properties due to the extensive presence of the coarse S(S0 ) phase. Therefore, it can be predicted that the mechanical properties of the CSed AA2024 particles reduce after FSP, which will be discussed in the next section.

3.2. Mechanical evolution 3.2.1. Micromechanical properties Nanoindentation was used to determine the micromechanical properties of the CSed and FSPed AA2024 particles in MMCs. Fig. 7 shows the load with depth curves in two states with the mean elastic modulus and hardness. The plastic deformation experienced by AA2024 particles during CS results in very large hardness and modulus values [34]. In order to calculate the yield strengths from indentation testing, the model proposed by Yu and Blanchard [35] is adopted:

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi s2GPB þ s2SðS0 Þ þ sdis

(8)

where s0 is base strength for aluminum; sHP is grain size contribution according to the Hall-Petch law; sss is the solid solution contribution; sGPB is hardening due to the GPB zone; sSðS0 Þ is hardening due to the S(S0 ) precipitate. During FSP, dislocation density decreases and grains become refined due to dynamic recrystallization with the reduction of sdis and increase of sHP . According to Jones et al. [37] and Zhang et al. [38], the GPB zone dissolves with the coarsening of the S(S0 ) precipitate which results in reducing hardening values. These reasons attribute to the reduction of yield stress with FSP. The existing studies suggest that for age-hardening aluminum alloys, yield stress is more dependent on precipitation hardening rather than grain size, unless nanograins develop with quick cooling [33,39]. In addition, it can be seen that hardness and elastic modulus also decrease with FSP. In a similar fashion, these phenomena are the result of the removal of work-hardening, the dissolution of the GPB zone and coarsening of the S(S0 ) precipitate. If the calculated characteristic strength Y is assumed to be equal to the yield strength of the material, then these would probably be unrealistically large. For example, AA2024 typically has yield strength under 350 MPa, which is much lower than those calculated here [28]. Extensively cold-worked AA2024 sheets possess a yield stress of approximately 550 MPa, whose elongation to fracture is reduced from 20% to 30% to under 3% [40]. Because of the severe plastic deformation during CS, the present deposits would be so embrittled that their elongation to fracture would be very low, which is inconsistent with their rather ductile response observed during indentation tests. The investigations on Ta by Bolelli et al. [41] show the same phenomenon. Therefore, it is more appropriate to follow the literature [41,42], where Y is a characteristic stress associated to a representative strain (8%). Although Y is a characteristic stress, it reflects the actual relationship between different yield stress values [41]. Therefore, it would be used in Section 3.3 to evaluate the strength contribution of interfaces between particles. The influence of FSP on Vickers hardness of CSed AA2024/Al2O3 MMCs is shown in Fig. 8. It is shown that FSP leads to a dramatic increase in hardening with microhardness increasing from 125HV to 145HV, compared to the CSed deposits. This result brings about an inconsistency with nanoindentation. Compared to nanoindentation, the scope of Vickers hardness test is relatively much large. Therefore, the increase of Vickers hardness is due to the uniform distribution and refinement of reinforced Al2O3 particles in the FSPed deposits. This effect has also been reported by Hodder et al. [17] and Huang et al. [18], which illustrated that FSP can increase the microhardness of CSed MMCs effectively through dispersing and decreasing the mean free path of reinforced ceramic particles.

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Fig. 7. Load vs. depth curves obtained from nanoindentation of AA2024 particles: (a) CSed and (b) FSPed states. Note: different lines in each sub-figure mean different indentations for one specimen.

Fig. 8. Vickers hardness values of AA2024/Al2O3 MMCs.

3.2.2. Tensile properties Fig. 9a shows the tensile test results of CSed and FSPed AA2024/ Al2O3 MMCs. It can be confirmed that FSP has a great improvement effect on CSed AA2024/Al2O3 composites, with a remarkable increase in UTS and elongation. This can be the following reasons: (1) reduced porosity [43]; (2) improved bonding between Al2O3/ AA2024 and AA2024/AA2024 particles which improves interfacial bonding [22,44]; (3) higher level of separation and dispersion of Al2O3 particles which reduces the probability of agglomeration of Al2O3 and crack nucleation [45]; (4) dispersion strengthening attributed to nanoscale Al2O3 particles; (5) grain refinement due to dynamic recrystallization occurring during FSP. Barmouz et al. [22] found that multi-pass FSP could effectively increase tensile properties of the Cu/SiC composites. Therefore, it is proposed that multipass FSP will further strengthen the CSed AA2024/Al2O3 MMCs. Fig. 10 shows the SEM images of fracture surfaces of CSed and FSPed composites. It is clearly seen that there exist pores and pulled-out largescale Al2O3 particles in the fracture surface of the CSed state as marked in Fig. 10b. These are the main reasons for the relatively low tensile property by causing crack nucleation and propagation. No dimples can be found in Fig. 10a and b with

Fig. 9. (a) Stess-elongation curves and (b) UTS and elongations with error bars of the CSed and FSPed AA2024/Al2O3 MMCs.

fracture taking place along interfaces between the deposited particles rather than through them, indicating a brittle rupture for the CSed state. Compared to the CSed state, there are a number of voids and dimples sizes in the FSPed case, which means a ductile fracture (Fig. 10d). However, there remain a few unbroken Al2O3 particles in the FSPed state marked in Fig. 10d. This could be a hint for the formation of weak bonding between Al2O3 and AA2024 matrix. Huang et al. [10] reported a large increase of 195.3% for UTS and 376.5% for elongation of the CSed Cu-Zn alloy after FSP. In this work, the presence of large ceramic particles is the main reason, which hinders a better tensile property of AA2024/Al2O3 MMCs. The three-dimensional topographies of fracture surfaces are shown in Fig. 11. Compared to the CSed state, a relatively rough surface was produced during fracture, which leads to a higher tensile property.

3.3. Discussion on strengthening mechanisms Several additional laws are available for estimating the total


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Fig. 10. SEM micrographs of fracture surface of tensile specimens for (a, b) CSed and (c, d) FSPed states. (b, d) are higher magnification images of (a, c), respectively.

Fig. 11. Three-dimensional topographies of fracture surfaces of tensile specimens for (a) CSed and (b) FSPed states.

strengthening [33,36]. In our case, a combined law, linear and Pythagorean, is suggested. The model proposed for the CSed state is:

sFSP ¼ s0 þ sHP þ sss þ

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi s2GPB þ s2SðS0 Þ þ sdis þ sOR sPIFSP (10)

sCS ¼ s0 þ sHP

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi þ sss þ s2GPB þ s2SðS0 Þ þ sdis sPICS

(9)

where sPICS is the special existence in CS, which means a negative contribution to the final tensile strength due to the interfaces between particles [6]. For CS, the existence of interfaces between the particles is the main reason for the poor performance of CSed deposits compared to the bulk material, namely those of AA2024/ Al2O3 and AA2024/AA2024. As the broken extent of Al2O3 particles is very limited during CS, dispersion strengthening can be ignored. For the FSPed state, the equation for total tensile strength becomes:

FSP can decrease interparticle spacing, with promoting bonding. So, the negative contribution to the final strength due to the interface is limited after FSP (sPIFSP ). During FSP, some Al2O3 particles are refined to nanoscale, which can influence dispersion strengthening or Orowan strengthening (sOR ). sOR can be calculated with the modified Orowan formula [46,47]:

sOR ¼

4 mGb qffiffiffiffiffi ln 2b p 1 1:18 2p 4 6f v

(11)

122


where m is taylor factor (approximately 3); G is shear modulus for aluminum (26.2 GPa); 4 is the average nanoscale Al2O3 size. In this study, particles size below 1 mm would be counted as nanoparticles. Therefore, 0.5 mm is considered as the average value for calculation; fv is the volume fraction of nanoscale Al2O3 particles (7.1%). From Eqs. (8e11), the contribution of dispersion strengthening (sOR ) is found to be 24 MPa. After FSP, the decrease of negative contribution due to interfaces (sPICS sPIFSP ) is about 150 MPa. From ‘Section 3.2.1’, it can be shown that the strength of CSed AA2024 matrix decreases after FSP, which can be due to the remove of work-hardening, dissolution of the GPB zone and coarsening of the S(S0 ) precipitate. Therefore, it can be concluded that the strengthening mechanisms are mainly through improving interparticle bonding due to decreased interparticle spacing and dispersion strengthening attributed to refined Al2O3 particles. It should be noted that although the negative contribution of sPI can be reduced with FSP, the existence of weak interparticle interfaces remains the main factor hampering the improvement of tensile property. According to literature [16,22], the higher ratio of tool rotation speed to traverse speed and number of passes are the means to further strengthen CSed composites. 4. Conclusions In this research, cold sprayed AA2024/Al2O3 MMCs were modified with FSP. The microstructural and mechanical evolutions in the CSed and FSPed conditions were investigated. The strengthening mechanisms were identified and quantified. The main conclusions are as follows: (1) FSP refines the Al2O3 particles and improves the dispersion of Al2O3 particles in AA2024 matrix with severe stirring action. Meanwhile, the dissolution of the GPB zone and formation and coarsening of the S(S0 ) phase happen during FSP. (2) After FSP, the hardness and elastic modulus of CSed AA2024 particles decrease. While the Vickers hardness of deposited MMCs increases due to the refinement of Al2O3 particles. (3) Tensile properties of CSed AA2024/Al2O3 MMCs are significantly enhanced with FSP, making FSP an effective process to strengthen CSed composites. (4) The improved interparticle bonding due to decreased interparticle spacing and dispersion strengthening attributed to refined Al2O3 particles are the main strengthening mechanisms. Therefore, weakening interparticle interfaces could further improve the tensile performance of CSed MMCs. Acknowledgements The authors would like to thank the financial support of the National Key Research and Development Program of China (2016YFB1100104), the National Natural Science Foundation of China (51574196), the fund of SAST (SAST2016043) and the 111 Project (B08040). References €rtner, T. Klassen, Cold spraying e a materials [1] H. Assadi, H. Kreye, F. Ga perspective, Acta Mater. 116 (2016) 382e407. [2] C. Lee, J. Kim, Microstructure of kinetic spray coatings: a review, J. Therm. Spray Technol. 24 (2015) 592e610. [3] R. Jones, N. Matthews, C.A. Rodopoulos, K. Cairns, S. Pitt, On the use of supersonic particle deposition to restore the structural integrity of damaged aircraft structures, Int. J. Fatig. 33 (2011) 1257e1267. [4] S. Yin, Y.C. Xie, J. Cizek, E.J. Ekoi, T. Hussain, D.P. Dowling, R. Lupoi, Advanced diamond-reinforced metal matrix composites via cold spray: properties and deposition mechanism, Compos. Part B Eng. 113 (2017) 44e54.

[5] G.J. Yang, C.J. Li, F. Han, W.Y. Li, A. Ohmori, Low temperature deposition and characterization of TiO2 photocatalytic film through cold spray, Appl. Surf. Sci. 254 (2008) 3979e3982. [6] K. Yang, W.Y. Li, X.P. Guo, X.W. Yang, Y.X. Xu, Characterizations and anisotropy of cold-spraying additive-manufactured copper bulk, J. Mater. Sci. Technol.. In press. [7] M. Yu, X. Suo, W. Li, Y. Wang, H. Liao, Microstructure, mechanical property and wear performance of cold sprayed Al5056/SiCp composite coatings: effect of reinforcement content, Appl. Surf. Sci. 289 (2014) 188e196. [8] J. Shockley, S. Descartes, P. Vo, E. Irissou, R.R. Chromik, The influence of Al2O3 particle morphology on the coating formation and dry sliding wear behavior of cold sprayed Al-Al2O3 composites, Surf. Coat. Technol. 270 (2015) 324e333. [9] E. Pialago, O. Kwon, C.W. Park, Cold spray deposition of mechanically alloyed ternary Cu-CNT-SiC composite powders, Ceram. Int. 41 (2015) 6764e6775. [10] C.J. Huang, W.Y. Li, Y. Feng, Y.C. Xie, M.P. Planche, H.L. Liao, G. Montavon, Microstructure evolution and mechanical properties enhancement of a coldsprayed Cu-Zn alloy coating with friction stir processing, Mater. Charact. 125 (2017) 76e82. [11] Z. Zhang, D.L. Chen, Contribution of Orowan strengthening effect in particulate-reinforced metal matrix nanocomposites, Mater. Sci. Eng. A 483e484 (2008) 148e152. [12] A.M. Redsten, E.M. Klier, A.M. Brown, D.C. Dunand, Mechanical properties and microstructure of cast oxide-dispersion-strengthened aluminum, Mater. Sci. Eng. A 201 (1995) 88e102. [13] R.Z. Huang, M. Sone, W.H. Ma, H. Fukanuma, The effects of heat treatment on the mechanical properties of cold-sprayed coatings, Surf. Coat. Technol. 261 (2015) 278e288. [14] Z.Y. Ma, Friction stir processing technology: a review, Metall. Mater. Trans. A 39A (2008) 642e658. [15] R.S. Mishra, Z.Y. Ma, I. Charit, Friction stir processing: a novel technique for fabrication of surface composite, Mater. Sci. Eng. A 341 (2003) 307e310. [16] N. Gangil, A.N. Siddiquee, S. Maheshwari, Aluminium based in-situ composite fabrication through friction stir processing: a review, J. Alloys Compd. 715 (2017) 91e104. [17] K.J. Hodder, H. Izadi, A.G. McDonald, A.P. Gerlich, Fabrication of aluminumalumina metal matrix composites via cold gas dynamic spraying at low pressure followed by friction stir processing, Mater. Sci. Eng. A 556 (2012) 114e121. [18] C.J. Huang, W.Y. Li, M. Yu, M. Planche, H. Liao, G. Montavon, Modification of a cold sprayed SiCp/Al5056 composite coating by friction stir processing, Surf. Coat. Technol. 296 (2016) 69e75. [19] T. Peat, A. Galloway, A. Toumpis, R. Steel, W.Z. Zhu, N. Iqbal, Enhanced erosion performance of cold spray co-deposited AISI316 MMCs modified by friction stir processing, Mater. Des. 120 (2017) 22e35. [20] T. Peat, A. Galloway, A. Toumpis, P. McNutt, N. Iqbal, The erosion performance of cold spray deposited metal matrix composite coatings with subsequent friction stir processing, Appl. Surf. Sci. 396 (2017) 1635e1648. [21] Z.Y. Ma, S.R. Sharma, R.S. Mishra, Effect of multiple-pass friction stir processing on microstructure and tensile properties of a cast aluminum-silicon alloy, Scr. Mater. 54 (2006) 1623e1626. [22] M. Barouz, M.K.B. Givi, Fabrication of in situ Cu/SiC composites using multipass friction stir processing: evaluation of microstructural, porosity, mechanical and electrical behavior, Compos. Part A Appl. S 42 (2011) 1445e1453. [23] K. Surekha, B.S. Murty, R.K. Prasad, Microstructural characterization and corrosion behavior of multipass friction stir processed AA2219 aluminum alloy, Surf. Coat. Technol. 202 (2008) 4057e4068. [24] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and wolfram, Acta Metall. 1 (1953) 22e31. €rtner, T. Stoltenhoff, H. Kreye, Bonding mechanism in cold gas [25] H. Assadi, F. Ga spraying, Acta Mater. 51 (2003) 4379e4394. [26] M.X. Huang, E.J. Pedro Rivera-Díaz-del-Castillo, B. Olivier, A constitutive model for high strain rate deformation in FCC metals based on irreversible thermodynamics, Mech. Mater. 41 (2009) 982e988. €rtner, T. Stoltenhoff, H. Kreye, Microstructural bonding fea[27] C. Borchers, F. Ga tures of cold sprayed face centered cubic metals, J. Appl. Phys. 96 (2004) 4288e4292. [28] Z.H. Zhang, W.Y. Li, Y. Feng, J.L. Li, Y.J. Chao, Global anisotropic response of friction stir welded 2024 aluminum sheets, Acta Mater. 92 (2015) 117e125. [29] A.K. Shukla, W.A. Baeslack, Orientation relationships and morphology of S phase in friction stir welded Al-Cu-Mg alloy, J. Mater. Sci. 44 (2009) 676e679. [30] Y.S. Sato, H. Kokawa, M. Enomoto, Microstructure evolution of 6063 aluminum during friction-stir welding, Metall. Mater. Trans. A 30 (1999) 2429e2437. [31] A. Chaudhuri, Y. Raghupathy, D. Srinivasan, S. Suwas, C. Srivastava, Microstructural evolution of cold-sprayed Inconel 625 superalloy coatings on low alloy steel substrate, Acta Mater. 129 (2017) 11e25. gue, A. Deschamps, W.J. Poole, On the coupling between [32] C. Genevois, D. Fabre precipitation and plastic deformation relation with friction stir welding of AA2024 T3 aluminium alloy, Mater. Sci. Eng. A 441 (2006) 39e48. [33] C. Genevois, A. Deschamps, A. Denquin, B. Doi, Quantitative investigation of precipitation and mechanical behavior for AA2024 friction stir welds, Acta Mater. 53 (2005), 2447e2548. [34] S.R. Bakshi, V. Singh, K. Balani, D.G. McCartney, S. Seal, A. Agarwal, Carbon nanotube reinforced aluminum composite coating via cold spraying, Surf. Coat. Technol. 202 (2008) 5162e5169.

K. Yang et al. / Journal of Alloys and Compounds 736 (2018) 115e123 [35] W.P. Yu, J.P. Blanchard, An elastic-plastic indentation model and its solutions, J. Mater. Res. 11 (1996) 2358e2367. [36] H.J. Liu, Y.Y. Hu, C. Dou, D.P. Sekulic, An effect of the rotation speed on microstructure and mechanical properties of the friction stir welded 2060-T8 Al-Li alloy, Mater. Charact. 123 (2017) 9e19. [37] M.J. Jones, P. Heurtier, C. Desrayaud, F. Montheillet, D. Allehaux, J.H. Driver, Correlation between microstructure and microhardness in a friction stir welded 2024 aluminum alloy, Scr. Mater. 52 (2005) 693e697. [38] Z. Zhang, B.L. Xiao, Z.Y. Ma, Hardness recovery mechanism in the heat-affected zone during long-term natural aging and its influence on the mechanical properties and fracture behavior of friction stir welded 2024Al-T351 joints, Acta Mater. 73 (2014) 227e239. [39] X.L. Feng, H.J. Liu, S.S. Babu, Effect of grain size refinement and precipitation reactions on strengthening in friction stir processed Al-Cu alloys, Scr. Mater. 65 (2011) 1057e1060. [40] M. Mirzaei, M.R. Roshan, S.A. Jenabali Jahromi, Microstructure and mechanical properties relation in cold rolled Al 2024 alloy determined by X-ray line profile analysis, Mater. Sci. Eng. A 620 (2015) 44e49.

123

[41] G. Bolelli, B. Bonferroni, H. Koivuluoto, L. Lusvarghi, P. Vuoristo, Depth-sensing indentation for assessing the mechanical properties of cold-sprayed Ta, Surf. Coat. Technol. 205 (2010) 2209e2217. [42] D. Tabor, The Hardness of Metals, Oxford University Press, Oxford (UK), 1951, pp. 105e106. [43] C. Tekmen, I. Ozdemir, U. Cocen, The mechanical response of Al/Si/Mg/SiCp composite: influence of porosity, Mater. Sci. Eng. A 360 (2003) 365e371. [44] I.S. Lee, P.W. Kao, N.J. Ho, Microstructure and mechanical properties of Al-Fe in situ Nanocomposite produced by friction stir processing, Intermetallics 16 (2008) 1104e1108. [45] K. Nakata, Y.G. Kima, H. Fujii, T. Tsumura, T. Komazaki, Improvement of mechanical properties of aluminum die casting alloy by multi-pass friction stir processing, Mater. Sci. Eng. A 437 (2006) 274e280. oz-Morris, O.C. Garcia, D.G. Morris, An analysis of strengthening [46] M.A. Mun mechanisms in a mechanically alloyed, oxide dispersion strengthened iron aluminide intermetallic, Acta Mater. 50 (2002) 2825e2836. [47] X.T. Luo, G.J. Yang, C.J. Li, Multiple strengthening mechanisms of cold sprayed cBNp/NiCrAl composite coating, Surf. Coat. Technol. 205 (2011) 4808e4813.