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Laser additive manufacturing of carbon nanotubes (CNTs) reinforced aluminum matrix nanocomposites: Processing optimization, microstructure evolution and mechanical properties
Additive Manufacturing 29 (2019) 100801
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Full Length Article
Laser additive manufacturing of carbon nanotubes (CNTs) reinforced aluminum matrix nanocomposites: Processing optimization, microstructure evolution and mechanical properties
T
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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 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
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Additive manufacturing Selective laser melting Carbon nanotubes (CNTs) Aluminum matrix nanocomposites Mechanical properties
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 0.5 wt.% 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 SLMprocessed 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.
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 new-generation materials 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
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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 plasma sintering to fabricate multi-walled CNTs (MWCNTs) reinforced Al-based composite. They found that the yield strength of MWCNTs/Al composites increased substantially 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 partially reacted CNTs led to a significant improvement of interfacial strength and attendant load transfer efficiency in AMCs. Although
Corresponding author. E-mail address: [email protected] (D. Gu).
https://doi.org/10.1016/j.addma.2019.100801 Received 15 December 2018; Received in revised form 14 July 2019; Accepted 21 July 2019 Available online 23 July 2019 2214-8604/ © 2019 Elsevier B.V. All rights reserved.
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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. An “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 ntwo neighbouring 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).
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 SLM processing of CNTs/Al-based nanocomposites to establish a relationship among processing parameters, microstructures and properties.
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 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 was 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.
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).
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
2.2. SLM process An independently developed SLM system by NUAA was used in this study, mainly consisting of a YLR-500-WC ytterbium fiber laser with a maximum power of ˜500 W, 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 control system (Fig. 3a). When the 2
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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.
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
were generally over 99% at a relatively high laser power of 350 W. The corresponding OM images also reveal 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
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. 3
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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.
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 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 CeC 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
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, 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. 4
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Fig. 5. XRD spectra of CNTs/Al-based composite powder and SLM-fabricated specimens at various processing parameters.
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 EDX scan in the marked 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 nonequilibrium 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
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.
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 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 accumulation. 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 5
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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) CNTs/AlSi10Mg powder 1.8 m/s 2.0 m/s 2.2 m/s 2.4 m/s
38.591 38.601 38.621 38.757 38.700 38.641
4047 1621 1220 595 1109
0.274 0.225 0.241 0.256 0.276
44.681 44.880 44.919 44.963 44.921 44.901
1750 1266 740 967 1425
0.298 0.235 0.249 0.255 0.268
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 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.
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.
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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.
3.4. Mechanical properties
Table 2 EDX results showing chemical compositions collected in different positions in Fig. 9f.
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
Element
Al Si Mg C
Mass fraction (wt.%) Spectrum A
Spectrum B
Spectrum C
86.30 10.76 0.30 2.64
77.76 10.92 0.28 11.04
70.04 25.97 0.29 3.70
Fig. 9. Characteristic microstructure of SLMprocessed 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/Albased 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.
Fig. 11. Microhardness distributions and corresponding indentation microstructures of SLM-fabricated specimens at different processing parameters.
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 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/Albased 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
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.
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 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). These defects could be 8
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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.
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 CNTs/Al composites.
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 highstrength 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, 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.
Declaration of Competing 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
4. Conclusions (1) Laser power and scan speed were dominant factors to determine 9
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China, and the Top-Notch Young Talents Program of China.
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Full Length Article
Laser additive manufacturing of carbon nanotubes (CNTs) reinforced aluminum matrix nanocomposites: Processing optimization, microstructure evolution and mechanical properties
T
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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 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
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Additive manufacturing Selective laser melting Carbon nanotubes (CNTs) Aluminum matrix nanocomposites Mechanical properties
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 0.5 wt.% 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 SLMprocessed 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.
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 new-generation materials 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
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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 plasma sintering to fabricate multi-walled CNTs (MWCNTs) reinforced Al-based composite. They found that the yield strength of MWCNTs/Al composites increased substantially 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 partially reacted CNTs led to a significant improvement of interfacial strength and attendant load transfer efficiency in AMCs. Although
Corresponding author. E-mail address: [email protected] (D. Gu).
https://doi.org/10.1016/j.addma.2019.100801 Received 15 December 2018; Received in revised form 14 July 2019; Accepted 21 July 2019 Available online 23 July 2019 2214-8604/ © 2019 Elsevier B.V. All rights reserved.
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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. An “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 ntwo neighbouring 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).
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 SLM processing of CNTs/Al-based nanocomposites to establish a relationship among processing parameters, microstructures and properties.
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 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 was 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.
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).
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
2.2. SLM process An independently developed SLM system by NUAA was used in this study, mainly consisting of a YLR-500-WC ytterbium fiber laser with a maximum power of ˜500 W, 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 control system (Fig. 3a). When the 2
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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.
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
were generally over 99% at a relatively high laser power of 350 W. The corresponding OM images also reveal 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
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. 3
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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.
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 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 CeC 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
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, 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. 4
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Fig. 5. XRD spectra of CNTs/Al-based composite powder and SLM-fabricated specimens at various processing parameters.
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 EDX scan in the marked 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 nonequilibrium 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
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.
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 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 accumulation. 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 5
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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) CNTs/AlSi10Mg powder 1.8 m/s 2.0 m/s 2.2 m/s 2.4 m/s
38.591 38.601 38.621 38.757 38.700 38.641
4047 1621 1220 595 1109
0.274 0.225 0.241 0.256 0.276
44.681 44.880 44.919 44.963 44.921 44.901
1750 1266 740 967 1425
0.298 0.235 0.249 0.255 0.268
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 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.
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.
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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.
3.4. Mechanical properties
Table 2 EDX results showing chemical compositions collected in different positions in Fig. 9f.
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
Element
Al Si Mg C
Mass fraction (wt.%) Spectrum A
Spectrum B
Spectrum C
86.30 10.76 0.30 2.64
77.76 10.92 0.28 11.04
70.04 25.97 0.29 3.70
Fig. 9. Characteristic microstructure of SLMprocessed 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/Albased 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.
Fig. 11. Microhardness distributions and corresponding indentation microstructures of SLM-fabricated specimens at different processing parameters.
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 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/Albased 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
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.
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 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). These defects could be 8
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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.
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 CNTs/Al composites.
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 highstrength 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, 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.
Declaration of Competing 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
4. Conclusions (1) Laser power and scan speed were dominant factors to determine 9
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China, and the Top-Notch Young Talents Program of China.
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