Microstructure, residual stress and tensile properties control of wire-arc additive manufactured 2319 aluminum alloy with laser shock peening

Accepted Manuscript Microstructure, residual stress and tensile properties control of wire-arc additive manufactured 2319 aluminum alloy with laser shock peening Rujian Sun, Liuhe Li, Ying Zhu, Wei Guo, Peng Peng, Baoqiang Cong, Jianfei Sun, Zhigang Che, Bo Li, Chao Guo, Lei Liu PII:

S0925-8388(18)30850-8

DOI:

10.1016/j.jallcom.2018.02.353

Reference:

JALCOM 45229

To appear in:

Journal of Alloys and Compounds

Received Date: 30 November 2017 Revised Date:

10 February 2018

Accepted Date: 28 February 2018

Please cite this article as: R. Sun, L. Li, Y. Zhu, W. Guo, P. Peng, B. Cong, J. Sun, Z. Che, B. Li, C. Guo, L. Liu, Microstructure, residual stress and tensile properties control of wire-arc additive manufactured 2319 aluminum alloy with laser shock peening, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.02.353. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Microstructure, residual stress and tensile properties control of wire-arc additive manufactured 2319 aluminum alloy with laser shock

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peening

Rujian Sun 1, Liuhe Li 1, Ying Zhu 1, Wei Guo 1*, Peng Peng 1, Baoqiang Cong 1, Jianfei Sun 1, Zhigang Che 2, Bo Li 3, Chao Guo 3, Lei Liu 3

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School of Mechanical Engineering and Automation, Beihang University, Beijing, 100191, China

Key Laboratory For High Energy Density Beam Processing Technology, Beijing Aeronautical

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Manufacturing Technology Research Institute, Beijing, 100024, China 3

AVIC The First Aircraft Institute, Xi’an 710089, China

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Corresponding author. Tel: +86 10 82317712; E-mail: [email protected] (Wei Guo)

Abstract: Wire-arc additive manufacturing can fabricate components with complex geometries efficiently compared with other manufacturing methods. However, the uncontrolled grain size and

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tensile residual stress in as-fabricated components have limited their applications. In this study,

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laser shock peening, an innovative surface treatment technique, was specially-combined with wire-arc additive manufacturing to refine microstructure, modify stress state and enhance tensile properties of as-printed 2319 aluminum alloy. After peening, the average grain size decreased from 59.7 µm to 46.7 µm, and the percentage of grains with low angle boundaries increased from 34% to 70%. High density of dislocations and mechanical twins were generated and resulted in the increase of micro-hardness. Residual stresses were modified from tensile to compressive state with a maximum value around 100 MPa. Yield strength was remarkably increased by 72%. This combined

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ACCEPTED MANUSCRIPT printing and peening manufacturing strategy provides microstructure and quality control of manufactured components for practical applications.

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Keywords: Laser shock peening; Wire-arc additive manufacturing; Microstructure; Residual stress; Tensile properties

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1. Introduction

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Aluminum alloy has a wide application and an irreplaceable position in aviation, aerospace, shipbuilding and nuclear industry due to its low density, high specific strength, good formability and low cost, etc. [1]. In recent decades, additive manufacturing (AM), also known as 3D printing, has emerged as an advanced alternative for manufacturing industry. It is a key technique of great

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potential in producing complex geometries close to their net-shape [2-4]. AM techniques can be classified into different categories according to feeding technique, heat source or feedstock material [2]. Wire arc additive manufacturing (WAAM) feeds a wire at a controlled speed into a plasma arc

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to melt the wire onto a substrate or the previously deposited layer, being developed for

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layer-by-layer manufacturing of functional metal with full density in aerospace and other industrial sectors [5]. Comparing with the powder based AM processes, WAAM has much higher material deposition rate and lower cost [6, 7]. However, it should also be noted that several problems need to be resolved to widen the adoption of WAAM in diverse industries, such as lack of fusion, uncontrolled grain size, tensile residual stress, cracks and delaminations [8, 9]. Among these defects, uncontrolled grain size and process induced tensile residual stress are the most frequently reported in many WAAM manufactured metallic pieces [10-12], since varying grain size can affect

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ACCEPTED MANUSCRIPT microstructure uniformity and tensile residual stress can weaken mechanical behaviors. To solve this problem, two types of typical attempt can be taken into consideration. First, AM parameter optimum. Vandenbroucke and Kruth pointed out that Ti-6Al-4V part with higher density can be

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achieved by increasing the applied laser energy [13]. Moat et al. systematically studied different AM process parameters to eliminate tensile residual stress [14]. Monroy investigated defect generation mechanisms based on surface characteristics [11]. Zhang et al. found porous structure

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with regular and rectangular pores can be tailored by adjusting the scan line spacing [15]. Second,

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synchronized or post processing. Colegrove et al. combined a high-pressure rolling together with WAAM, and reported that peak residual stress was reduced and refined microstructure was obtained [16]. Almangour and Yang proposed that severe plastic deformation induced by a shot-peening treatment can effectively enhances the roughness, hardness, compressive yield strength, and wear

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resistance of the direct laser sintered 17-4 stainless steel [17]. Leuders et al. discovered that the combination of hot isostatic pressing and shot peening was a suitable post-treatment for SLM-processed Ti-6Al-4V, allowing elimination of various kinds of process-induced defects [18].

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Zhang et al. evaluated fatigue crack propagation behavior in WAAMed Ti-6Al-4V using numerical

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simulation method [19]. From above open literature, it is believed that synchronized or post processing has a much better improvement in microstructure and mechanical properties of AMed components. Efforts are still in demand in this area to fulfill the gap between experiment study and industrial application.

Laser shock peening (LSP) is an innovative surface treatment technique, which applies a pulse laser with high power density (in GW/cm2 range) and ultra-short duration (usually nanoseconds) onto the surface of metallic components to induce large depth and high value compressive residual stress [20,

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ACCEPTED MANUSCRIPT 21]. Grain refinement can be achieved after multiple LSP treatments [22]. Enhancements in fatigue strength [23, 24], tensile strength [25], wear resistance [26] and corrosion resistance [27, 28] have been intensively reported. Applying LSP on WAAMed components is likely to modify surface

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stress state from tensile to compressive residual stress, and further enhance the fatigue resistance and tensile properties. In addition, with laser shock wave induced on the surface and transmitted within the substrate, coarse grain can be largely refined. Therefore, it is of great value to study the

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effect LSP on microstructure, residual stress and mechanical properties of additively manufactured

2. Experiments 2.1 Materials and WAAM process

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components to superiorly promote the material for a wider range of applications.

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In this study, 2319 aluminum specimens were prepared by WAAM, and their chemical compositions are listed in Table 1. The operating parameters of producing WAAMed aluminum specimens are given in Table 2. Moreover, 2A12 aluminum plates with a thickness of 12 mm were

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applied as substrates for the layer-by-layer deposition.

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Table 1 Chemical compositions of 2A12 aluminum substrate and 2319 aluminum wire (wt %). Elements

Cu

Mn

Zr

Si

Mg

Zn

Ti

Al

Substrate

3.9-4.8

0.3-0.9

≤0.3

≤0.5

1.2-1.8

≤0.3

≤0.15

Bal.

Wire

5.96

0.3

0.12

0.04

-

0.1

0.17

Bal.

Table 2 Parameters used for WAAM experiment. WAAM parameters 4

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Electrode

Tungsten, 3 mm diameter

Wire diameter

1.2 mm

Wire feed speed

1.5 m/min

Shielding gas

Argon, 99.99% purity

Shielding gas flow rate

15 L/min

Travel speed

0.3 m/min

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Welder type

Fig. 1a illustrates a WAAM setup. Each deposition is created by moving the torch in a linear direction and feeding wire at a controlled speed into the molten pool, which solidifies to build a layer. A subsequent layer is then deposited over the former deposited ones by increasing the height

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of the torch. Fig. 1b shows the as-printed wall after 90 layers of deposition with a height around 105 mm. The wall was then prepared by double-sided milling to a 3 mm thick plate. LSP experiments

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were carried out shortly after the preparation.

Fig. 1 (a) Schematic diagram of layer-by-layer WAAM deposition and (b) image of additively manufactured 2319 aluminum specimen and the substrate.

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ACCEPTED MANUSCRIPT 2.2 LSP procedure Schematic diagram of LSP is detailed in Fig. 2a. The LSP experiments were carried out by a Q switched Nd: YAG high power pulse laser with a wavelength of 1064 nm, a pulse duration of 15 ns,

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a circular spot size of 4 mm, an overlapping rate of 50% and a pulse energy of 15 J. Thus, the peak power density of this laser equals to 7.95 GW/cm2. Running water layer with a thickness about 2 mm was used as the transparent layer to increase the peak pressure of laser shock wave. To avoid

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possible damage or roughening of the specimen surface by laser irradiation and to reduce the

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reflection loss of incident laser light, a protective coating, usually aluminum foil or black tape, is typically attached on the specimen surface. Hence, a 3M aluminum foil with a thickness of 100 µm was used as absorbing layer. The specimen was double side peened with a zigzag laser path covering an area of 60 mm × 105 mm on each side, shown in Fig. 2b. The LSP treated specimen is

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shown in Fig. 2c.

Fig. 2 (a) Schematic diagram of LSP, (b) laser path illustrating an overlapping rate of 50% and (c) image of laser shock peened WAAM 2319 aluminum specimen.

2.3 Material characterization Electron backscattered diffraction (EBSD), integrated in a scanning auger nanoprobe (AES, PHI710) was adopted to study the grain refinement, misorientation and texture of specimens before and after

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ACCEPTED MANUSCRIPT LSP with a detection area of 260 µm × 260 µm and a step size of 1 µm. Transmission electron microscope (TEM, JEM-2100, JEOL) was employed to investigate microstructure evolution in specimens before and after LSP. Two LSPed TEM observation thin discs were prepared from layers

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in a depth of around 10 µm and 500 µm below peening surface. Micro-vickers hardness tester (FM-800) was used to study the evolution of micro-hardness in depth direction before and after LSP with a test weight of 200 g and dwell time of 10 s. It was measured

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from the peened surface to a depth with 1.5 mm with a step distance of 0.1 mm. To eliminate the

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deviation, each measurement was conducted for three times and the average values were presented in this study.

Prism residual stress measurement system based on incremental hole drilling method was applied to measure residual in the depth direction (-Y direction) of specimen before and after LSP. It is a

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destructive residual stress measurement. A subsurface residual stress profile is obtained first by bonding a strain gauge rosette to the surface of the specimen. A small drilling rig is then used to drill a blind, flat-bottomed hole and removes materials at incremental depths in the center of the

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strain gauge rosette. With each incremental depth of material that is removed from the hole, a

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change in the strain field occurs accordingly due to the relaxation of the surrounding material. The change in strain is measured by the strain gauges and the relaxation can be correlated with the residual stress in the surface of the material. Residual stresses in the topmost section and middle section were measured from the surface to a depth of 0.75 mm, as illustrated in Fig. 3a. INSTRON 8801 servo hydraulic test system was utilized to perform tensile experiments of specimens before and after LSP at room temperature with a constant strain rate of 10-3 s-1 and a tensile rate of 1.5 mm/min. The 3 mm thickness specimens with a gauge length of 32 mm,

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ACCEPTED MANUSCRIPT displayed in Figs. 3b and 3c, were wire-electrode cut from the as-printed wall. The peening was conducted on double side covering the gauge length region. Three specimens were tested for each condition to obtain reliable results of tensile properties. The yield strength (YS), ultimate tensile

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by scanning electron microscope (SEM, SUPRA55, ZEISS).

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strength (UTS), and ductility (elongation, E) were evaluated, and fracture surface was characterized

Fig. 3 Schematic diagram of (a) cross-section view for micro-hardness and residual stress tests, (b) specimen

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preparation for tensile test and (c) detailed dimension of tensile test specimen.

3. Results and Discussion 3.1 Microstructural evolution 3.1.1 Grain refinement and grain misorientation Fig. 4 illustrates the surface EBSD maps of specimens before and after LSP. Grain size distribution, inverse pole figure, texture and grain misorientation of specimens before and after LSP were investigated. [100] was adopted as the reference direction for the inverse pole figure to code the

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ACCEPTED MANUSCRIPT color pattern (red: [001], blue: [111], green: [101]), as indicated by the triangle in the top right corner of Figs. 4a and 4b. The initial microstructure, displayed in Fig. 4a, was consisted of coarse equiaxed grains. This can be resulted from the latter thermal input for the building of the layer,

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acting as the low-temperature heat treatment for the previously deposited layers [7], and hence, coarse grains were generated. In contrast, refined microstructure was obtained after LSP, as shown in Fig. 4b. Figs. 4c and 4d present histograms of area fraction versus grain size before and after LSP.

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47% of the grains exceeded 60 µm before LSP, while only 24% of the grains exceeded 60 µm after

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LSP. In addition, the average size decreased by 22%, from 59.7 µm to 46.7 µm after LSP. For the texture before and after LSP, both of the specimens exhibited weak texture, as the maximum values were 3.588 and 3.512 (Figs. 4e and 4f). This is because no preferred orientation is generated during WAAM. Antonysamy et al. also observed a weak α-texture resulting from a random

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distribution across habit variants in Ti6Al4V titanium alloy [29]. The misorientation angle is defined as the minimum rotation angle required for two neighboring crystal lattices to coincide and the number fraction at a certain misorientation angle is defined as the total number of pixels along

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all grain boundaries within the analysis area divided by the number of pixels along the boundaries

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within this misorientation angle [30]. For the specimens before LSP, 34% of grain boundaries were found with low misorientation angles (less than 15°, also known as low angle grain boundaries). In particular, 2° misorientation angles of grain boundaries are maximal with 18%. Meanwhile, more than 60% of boundaries have high misorientation angles ranging from 20° to 60°, as shown in Fig. 4g, which illustrates a random distribution of grains. In contrast, nearly 70% of grain boundaries were found with low misorientation angles. Especially for 2° misorientation angles, it has a fraction of 43%, as presented in Fig. 4h. This percentage is approximately 2.5 times higher than that in the

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ACCEPTED MANUSCRIPT original specimen. This distinct difference is due to numerous laser-beam interactions with the material on the surface and the high-energy plasma and consecutive shock waves at the interaction region. The original coarse grains refined into small grains including sub-grains during LSP. The

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new misorientation angles are the small grains from the same coarse grain, and therefore, a large number of misorientations around 2° were achieved [31]. The increase of low angle grain

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boundaries also proves that LSP is an effective method for grain refinement.

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Fig. 4 EBSD analysis of inverse pole figure, grain size, texture and grain boundary misorientation. (a), (c), (e) and (g) for specimen before LSP and (b), (d), (f) and (h) for specimen after LSP.

3.1.2 TEM observation

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ACCEPTED MANUSCRIPT Fig. 5 shows typical TEM images of specimen before LSP. It can be observed from Fig. 5a that a small number of mechanical twins are randomly distributed in the original grains, which can be considered growth twins since no plastic deformation has occurred in the substrate [25]. Fig. 5b is

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the corresponding pattern of selected area electron diffraction (SAED) taken from the ellipse [A] in Fig. 5a, in which [011] is detected as zone axis. Fig. 5d is the inverse fast fourier transform (IFFT) image in crystal orientation of (1 1 1) from high resolution transmission electron microscopy

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(HRTEM) image (Fig. 5c) taken from the ellipse [A]. Interatomic spacing is measured from 10

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atom layers in order to acquire an accurate result. It can be seen that atom layers were uniformly

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distributed with an interatomic spacing being 0.24 nm with uniform profile showing in Fig. 5d.

Fig. 5 Typical TEM images of specimen before LSP. (a) mechanical twins randomly distributed in the original grain, (b) corresponding SAED pattern of Ellipse [A] in (a), (c) HRTEM image, (d) IFFT image from crystal orientation of (1 1 1) , (e) profile of line in IFFT image.

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ACCEPTED MANUSCRIPT Typical TEM images in the surface layer of specimen after LSP are illustrated in Fig. 6. The surface layer is considered severe plastic deformation layer. Dislocation wall and dislocation tangles were clearly observed. However, no obvious dislocation lines detected because the surface layer was so

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severely deformed during the LSP process that it was difficult to find a zone with clear dislocation line structures [32]. Sub-grain was also obtained, surrounded by high-density dislocations. Fig. 6b is the corresponding SAED pattern captured from the ellipse [B] in Fig. 6a with [011] being the

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direction of crystal axis. Fig. 6d presents the IFFT image in crystal orientation of (1 1 1) from

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HRTEM image (Fig. 6c) taken from the ellipse [B] in Fig. 6a. Interatomic spacings of line I and II illustrated different atomic arrangement in the sub-grain and in the dislocation tangles, respectively. In the sub-grain, interatomic spacing was uniformly distributed with an interatomic spacing being 0.24 nm, same as the spacing in the specimen before LSP. However, interatomic spacing in the

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dislocation tangle decreased to 0.21 nm due to the presence of dislocation network suggesting that significant plastic deformation occurred in the specimen during the LSP process [33]. During LSP, the laser-driven shock waves propagated into the material and generated a high magnitude of

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shock-induced pressure, and the shock wave was the origin of the plastic deformation in the

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near-surface region. Reduced IFFT is revealed to obtain a clear observation of atomic dislocations, as shown in Fig. 6d. Figs. 6e and 6f illustrate profiles of line I and II in Fig. 6d.

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Fig. 6 Typical TEM images in the surface layer of specimen after LSP. (a) high density of dislocations, (b) corresponding SAED pattern of Ellipse [B] in (a), (c) HRTEM image, (d) IFFT image from crystal orientation of

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(1 1 1) , (e) reduced IFFT image of (d), (f) and (g) profile of line I and II in IFFT image.

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Typical TEM images at the depth around 500 µm below the top surface, known as minor plastic deformation layer, are depicted in Fig. 7. Fig. 7a shows clearly that a large number of mechanical

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twins in the direction of (1 1 1) were induced. It is also worth noting that a Cu enrichment region has been observed, which can be further proved by the energy dispersive spectrometer (EDS) maps in Figs. 7c and 7d. It can be found that Al and Cu uniformly distributed in the grain, with a small region of Cu enrichment inside. The corresponding SAED pattern captured from the ellipse [C] in Fig. 7a with [ 1 12 ] being the direction of crystal axis is shown Fig. 7b. Figs. 7f and 7g present the IFFT and reduced IFFT images in crystal orientation of (1 1 1) from HRTEM image (Fig. 7e) taken from the ellipse [C] in Fig. 7a. Interatomic spacings of line III and IV stand for different 14

ACCEPTED MANUSCRIPT atomic arrangements in the undeformed region and in the mechanical twin region, respectively. In the mechanical twin region, interatomic spacing was disorderly distributed with an interatomic spacing being 0.23 nm. A small number of atomic dislocations have been detected in the reduced

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IFFT image. Figs. 7h and 7i illustrate profiles of lines III and IV in Fig. 7f.

Fig. 7 Typical TEM images at a depth around 500 µm below the top surface of specimen after LSP. (a) high density of mechanical twins, (b) corresponding SAED pattern of Ellipse [C] in (a), (c) EDS map of element Al in (a), (d) EDS map of element Cu in (a), (e) HRTEM image, (f) IFFT image from crystal orientation of

(1 1 1) , (g)

reduced IFFT image of (f), (h) profile of line I in IFFT image and (i) profile of line II in IFFT image. 15

ACCEPTED MANUSCRIPT Distinct differences with the observation of dislocations in the surface layer, mechanical twins in the layer at the depth around 500 µm below the top surface, and randomly distributed twins in the specimen before LSP. Fig. 8 summarizes the interatomic spacing at different depths in WAAMed

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aluminum alloy after LSP. It is worth noting that the interatomic spacing increased from around 0.21 nm at the surface to around 0.23 nm at 500 µm in depth layer and further to more than 0.24 nm in the base metal. This increase in interatomic spacing indicates different intensiveness of

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interaction between shock wave and the target material during its transmission. Severe plastic

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deformation layer and minor plastic deformation layer were generated at the surface and 500 µm in depth, respectively, as the dislocations shown in Fig. 6a and the twins shown in Fig. 7a. In addition,

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the decrease of interatomic spacing indicated the existence of compressive residual stress.

Fig. 8 Interatomic spacing from LSP treated surface to 500 µm in depth and to the base metal of WAAMed aluminum alloy.

3.2 Micro-hardness Fig. 9 depicts the micro-hardness in depth direction of the topmost section and the middle section, measured to a depth of 1.5 mm, before and after LSP. It can be found that LSP has a dramatic

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ACCEPTED MANUSCRIPT influence on micro-hardness. The micro-hardness of the specimen before LSP was around 75 HV. After Peening, the surface micro-hardness increased to around 110 HV in both the topmost section and the middle section. Then the micro-hardness gradually decreased to a depth of around 1.2 mm

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when it stabilized around 75 HV. Hence, the corresponding hardening layer reached a thickness of 1.2 mm. Ding et al. also reported an improvement of micro-hardness in WAAMed nickel aluminum bronze alloy subjected to LSP treatment [34]. The difference in the improvement rate in different

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depth can be attributed that the laser shock wave has a more intensive interaction with the material

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in the surface layer, while the laser shock wave attenuates when it reaches the inner layer, generating less dislocations and mechanical twins. According to Hall-petch theory [35], micro-hardness value, HV, can be expressed by the following Eq1:

HV = HV0 + aGbρ 1/ 2

(1)

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where HV0 is the micro-hardness of an ideal material without any defects, a is a constant of material, G is the shearing modulus, and ρ is the dislocation density. In this study, higher densities of dislocations were generated in the surface layer than that in the deeper layers. Hence, a larger

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enhancement in the surface layer than that of the inner layer can be achieved [36]. Peyre et al.

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reported the similar result that LSP treatment is a feasible method to obtain higher micro-hardness and larger depth of the hardened layer [37].

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Fig. 9 Micro-hardness distributions in the depth direction before and after LSP.

3.3 Residual stress

Fig. 10 exhibits the comparison of residual stresses in the topmost section and middle section from

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surface to a depth of 0.75 mm before and after LSP. Tensile residual stresses in both the topmost section and middle section were generated before LSP [38], and their value fluctuated between 0

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and 40 MPa, and 0 and 60 MPa, respectively. Shrinkage of the cooling material behind the molten pool is one of the main reason for generating tensile residual stress [39]. The heat treatment resulted

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from the thermal input during the wall-printing could also induce the redistribution of residual stress. In Martina’s study, they reported that the wall produced by scanning rather than a single pass, could also affect the generation of residual stress [40]. However, efforts are still in demand to reveal the mechanism of residual stress generation and redistribution in 3D printing process. For the specimen after LSP, compressive residual stresses with a maximal value around 100 MPa and an affected depth of more than 0.75 mm in the topmost section and middle section were obtained. The surface residual stress of the topmost section and middle section were similar due to severe plastic 18

ACCEPTED MANUSCRIPT deformation generated by LSP. Kalentics et al. integrated LSP with selective laser melting and found that higher and deeper compressive residual stress was generated in the subsurface of the

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produced part as well [41].

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Fig. 10 In-depth residual stress in the topmost section and middle section before and after LSP.

During LSP, laser shock wave propagating into the material is equivalent to a Gauss distributed

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pressure pulse loaded on the surface. When the pressure pulse is higher than Hugoniot elastic limit

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(HEL) [42], it creates almost pure uniaxial compressive plastic deformation in the depth direction and tensile extension in the plane parallel to the surface [43], forming compressive residual stress at the material surface. As the shock wave dissipates in the material, the peak pressure of the shock wave decreases, while the deformation of material continues until the peak pressure attenuates below the dynamic yield strength [44]. Hence, a compressive stress field is generated within the affected depth [45]. When the materials are stabilized, the plastic deformation caused by shock waves leads to strain hardening and imparts residual stresses in the material. Fig. 11 illustrates the

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ACCEPTED MANUSCRIPT residual stress map in the as-built wall. Before LSP, it has tensile residual stress in the surface layer, while the stress in the surface layer turns into compressive after LSP. Due to the limit of residual stress measuring equipment, the residual stress below 0.75 mm was not measured. However, The

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compressive-tensile stress structure has been reported in our previous study with finite simulation

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[21]. The tensile state plays a role in balancing the generated compressive stress.

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Fig. 11 Schematic diagram of residual stress distribution in the WAAMed wall.

3.4 Tensile properties

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The room temperature tensile properties of specimens before and after LSP are shown in Fig. 12. Tensile force was applied in deposit direction in order to reveal the mechanical property between different layers. From Fig. 12a, it can be found that the engineering stress has a similar relationship with the engineering strain when the engineering strain is lower than 0.4% for specimens before and after LSP. Moreover, no significant change is observed in the slope of linear region for specimens before and after LSP, which suggests that LSP has little effect on the Young's modulus [46]. The serrated pattern on both specimens before and after LSP indicates that dynamic strain aging were 20

ACCEPTED MANUSCRIPT generated, which is a strengthening phenomenon in metals and alloys caused by the interaction between the diffusive solute atoms and the moving dislocation [47]. It is also worth noting that the dynamic strain aging in specimen before LSP is more severe that in LSPed specimen. This

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phenomenon can be attributed that LSP generated strengthening compensated for the dynamic strain aging. Fig. 12b summarized the YS, UTS and E of specimens before and after LSP. It can be discovered that the YS, UTS and E were 103.7 MPa (point A), 247.7 MPa (point C) and 12.3%

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(line segment GH) before LSP, respectively, while a significant increase of 72% in YS, from 103.7

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MPa to 178.3 MPa (point B), a similar UTS changing from 247.7 MPa to 240.3 MPa (point D) and a dramatically drop in E from 12.3% to 6.0% (line segment MN) were achieved after LSP. Ye et al. investigated tensile strength of 7075 aluminum alloy subject to warm LSP, and found the similar results in YS and E [48]. Luo et al. studied coverage area on tensile properties by increasing LSPed

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length on the gauge, and discovered the same tendencies when increasing LSPed length from 20 mm to 30 mm [49]. In another study, they discussed the effect of sheet thickness on tensile properties, finding that the generated residual stress distribution and grain size arrangement in the

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depth direction were two important factors [46]. In current study, compressive residual stress layers

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were introduced on both side of the specimen and grain refinement was also achieved, which can be one of the main factor for enhancement in the tensile properties. Besides, LSP induced localized strain hardening in the peened regions is another factor [50-53]. This strain hardening effect can be expounded as following. For specimen before LSP, the applied stress was first elastically loaded to its elastic limit, and then it reached the yielding point A. After yielding, the stress further went up to the ultimate tensile point C. Therefore, points A and C were YS and UTS, respectively. Its elongation was GH as indicated in Fig.12a. For specimen after LSP, a right translationally moved

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ACCEPTED MANUSCRIPT stress-strain curve was draw, as dotted line shown in Fig.12a. The applied force was similarly loaded to the elastic limit and reached the yielding point B’. After yielding, the stress-strain curve coincided with that for specimen before LSP and broke at point D. The correspondent YS and UTS

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were point B’ and point D, and the elongation MN. Hence, LSP can remarkably enhance tensile property in YS from point A to point B’, while cause a dramatically decrease in E and little

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influence on UTS.

Fig. 12 (a) Engineering stress-strain curves of specimens before and after LSP, and (b) comparison of YS, UTS and E before and after LSP.

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ACCEPTED MANUSCRIPT Figs. 13b to 13d are magnified rectangle [b] to [d] images in Fig. 13a, which show typical SEM fracture morphologies of specimens before LSP. It was featured by large number of homogeneous dimples and some micro-voids, which indicated that a ductile fracture occurred before LSP. Figs.

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13f to 13h are magnified rectangle [f] to [h] images in Fig. 13e, which demonstrate typical SEM fracture morphologies of specimens after LSP. Smooth cleavage planes, cleavage steps and tearing ridges are important features on the fracture surface. These features are the typical characteristic of

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cleavage fracture mode, indicating relatively poor plasticity. In addition, small number of shallow

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dimples and micro-voids can also be observed, suggesting that ductile fracture also existed in the specimen after LSP. Therefore, the fracture transformed from a ductile fracture to a mixture of

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ductile and cleavage fracture after LSP.

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Fig. 13 SEM fracture morphologies of specimens before and after LSP. (a) to (d) before LSP, and (e) to (h) after LSP.

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4. Conclusions In this study, LSP was introduced as an innovative post processing method for WAAM to not only modify its microstructure and stress state but also enhance its tensile properties. The main findings

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are listed below. (1) LSP can significantly refine the microstructure with the average size decreased from 59.7 µm to 46.7 µm after LSP. High density of dislocations and mechanical twins were generated in surface

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layer and 500 µm depth layer, respectively. A less misorientation grains with weak texture were

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achieved as well.

(2) Dramatic improvements of micro-hardness in the surface layer with a 1.2 mm hardening layer were obtained due to high density of dislocations generated by LSP. Residual stresses in both the topmost section and low section were modified from tensile residual stress to compressive residual

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stress with its maximum value around 100 MPa and an affected depth of more than 0.75 mm. (3) Tensile properties of YS was remarkably enhanced by 72%, while a dramatically decrease in E and little or no influence on UTS were found.

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This study proves that LSP is an effective surface enhancement method for improving mechanical

fields.

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properties of additive manufactured components applied in diversified scientific and industrial

Acknowledgments

The authors would like to acknowledge the support of this research work from the National International Collaborative Science and Technology Research Project China (Grant No. 2013DFR50590), Civil aircraft special research projects (Grant No. MJ-2016-F-16) and National

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ACCEPTED MANUSCRIPT Natural Science Foundation of China (Grant No. 11372019). In addition, special thanks go to Dr. Li Yan for linguistic assistance during the preparation of this manuscript.

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List of Figure captions Fig. 1 (a) Schematic diagram of layer-by-layer WAAM deposition and (b) image of additively manufactured 2319 aluminum specimen and the substrate.

shock peened WAAM 2319 aluminum specimen.

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Fig. 2 (a) Schematic diagram of LSP, (b) laser path illustrating an overlapping rate of 50% and (c) image of laser

Fig. 3 Schematic diagram of (a) cross-section view for micro-hardness and residual stress tests, (b) specimen

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preparation for tensile test and (c) detailed dimension of tensile test specimen.

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Fig. 4 EBSD analysis of inverse pole figure, grain size, texture and grain boundary misorientation. (a), (c), (e) and (g) for specimen before LSP and (b), (d), (f) and (h) for specimen after LSP.

Fig. 5 Typical TEM images of specimen before LSP. (a) mechanical twins randomly distributed in the original grain, (b) corresponding SAED pattern of Ellipse [A] in (a), (c) HRTEM image, (d) IFFT image from crystal

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orientation of (1 1 1) , (e) profile of line in IFFT image.

Fig. 6 Typical TEM images in the surface layer of specimen after LSP. (a) high density of dislocations, (b) corresponding SAED pattern of Ellipse [B] in (a), (c) HRTEM image, (d) IFFT image from crystal orientation of

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(1 1 1) , (e) reduced IFFT image of (d), (f) and (g) profile of line I and II in IFFT image.

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Fig. 7 Typical TEM images at a depth around 500 µm below the top surface of specimen after LSP. (a) high density of mechanical twins, (b) corresponding SAED pattern of Ellipse [C] in (a), (c) EDS map of element Al in (a), (d) EDS map of element Cu in (a), (e) HRTEM image, (f) IFFT image from crystal orientation of

(1 1 1) , (g)

reduced IFFT image of (f), (h) profile of line I in IFFT image and (i) profile of line II in IFFT image. Fig. 8 Interatomic spacing from LSP treated surface to 500 µm in depth and to the base metal of WAAMed aluminum alloy. Fig. 9 Micro-hardness distributions in the depth direction before and after LSP.

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ACCEPTED MANUSCRIPT Fig. 10 In-depth residual stress in the topmost section and middle section before and after LSP. Fig. 11 Schematic diagram of residual stress distribution in the WAAMed wall. Fig. 12 (a) Engineering stress-strain curves of specimens before and after LSP, and (b) comparison of YS, UTS

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and E before and after LSP. Fig. 13 SEM fracture morphologies of specimens before and after LSP. (a) to (d) before LSP, and (e) to (h) after

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LSP.

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List of Table captions Table 1 Chemical compositions of 2A12 aluminum substrate and 2319 aluminum wire (wt %).

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Table 2 Parameters used for WAAM experiment.

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ACCEPTED MANUSCRIPT Laser shock peening and wire arc additive manufacturing are specially-combined.




Grain refinement with higher percentage of low angle boundaries is achieved.




Tensile residual stresses are modified to beneficial compressive residual stresses.




Yield strength is remarkably increased by 72%.

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