Accepted Manuscript

Title: Laser additive manufacturing of carbon nanotubes(CNTs) reinforced aluminum matrix nanocomposites:Processing optimization, microstructure evolution andmechanical properties

Authors: Dongdong Gu, Xiangwei Rao, Donghua Dai,Chenglong Ma, Lixia Xi, Kaijie Lin

PII: S2214-8604(18)31039-XDOI: https://doi.org/10.1016/j.addma.2019.100801Article Number: 100801

Reference: ADDMA 100801

To appear in:

Received date: 15 December 2018Revised date: 14 July 2019Accepted date: 21 July 2019

Please cite this article as: Gu D, Rao X, Dai D, Ma C, Xi L, LinK, Laser additive manufacturing of carbon nanotubes (CNTs) reinforcedaluminum matrix nanocomposites: Processing optimization, microstructureevolution and mechanical properties, Additive Manufacturing (2019),https://doi.org/10.1016/j.addma.2019.100801

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Revised Manuscript ADDMA_2018_1000_R2

For Re-consideration in Additive Manufacturing

Laser additive manufacturing of carbon nanotubes (CNTs) reinforced aluminum

matrix nanocomposites: Processing optimization, microstructure evolution and

mechanical properties

Dongdong Gua,b,*, Xiangwei Raoa,b, Donghua Daia,b, Chenglong Maa,b, Lixia Xia,b,

Kaijie Lina,b

a College of Materials Science and Technology, Nanjing University of Aeronautics and

Astronautics (NUAA), Yudao Street 29, Nanjing 210016, Jiangsu Province, PR China

b Jiangsu Provincial Engineering Laboratory for Laser Additive Manufacturing of High-

Performance Metallic Components, Nanjing University of Aeronautics and Astronautics

(NUAA), Yudao Street 29, Nanjing 210016, Jiangsu Province, PR China

*Corresponding author. E-mail: dongdonggu@nuaa.edu.cn (D. Gu).

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Abstract:

In this study, a laser-based additive manufacturing route of selective laser melting

(SLM) was applied to fabricate carbon nanotubes (CNTs) reinforced Al-based

nanocomposites with tailored microstructures and excellent mechanical properties. The

densification behavior, microstructure features and mechanical properties were

investigated and the relationship between process and property was established. The

results showed that the applied laser power and scan speed were the governing factors

of the densification behavior of SLM-processed Al-based nanocomposites. SLM

processing of 0.5wt.% CNTs/AlSi10Mg nanocomposite powder led to the formation of

three typical microstructures including the primary Al9Si cellular dendrites decorated

with fibrous Si, the in situ Al4C3 covered on CNTs, and the precipitated Si inside the

cellular grains. As the optimal SLM processing parameters of laser power of 350 W and

scan speed of 2.0 m/s were applied, the fully dense SLM-processed CNTs/Al-based

nanocomposites exhibited high microhardness of 154.12 HV0.2, tensile strength of 420.8

MPa and elongation of 8.87%, due to the formation of high densification and ultrafine

microstructure. The grain refinement effect, Orowan looping system and load transfer

are considered as three strengthening mechanisms occurred simultaneously during

tensile tests, leading to excellent mechanical properties of SLM-processed CNTs/Al-

based nanocomposites.

Key words: Additive manufacturing; Selective laser melting; Carbon nanotubes

(CNTs); Aluminum matrix nanocomposites; Mechanical properties

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1. Introduction

In recent years, with an increasing demand for lightweight and high strength

materials, the development of aluminum matrix composites (AMCs) is of particular

interest. AMCs, which combine the excellent properties of metallic matrix and

reinforcing phases, are expected to exhibit higher strength, stiffness, wear resistance,

but lower thermal expansion coefficient as relative to unreinforced aluminum alloys [1].

AMCs have demonstrated significant potential in engineering applications in

automotive, aerospace and defense industries [2]. Among various reinforcing materials,

carbon nanotubes (CNTs) are regarded as a new-generation material since the first

discovery by Iijima in 1991 [3], because of their unique properties such as ultrahigh

strength (up to 100 GPa), ultrahigh Young’s modulus (up to 1 TPa) and large aspect

ratio (50-500) [4]. However, several challenges are still difficult to overcome in

preparing CNTs reinforced AMCs [5]. One significant difficulty is that CNTs tend to

agglomerate due to the considerably high aspect ratio. Another problem is the poor

wetting between CNTs and molten metals caused by a large difference in surface

tensions, resulting in weak interfacial bonding and low densification level.

A variety of processing techniques were applied to produce CNTs reinforced

AMCs in the past decade, among which powder metallurgy (PM) was a commonly used

method [6]. Wu et al. [7] successfully synthesized Al6061-CNTs composite through

semi-solid powder processing from the mechanically alloyed powders at different

durations. The results showed that the mechanical alloying could crush the

agglomerated CNTs and accordingly disperse CNTs uniformly. Zhou et al. [8] used

spark plasm sintering to fabricate multi-walled CNTs (MWCNTs) reinforced Al-based

composite. They found that the yield strength of MWCNTs/Al composites was

substantially increased with an appropriate quantity of Al4C3 produced at the MWCNT-

Al interface. Chen et al. [9] also confirmed that the formation of interfacial Al4C3 on the

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partially reacted CNTs led to a significant improvement of interfacial strength and

attendant load transfer efficiency in AMCs. Although these conventional processing

techniques can be applied to produce CNTs reinforced AMCs parts, the obtainable

microstructures are typically coarsened due to the relatively long heating duration,

thereby weakening the reinforcing effect of CNTs. The flexibility of traditional forming

processes is also limited and, therefore, it is rather difficult to build AMCs parts with

complex geometries.

Selective laser melting (SLM), as an emerging additive manufacturing (AM)

method, shows considerable potential in fabricating metallic components with desired

structures and properties [10–12]. During SLM process, a laser beam controlled by

computer selectively scans a layer of metallic powder, fusing and consolidating the

powder particles into the designed configurations according to computer aided design

(CAD) data. In view of processing capability, SLM can realize a quick fabrication of

complex shaped components [13,14]. Moreover, SLM can produce fully dense metallic

parts within one-step manufacturing, due to the metallurgical mechanism of the

complete melting of powder [15,16]. The cooling rate during SLM process is extremely

high (up to 105-7K/s) [17] and, therefore, SLM is capable of forming very fine and

unique microstructure compared with conventional processing techniques [18–20].

In this work, SLM AM of CNTs reinforced AlSi10Mg powder was performed to

produce high-performance AMCs components. As a typical Al-Si casting alloy,

AlSi10Mg is relatively easy to process by laser AM, due to the near eutectic

composition of Al and Si that leads to a small solidification temperature range [21].

Nevertheless, due to the incorporation of CNTs reinforcing phase, it is rather difficult to

optimize the SLM processing parameters for fabricating high-performance CNTs/Al-

based nanocomposites. The densification behavior, processing optimization,

microstructure evolution and mechanical properties were systematically investigated for

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SLM processing of CNTs/Al-based nanocomposites to establish a relationship among

processing parameters, microstructures and properties.

2. Experimental procedures

2.1 Powder materials

Gas-atomized AlSi10Mg powder (Fig. 1a) with a particle size range from 5 μm to

70 μm (Fig. 1b) and commercial multi-walled carbon nanotubes (MWCNTs) with an

outer diameter of 10-20 nm and a length of 10-30 μm (Fig. 1c) were used as raw

materials for preparation of CNTs/Al-based nanocomposite powder. A QM-3SP4

planetary ball milling machine (Nanjing NanDa Instrument Plant, China) was used to

homogenously disperse CNTs on the surfaces of AlSi10Mg powder particles. The

powder mixture with a mass fraction of 0.5% of CNTs was sealed in a stainless steel

bowl with a ball-to-powder ratio of 2:1. The rotation speed was set at 200 rpm and the

total milling time was 4 h. An interval of 5 min was set after each 15 min of milling in

order to avoid overheating of powder that may cause damage to the structural integrity

of CNTs. Fig. 2a shows the morphology of Al-based nanocomposite powder after ball

milling, indicating a very slight deformation of AlSi10Mg matrix powder during

milling. A high sphericity of powder guaranteed a high flowability and resultant SLM

processability. The elemental distribution mapping of carbon element revealed a

uniform distribution of CNTs on the surface of AlSi10Mg powder after ball milling

(Fig. 2b-e).

2.2 SLM process

An independently developed SLM system by NUAA mainly consisted of a YLR-

500-WC ytterbium fiber laser with a maximum power of ~500W, a spot size of ~70 μm

and a continuous wavelength of 1070 ± 10 nm (IPG Laser GmbH, Germany), a

hurrySCAN 30 scanner with a scan speed up to 7.0 m/s (SCANLAB GmbH, Germany),

an automatic powder layering device, an inert gas protection system, and a computer

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control system (Fig. 3a). When the specimens were to be prepared, an aluminum

substrate was leveled and fixed on the building platform. Afterwards, the building

chamber was evacuated and filled with argon gas to prevent oxidation during

processing. A “island” laser scanning strategy was applied for SLM, with island size of

7.5 mm × 7.5 mm, hatching spacing of 50 μm, and scan direction rotation of 37°

between neighboring two layers (Fig. 3b). In order to optimize the laser processing

parameters, a series of variables of laser power (P) and scan speed (v) were settled in

Fig. 3c. The powder layer thickness was fixed at 50 μm. The first batch of cubic

specimens with dimensions of 10 mm×10 mm×5 mm was prepared for process

optimization and microstructural characterization. Another batch of rectangular

specimens with dimensions of 70 mm×14 mm×5 mm was further built using the

optimized SLM parameters for the preparation of standard tensile test samples (Fig. 3d).

2.3 Microstructural characterization and mechanical properties tests

After SLM process, all specimens were cut from the substrate by wire electro

discharge machining (EDM) for the following microstructural and mechanical tests.

Specimens for microstructural characterization were ground and polished according to

standard procedures, and then etched by Keller reagent (HNO3 2.5 ml, HCl 1.5 ml, HF 1

ml and H2O 95 ml) for metallographic examination. The low-magnification

microstructures and densification behaviors of specimens were characterized by a

PMG3 optical microscopy (OM, Olympus Corporation, Japan). Phase identification was

performed by X-ray diffraction (XRD) using a D8 Advance X-ray diffractometer

(Bruker AXS GmbH, Germany) with Cu Kα radiation at 40 kV and 40 mA. The scan

mode was continuous and the scan speed was set at 4° min-1 with 2θ range of 20-90°.

The structural integrity of MWCNTs after ball milling was measured by a Renishaw

RM 2000 Raman spectroscopy, with excitation by Ar + laser line of 514 nm.

Microstructure observation was performed using a Zeiss Sigma 300 field emission

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scanning electron microscope (FE-SEM, Carl Zeiss AG, Germany). Chemical

compositions were characterized using a Bruker XFlash 6160 Energy Dispersive X-ray

Spectroscopy (EDX, Bruker Daltonics Inc., USA).

The relative density was calculated by the ratio of the actual density measured by

Archimedes principle to the theoretical density. The Vickers hardness were measured on

the cross-sections of SLM-processed specimens along the building direction, using a

HXS-1000AY microhardness tester (AMETEK, Shanghai, China) with a load of 200 g

and a dwell time of 15 s. Twenty indentations with an interval of 0.15 mm were

obtained for each specimen. The as-fabricated rectangular specimens were machined

into dog-bone shaped testing samples according to ASTM E8 standard (Fig. 3d). Three

specimens were fabricated and tested under each SLM processing parameters. The

tensile tests were carried out at room temperature using a CMT5205 testing machine

(MTS Industrial Systems, China) with a cross head velocity fixed at 0.2 mm/min. After

tensile tests, the morphologies of fracture surfaces of samples were studied by a Zeiss

Sigma 300 FE-SEM.

3. Results and discussion

3.1 Densification behavior and process optimization

Fig. 4a shows the variation of relative density of SLM-processed CNTs/Al-based

nanocomposite parts under a series of processing parameters. Generally, for a fixed laser

scan speed, the relative density increased with increasing laser power. For a fixed laser

power, the relative density firstly increased and then decreased with an increase in laser

scan speed. It is known that aluminum has the intrinsic properties of high reflectivity

and thermal conductivity, which may cause the deficiency of laser energy input and

resultant incomplete fusion of powder during SLM process. A high laser power ensured

a sufficient laser energy input and, therefore, the relative densities of specimens were

generally over 99% at a relatively high laser power of 350 W. The corresponding OM

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images also reveals an enhanced densification activity of the specimens with increasing

laser powers. At a relatively low laser power of 300 W, large-sized irregular pores

appeared along the molten pool boundaries. As the laser power increased to 325 W, the

large-sized pores changed into a few small-sized pores. When laser power further

increased to 350 W, the as-fabricated specimen was fully dense and free of any obvious

pores and cracks.

Due to the line-by-line and layer-by-layer manufacturing manner of SLM process,

a sufficient overlap between scan tracks is crucial for a high densification level of as-

built part. The wetting behavior of liquid phase, that is determined by the interaction

between laser beam and powder bed, has a strong impact on the interlayer bonding

ability. A high laser power yields a high energy input and resultant high temperature of

liquid phase, which favors the breaking of the oxide film on aluminum melt and the

decrease of dynamic viscosity (μ) of the melt [22]. A lower μ ensures a sufficient

spreading of liquid on the previously solidified layer, enhancing the inter-layer

wettability and consolidation level. A high laser energy input also increases the width

and depth of the molten pool, thereby increasing the overlap rate between the

neighboring tracks and layers to decrease the pore formation. On the other hand, at a

constant laser power, the variation curves of relative density indicate that a critical value

of scan speed exists to yield a maximum densification level. When the applied scan

speed is low, the dwelling time of laser beam becomes longer and the over-melting of

powder tends to occur. The degree of superheating of melt and resultant material

vaporization intensify, inducing defects such as keyholes and thermal cracks and

lowering the densification level. Furthermore, the time interval between melting and

solidification becomes longer at a lower scan speed, resulting in a more significant

growth of the trapped hydrogen pores and a limited densification rate [23]. When the

scan speed becomes too high, the capillary instability in the molten pool intensifies,

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causing the splash of liquid droplets on the surface of pool and the resultant balling

phenomena [24]. The occurrence of balling phenomena results in the irregular and

rough surface, which hinders the sufficient spreading of the melt on the previously

processed layer and hence lowers the densification level. Fig. 4b shows the 3D view of

an etched specimen fabricated with process optimization at a laser power of 350 W and

a scan speed of 2.0 m/s. No apparent pores or cracks are found along the boundaries of

molten pools and between the neighboring layers.

3.2 Phase identification

Typical XRD spectra of initial powder and SLM-fabricated specimens at various

parameters are depicted in Fig. 5. The strong diffraction peaks corresponding to Al9Si

(PDF#65-8554) and Si (PDF#27-1402) were generally identified. The diffraction peaks

of CNTs were not detected due to the addition of a very low content of CNTs. The

characteristic diffraction peaks of Al9Si showed a trend to shift to the larger 2θ locations

as relative to the standard peaks (Table 1), which indicated that the distance of lattice

plane changed after SLM process according to Bragg’s law [25]. Compared with the

starting Al-based powder, the intensity of Al9Si diffraction peaks in SLM-processed

specimens typically decreased when the applied scan speed increased. As the scan speed

was above 2.2 m/s, the intensity of Al9Si peak at ~44.9° became higher than the peak at

38.7°, suggesting that the preferred crystallization direction of Al9Si transferred from

(111) to (200) lattice plane. Table 1 also reveals that the full width at half maxima

(FWHM) increased with increasing the applied laser scan speed, indicating a more

significant grain refinement effect of Al9Si at a higher scan speed. For the detected Si

phase, Fig. 5 shows that its peak intensity lowered and the peak width broadened when

the applied laser scan speed increased, implying the formation ultrafine Si crystals

during SLM process. Generally, a higher laser scan speed yields a superfast cooling rate

of the molten pool during SLM, thereby hindering the grain growth of the precipitated

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Si phase.

Raman spectroscopy is a useful method to evaluate the structural integrity of CNTs

and to characterize the formation of carboniferous compounds [26]. The D-band at 1340

cm-1 is known as the damage peak that is generated by the structural defects of C-C

bonding, whereas the G-band at 1570 cm-1 is regarded as a signal of the intact structure

of CNTs [27]. The ratio of the intensity of D-band and G-band (ID/IG) is normally used

as a criterion to characterize the state and performance of CNTs. Fig. 6 depicts the

Raman spectra collected from the initial MWCNTs, ball milled CNTs/Al-based powder

and SLM-fabricated specimens. As relative to the ID/IG of the initial MWCNTs of 0.78,

the ID/IG of ball milled CNTs/Al-based composite powder increased to 0.99, which

indicated that ball milling would induce structural defects of MWCNTs. The SLM-

processed specimens demonstrated a considerable increase in ID/IG, especially when a

relatively low scan speed of 1.8 m/s was applied. This phenomenon implied that a high

energy density laser irradiation and resultant high SLM working temperature could

induce a large number of crystal defects into the MWCNTs. The characteristic peaks of

Al4C3 that usually presented at 490 cm-1 and 860 cm-1 were observed in SLM-fabricated

specimens. When a relatively high scan speed of 2.4 m/s was applied, the Al4C3 peaks

became insignificant, indicating that the reaction between CNTs and matrix was very

limited in this instance. However, the characteristic peaks of Al4C3 became sharp and

evident when the laser energy input increased by lowering scan speed, indicating that a

high laser energy input contributed to the in situ reaction of CNTs with Al matrix.

3.3 Microstructural evolution and underlying mechanisms

Fig. 7 shows the characteristic microstructures of SLM-processed CNTs/Al-based

part along the layer-by-layer building direction. The molten pools in SLM layers

showed a clear “fish-scale” structural feature (Fig. 4b) and the profile of the pools is

revealed in Fig. 7a. Fig. 7b, as the detailed view of dotted section A in Fig. 7a, exhibits a

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hierarchical microstructure featured by three different microstructural zones. Generally,

the microstructure consisted of cellular dendrites; the grey cells were primary Al9Si

phase that decorated with white fibrous Si phase. Fig. 7c and d present the high-

magnification microstructures of three different zones. The transition zone was induced

by the heating effect from the adjacent scan track or the subsequent layer. Therefore, the

cellular boundaries of fibrous Si phase in transition zone became discontinuous due to

the heat treatment. Inside a molten pool, the cellular dendrites grew towards the center

of the pool because of the temperature gradient. The cells near the molten pool

boundaries were relatively coarse, while the cells near the top of the pool were

considerably fine. According to Liu et al.’s study [28], the cooling rate in the bottom

area of the molten pool was much lower than that in the top area, resulting in the

slightly coarse dendritic structure. The hierarchical microstructure of SLM-processed

CNTs/Al-based composites is similar to the microstructure of SLM-processed

AlSi10Mg aluminum alloy [29].

Fig. 8 reveals the effect of SLM processing parameters on the microstructure

evolution of as-fabricated CNTs/Al-based specimens. All these SEM images were

obtained from the fine cellular zones. Generally, the specimens showed the considerably

fine microstructures due to the superhigh cooling rate of SLM process. With an increase

in laser scan speed, the size of cellular dendrites became smaller (Fig. 8). The dwelling

time of laser beam at a certain radiation position decreased when the applied laser scan

speed increased, thereby intensifying the cooling rate in the molten pool and resultant

grain refinement effect of cellular dendrites. Fig. 9a to e illustrate the elemental

distributions in the fine cellular zone of SLM-processed specimen at a scan speed of 1.8

m/s. The Si element had a high concentration at the primary Al9Si cell boundaries,

while the Mg element had a uniform distribution throughout the microstructure. The C

element concentrated mainly inside the grey cells. The EDS scan in the marked

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positions in Fig. 9f further confirmed the compositions of different structures (Table 2).

The spectrum A collected from the grey cells indicated the presence of Al9Si and the

spectrum B collected from the rod-like structure revealed a relatively high content of C

element. The spectrum C collected from the white phase showed a high content of Si

element, which indicated that the Al9Si phase was decorated with the continuously

distributed Si particles. The high-magnification SEM image in Fig. 9g reveals three

characteristic constitutional structures: (i) the cellular dendritic Al9Si decorated with

fibrous Si particles; (ii) the Al4C3 covered on CNTs; and (iii) some precipitated Si

particles inside the cells.

Fig. 10 depicts schematically the formation mechanism of three different phases

and structures during SLM process. Due to the non-equilibrium metallurgical

characteristics of SLM process, the original Al was supersaturated with Si, thereby

forming the Al9Si solid solution after SLM. The presence of some precipitated Si

particles in the cells was caused by the accumulated heating effect during line-by-line

and layer-by-layer scanning and the resultant decrease in the solid solubility of Si in Al.

During ball milling preparation of CNTs/Al-based powder, large clusters of CNTs were

divided into a number of small bunches consisting of several individual CNTs, leading

to a uniform dispersion of the small bunches of CNTs around the surface of AlSi10Mg

powder particles. During SLM process, the laser energy was firstly absorbed by CNTs

on particle surface and the tubular structure of CNTs might be destroyed since the SLM

process normally involved a high working temperature. Furthermore, the CNTs had

many crystalline defects after ball milling, which further lowered the stability of CNTs

under high processing temperature [30]. A partial decomposition of CNTs and the

attendant diffusion of carbon atoms tended to occur in laser-induced molten pool. In situ

reaction thus occurred at the interface between CNTs and Al matrix, forming aluminum

carbide Al4C3 on the outer layer of MWCNTs. A relatively high scan speed produced a

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lower laser energy input and resultant limited diffusion of carbon atoms, resulting in a

low degree of in situ reaction between C and Al. When a relatively low scan speed was

applied, the elevated laser energy density and attendant high SLM temperature would

promote the atomic diffusion in the molten pool. Moreover, the intensity of Marangoni

flow in the molten pool also increased at a higher SLM temperature [31], contributing to

the vibration of melt and resultant increase in the nucleation rate of Al4C3. It was noted

that when the laser scan speed further decreased, the laser energy input might be

considerably high, which tended to cause splashing or even evaporation of CNTs.

Therefore, either at a relatively high scan speed of 2.4 m/s or at a relatively low speed of

1.8 m/s, the content of in situ Al4C3 covered on CNTs was lower than that formed at

scan speeds of 2.0 m/s and 2.2 m/s.

3.4 Mechanical properties

Fig. 11 shows the microhardness of SLM-fabricated specimens at different

processing parameters. When a relatively high scan speed of 2.4 m/s was applied, the

average microhardness, 133.35 HV0.2, was lowest among four specimens. With the scan

speed decreased from 2.2 m/s to 2.0 m/s, the average microhardness increased from

140.52 HV0.2 to 154.12 HV0.2. A further decrease in scan speed to 1.8 m/s, however,

resulted in a decrease in the average microhardness to 146.43 HV0.2. Nevertheless, all

SLM-processed CNTs/Al-based specimens exhibited higher hardness than SLM-

processed AlSi10Mg alloy parts (127 ± 3 HV0.5 [29]) and CNTs/Al parts fabricated via

powder metallurgy (<100 HV) [32]. The densification level and microstructural feature

were found to have a significant effect on the microhardness distributions. At a

relatively high scan speed, a significant fluctuation of microhardness distribution was

observed, which was caused by the limited densification level in this instance. As an

optimal scan speed of 2.0 m/s was applied, the combined effect of the elevated

densification and the reinforcement of precipitated Si particles and CNTs covered with

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Al4C3 led to a high microhardness. The main factor that governed the increase of

microhardness was the content of Al4C3 covered on CNTs. The existence of Al4C3 on

the outer layer of CNTs improved the interfacial bonding between reinforcement and

matrix [33], which further enhanced the plastic deformation resistance of the matrix and

the attendant microhardness.

Fig. 12 depicts the tensile properties of SLM-processed CNTs/Al-based parts at

different parameters. A contrastive specimen of unreinforced AlSi10Mg alloy was

fabricated at a laser power of 350 W and a scan speed of 2.0 m/s. The tensile strength

and elongation of AlSi10Mg specimen were 348.94 MPa and 5.56%, respectively. The

CNTs/Al-based specimen fabricated at a scan speed of 2.0 m/s showed the highest

tensile strength of 420.8 MPa and ductility (8.87%) among all tensile specimens.

Although the specimen fabricated at 1.8 m/s also had a high strength of 400.1 MPa, the

elongation decreased to 7.67%. The CNTs/Al-based specimen fabricated at 2.4 m/s

exhibited a strength of 341.4 MPa that was comparable to AlSi10Mg alloy and an

elongation of 7.99% that was significantly higher than AlSi10Mg alloy without addition

of CNTs. In general, all SLM-processed CNTs/Al-based specimens exhibited excellent

tensile strength as relative to CNTs/Al-based composites fabricated via conventional

powder metallurgy route (normally lower than 200 MPa) [34]. The SLM-fabricated

CNTs/Al-based specimen using optimal processing parameters had considerably higher

strength and ductility than the unreinforced AlSi10Mg alloy specimen.

In order to further elucidate the fracture mechanisms, SEM images of the typical

morphologies of fracture surfaces of SLM-fabricated specimens under different

processing parameters are depicted in Fig. 13. For the specimen fabricated at 1.8 m/s, a

typical river pattern consisted of several cleavage planes implied a mechanism of

cleavage fracture (Fig. 13a), indicating a brittle fracture. For the specimen fabricated at

2.0 m/s, the SEM fractograph showed a large number of fine equiaxed dimples on the

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fracture surface (Fig. 13b), confirming the operation of a mechanism of ductile fracture.

For specimens fabricated at scan speeds of 2.2 m/s and 2.4 m/s, although a general

formation of fine dimples was observed, several residual pores or microcracks were

present on the fracture surfaces (Fig. 13c and d). Those defects could be the source of

cracking and initiated crack propagation, hence degrading the strength of as-fabricated

specimens. Therefore, realizing the production of a fully dense part is the precondition

of achieving high-strength CNTs/Al-based composite parts.

The strengthening mechanisms of SLM-processed CNTs/Al-based parts are

summarized as follows. The grain refinement is the first operative strengthening

mechanism. According to the Hall-Petch relationship, a decrease in average grain size

can improve the strength of metallic parts [35]. During SLM process, an extremely high

cooling rate leads to the formation of ultra-fine microstructure (Figs. 7 and 8). The

strength of SLM-processed CNTs/Al-based composites is much higher than that

fabricated through conventional powder metallurgy route. The second strengthening

mechanism is Orowan looping system. Fig. 9g shows the formation of nano-sized

structures of eutectic Si particles, precipitated Si particles and Al4C3 covered on CNTs.

It is believed that these ultra-fine nanostructures strengthen the Al matrix via the

mechanism of Orowan looping. During plastic deformation, these nanoscale structures

obstruct the motion of dislocations, leading to the “dislocation bending” between

nanostructures. The “dislocation bending” produces a back stress, which in turn

prevents further dislocation migration and increases the yield stress [36]. Load transfer

is the third strengthening mechanism of SLM-processed CNTs/Al-based composites.

The interfaces between reinforcement and Al matrix play a significant role in material

strengthening. A stable interfacial bonding assists in transferring load from matrix to

reinforcement during loading. Zhou et al. have confirmed that an appropriate quantity of

Al4C3 formed on the outer layer of MWCNTs can improve the interfacial stability,

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effectively enhancing the load transfer efficiency in MWCNTs/Al-based composites [8].

Therefore, the strength of SLM-processed CNTs/Al-based composites are substantially

elevated due to the efficient load transfer by producing Al4C3 at the interface between

CNTs and Al. The above-mentioned three strengthening mechanisms simultaneously

occur during deformation, favoring the improvement of tensile strength and ductility of

the optimally prepared CNTs/Al-based specimen by SLM.

4. Conclusions

(1) Laser power and scan speed were dominant factors to determine SLM densification

level. With an increase in laser power, large irregular interlayer pores disappeared

and densification rate increased. For a fixed laser power, a too high scan speed

caused the occurrence of balling phenomenon, while a too low scan speed resulted

in overheating and vaporization of materials. A laser power of 350 W and a scan

speed of 2.0 m/s were optimized to yield a fully dense CNTs/Al-based part.

(2) The major phases of SLM-processed CNTs/Al-based composites were Al9Si and Si.

SLM process introduced an in situ reaction between CNTs and Al, forming Al4C3

on the outer layer of CNTs when laser energy input was sufficiently high.

(3) The SLM-processed CNTs/Al-based composites demonstrated a hierarchical

microstructural feature. A typical solidified molten pool had three different zones

including coarse cellular zone, transition zone and fine cellular zone. Multiple

reinforcing phases including the Al4C3 covered on CNTs and the precipitated Si

were present in SLM-processed CNTs/Al-based composites.

(4) High densification level and significant grain refinement contributed to the high

microhardness (154.12 HV0.2) of SLM-processed CNTs/Al-based composites.

Three strengthening mechanisms of grain refinement, Orowan looping system and

load transfer occurred simultaneously during deformation, leading to considerably

high tensile strength of 420.8 MPa and elongation of 8.87% for SLM-processed

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CNTs/Al composites.

Conflict of Interest: The authors declare no competing interests.

Acknowledgments

This work is supported by the National Natural Science Foundation of China (grant

number 51735005), the National Key Research and Development Program “Additive

Manufacturing and Laser Manufacturing” (grant numbers 2016YFB1100101 and

2018YFB1106302), and the Equipment Pre-Research Field Fund (grant number

61409230311). D.D. Gu acknowledges the support from the National High-level

Personnel of Special Support Program of China, the Cheung Kong Young Scholars

Program of Ministry of Education of China, and the Top-Notch Young Talents Program

of China.

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Figure and Table Captions

Fig. 1. (a) SEM image showing the morphology of starting AlSi10Mg powder; (b)

Particle size distribution of AlSi10Mg powder; (c) SEM image showing the morphology

of starting CNTs.

Fig. 2. (a) SEM image showing the morphology of 0.5 wt.% CNTs/AlSi10Mg mixed

powder. EDX mapping showing the elemental distributions of (b) Al, (c) Si, (d) Mg and

(e) C on particle surface.

Fig. 3. Schematic of SLM equipment and process (a); Schematic of island scanning

strategy applied for SLM (b); Photography of SLM-fabricated specimens at different

processing parameters (c); (d) Dimensions of tensile test sample based on ASTM E8

standard.

Fig. 4. (a) Effects of laser power and scan speed on densification rate and corresponding

OM microstructure of SLM-fabricated specimens; (b) 3D OM view of etched specimen

fabricated at laser power of 350 W and laser scan speed of 2.0 m/s.

Fig. 5. XRD spectra of CNTs/Al-based composite powder and SLM-fabricated

specimens at various processing parameters.

Fig. 6. Raman spectra of initial MWCNTs, ball milled CNTs/Al-based composite

powder and SLM-fabricated specimens at different processing parameters.

Fig. 7. SEM image showing the hierarchical microstructure of specimen fabricated at P

= 350 W, v = 2.4 mm/s (a); Local magnification of area A in Fig. 7a revealing

microstructural difference of grains in three representative zones (b); High-

magnification SEM images showing the characteristic microstructures in (c) the coarse

cellular zone and transition zone and (d) the fine cellular zone.

Fig. 8. SEM images showing the evolution of characteristic microstructures of SLM-

fabricated CNTs/Al-based specimens using different processing parameters: (a) P = 350

W, v = 1.8 m/s; (b) P = 350 W, v = 2.0 m/s; (c) P = 350 W, v = 2.2 m/s; and (d) P = 350

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W, v = 2.4 m/s.

Fig. 9. Characteristic microstructure of SLM-processed CNTs/Al-based specimen

fabricated at P = 350 W and v = 1.8 m/s (a); EDX mapping showing the elemental

distributions of (b) Al, (c) Si, (d) Mg and (e) C elements; Typical microstructures of

SLM-processed CNTs/Al-based specimen produced at P = 350 W and v = 2.0 m/s,

observed at a relatively low magnification (f) and a higher magnification (g).

Fig. 10. Schematic of the formation mechanisms of different microstructures during

laser melting and solidification process.

Fig. 11. Microhardness distributions and corresponding indentation microstructures of

SLM-fabricated specimens at different processing parameters.

Fig. 12. Stress-strain curves (a) and tensile properties including ultimate tensile strength

(UTS) and elongation (b) of SLM-fabricated CNTs/Al-based specimens using different

processing parameters. Tensile properties of unreinforced AlSi10Mg specimen

fabricated at P = 350 W and v = 2.0 m/s are included for comparison.

Fig. 13. SEM images showing the characteristic fracture surface morphologies of SLM-

fabricated CNTs/Al-based specimens at different processing parameters: (a) P = 350 W,

v = 1.8 m/s; (b) P = 350 W, v = 2.0 m/s; (c) P = 350 W, v = 2.2 m/s; (d) P = 350 W, v =

2.4 m/s.

Table 1.

XRD data showing the displacement and intensity variation of the identified peaks for

Al9Si phase in SLM-processed specimens using different processing parameters.

Table 2.

EDX results showing chemical compositions collected in different positions in Fig. 9f.

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Figures and Tables ADDMA_2018_1000_R2

For Re-consideration in Additive Manufacturing

(a) (b)

(c)

Fig. 1. (a) SEM image showing the morphology of starting AlSi10Mg powder; (b)

Particle size distribution of AlSi10Mg powder; (c) SEM image showing the morphology

of starting CNTs.

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Fig. 2. (a) SEM image showing the morphology of 0.5 wt.% CNTs/AlSi10Mg mixed

powder. EDX mapping showing the elemental distributions of (b) Al, (c) Si, (d) Mg and

(e) C on particle surface.

(a) (b)

(c) (d)

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Fig. 3. Schematic of SLM equipment and process (a); Schematic of island scanning

strategy applied for SLM (b); Photography of SLM-fabricated specimens at different

processing parameters (c); (d) Dimensions of tensile test sample based on ASTM E8

standard.

(a)

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(b)

Fig. 4. (a) Effects of laser power and scan speed on densification rate and corresponding

OM microstructure of SLM-fabricated specimens; (b) 3D OM view of etched specimen

fabricated at laser power of 350 W and laser scan speed of 2.0 m/s.

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Fig. 5. XRD spectra of CNTs/Al-based composite powder and SLM-fabricated specimens

at various processing parameters.

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Fig. 6. Raman spectra of initial MWCNTs, ball milled CNTs/Al-based composite powder

and SLM-fabricated specimens at different processing parameters.

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(a) (b)

(c) (d)

Fig. 7. SEM image showing the hierarchical microstructure of specimen fabricated at P

= 350 W, v = 2.4 mm/s (a); Local magnification of area A in Fig. 7a revealing

microstructural difference of grains in three representative zones (b); High-magnification

SEM images showing the characteristic microstructures in (c) the coarse cellular zone

and transition zone and (d) the fine cellular zone.

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(a) (b)

(c) (d)

Fig. 8. SEM images showing the evolution of characteristic microstructures of SLM-

fabricated CNTs/Al-based specimens using different processing parameters: (a) P = 350

W, v = 1.8 m/s; (b) P = 350 W, v = 2.0 m/s; (c) P = 350 W, v = 2.2 m/s; and (d) P = 350

W, v = 2.4 m/s.

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Fig. 9. Characteristic microstructure of SLM-processed CNTs/Al-based specimen

fabricated at P = 350 W and v = 1.8 m/s (a); EDX mapping showing the elemental

distributions of (b) Al, (c) Si, (d) Mg and (e) C elements; Typical microstructures of SLM-

processed CNTs/Al-based specimen produced at P = 350 W and v = 2.0 m/s, observed at

a relatively low magnification (f) and a higher magnification (g).

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Fig. 10. Schematic of the formation mechanisms of different microstructures during laser

melting and solidification process.

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Fig. 11. Microhardness distributions and corresponding indentation microstructures of

SLM-fabricated specimens at different processing parameters.

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(a)

(b)

Fig. 12. Stress-strain curves (a) and tensile properties including ultimate tensile strength

(UTS) and elongation (b) of SLM-fabricated CNTs/Al-based specimens using different

processing parameters. Tensile properties of unreinforced AlSi10Mg specimen fabricated

at P = 350 W and v = 2.0 m/s are included for comparison.

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(a) (b)

(c) (d)

Fig. 13. SEM images showing the characteristic fracture surface morphologies of SLM-

fabricated CNTs/Al-based specimens at different processing parameters: (a) P = 350 W,

v = 1.8 m/s; (b) P = 350 W, v = 2.0 m/s; (c) P = 350 W, v = 2.2 m/s; (d) P = 350 W, v =

2.4 m/s.

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Table 1.

XRD data showing the displacement and intensity variation of the identified peaks for

Al9Si phase in SLM-processed specimens using different processing parameters.

Specimen 2θ

(deg.)

Intensity

(CPS)

FWHM

(deg.)

2θ

(deg.)

Intensity

(CPS)

FWHM

(deg.)

Standard (PDF#65-

8554)

38.591 44.681

CNTs/AlSi10Mg powder 38.601 4047 0.274 44.880 1750 0.298

1.8 m/s 38.621 1621 0.225 44.919 1266 0.235

2.0 m/s 38.757 1220 0.241 44.963 740 0.249

2.2 m/s 38.700 595 0.256 44.921 967 0.255

2.4 m/s 38.641 1109 0.276 44.901 1425 0.268

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Table 2.

EDX results showing chemical compositions collected in different positions in Fig. 9f.

Element Mass fraction (wt.%)

Spectrum A Spectrum B Spectrum C

Al 86.30 77.76 70.04

Si 10.76 10.92 25.97

Mg 0.30 0.28 0.29

C 2.64 11.04 3.70

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