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S136 metal matrix composites with tailored microstructures and mechanical properties
Materials and Design 175 (2019) 107811
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Selective laser melting of reduced graphene oxide/S136 metal matrix composites with tailored microstructures and mechanical properties Shifeng Wen a, Keyu Chen a, Wei Li a,c, Yan Zhou a,b,⁎, Qingsong Wei a, Yusheng Shi a a b c
State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China Faculty of Engineering, China University of Geosciences, Wuhan 430074, China School of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The fully dense RGO/S136 composites via molecular-level mixing were firstly fabricated by selective laser melting. • EBSD patterns revealed the refined grain and tailored microstructure for SLM RGO/S136 composites. • SLM RGO/S136 composites with 0.1 wt% RGO exhibited maximum tensile properties and hardness.
a r t i c l e
i n f o
Article history: Received 17 March 2019 Received in revised form 7 April 2019 Accepted 19 April 2019 Available online 28 April 2019 Keywords: Selective laser melting S136 steel Reduced graphene oxide (RGO) Microstructure Mechanical properties
a b s t r a c t In this work, a novel approach combining liquid deposition with selective laser melting (SLM) is used for fabricating reduced graphene oxide (RGO)/S136 metal matrix composites (MMCs). The grain sizes, crystallographic textures, phase compositions and mechanical properties can be tailored by controlling the RGO content in the RGO/S136 MMCs. The results show that the average grain size reaches its smallest size of 0.75 μm when 0.1 wt % RGO was added to the RGO/S136 MMCs. As the RGO content is increased from 0 wt% to 0.5 wt%, a continuous transition of the grains from the (001) orientation to the (101) and (111) orientations is observed. In addition, the cellular dendritic grains transform into equiaxed fine grains with increasing RGO content. The SLMprepared RGO/S136 MMCs are dominated by high-angle grain boundaries (˃15°) and the martensite (bcc) phase. The hardness, ultimate tensile strength and yield strength of the SLM RGO/S136 MMCs exhibit trends that initially increase and then decrease, with maximum values of 580.6 HV, 535.3 MPa and 515.8 MPa, respectively. This paper highlights the possibility of controlling the RGO content to achieve the desired microstructural characteristics and mechanical properties of RGO/S136 MMCs fabricated by the SLM process. © 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).
1. Introduction
⁎ Corresponding author at: State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China. E-mail address: [email protected] (Y. Zhou).
Recently, selective laser melting (SLM) technique has attracted great attention from researchers because of its high material utilization rate, short manufacturing cycle, and the ability to form metal parts with complex shapes [1-3]. A typical process of SLM technique is to design the component using a three-dimensional computer-assisted design
https://doi.org/10.1016/j.matdes.2019.107811 0264-1275/© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).
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method, follow by melting the material through a high-energy laser beam and accumulate it layer-by-layer to form a desired threedimensional entity. S136 is a kind of martensitic carbon steel widely used in the plastic mould industry. It is highly valued owing to its excellent corrosion resistance, which can tolerate the harsh corrosion environment of plastic mould processing [4,5]. However, extremely high mould pressures and impact loads, as well as frequent work times, lead to common failure modes in the plastic moulds, such as surface wear, deformation and fracturing. To improve the mould life, mould materials are required to have high tensile strength, hardness and wear resistance [6–8]. RGO, which exhibits a very high tensile strength (130 GPa), Young's modulus (0.5–1 TPa) and thermal conductivity (5000 Wm−1 k−1) [9] while being chemically inert (not easily reacted with the metal matrix) [10], has been considered an ideal reinforcement material for producing metal matrix composites (MMCs) [11–14]. Several works have been reported on the fabrication of RGOreinforced MMCs (such as Al, Cu and Mg) by powder metallurgy, spark plasma sintering, laser sintering, hot extrusion, selective laser melting and so on [15–21]. Li et al. fabricated RGO/Al MMCs via a composite powder assembly approach, after which the stiffness and tensile strength were significantly improved (~50% enhancement over the Al matrix) [22]. Hwang et al. innovatively adopted a molecular-level mixing process and spark plasma sintering (SPS) to fabricate RGO/Cu MMCs. The elastic modulus and yield strength of this material were 131 GPa and 284 MPa, respectively, which were 1.3 and 1.8 times higher than those of pure Cu [23]. Xiang et al. prepared graphene nanoplatelet (GNP)-reinforced Mg6Zn MMCs and the mechanical properties of the Young's modulus, YS and UTS were improved by 20%, 166% and 35%, respectively, compared to those of the base material [24]. However, to the best of our knowledge, no research regarding SLM RGO/S136 MMCs has been reported. Therefore, the present study aims to investigate the feasibility of RGO-reinforced S136 stainless steel prepared by SLM for the first time.
However, due to the unique monoatomic layered structure of RGO, aggregation and entanglement are likely to occur in RGO MMCs [25,26]. Herein, a novel process combining molecular-level mixing techniques based on liquid deposition methods [27] and SLM techniques is used to investigate the effect of RGO content on the microstructures and mechanical properties of RGO/S136 MMCs and is presented in this study. In this paper, the RGO content in RGO/S136 MMCs is optimized; that is, the desired microstructural features and mechanical properties are achieved by tailoring the RGO content. 2. Experimental 2.1. Fabrication of RGO/S136 mixed powders The gas-atomized S136 steel powder (Changsha Hualiu Metallurgy Powder Co., Ltd., China) exhibited a spherical or nearly spherical shape with an average particle size of 15 μm, as shown in Fig. 1a and b. Its chemical composition is shown in Table 1. The quality and morphology of the RGO (Arke Inc.) are shown in Fig. 1c and d. The RGO was a single layer of graphene approximately 6–8 nm in thickness and 5 μm across. Sodium dodecyl benzene sulfonate (SDBS) with a purity of no less than 90% was obtained from Tianjin Kaitong Chemical Reagent Co., Ltd. Several steps were required for preparing the RGO/S136 mixed powders, and a schematic diagram of the process is shown in Fig. 2. In step 1, 1 g of RGO and 0.5 g of SDBS were placed into a beaker, which was then filled with 5000 ml of deionized water followed by mechanical stirring and ultrasonication for 40 min. Then, S136 powder was added to the suspension prepared above, and the container was placed onto a thermostatic heating platform at a constant temperature of 80 °C. During the entirety of step 2, ultrasonic stirring was maintained until the deionized water had completely evaporated. Finally, a RGO/S136 powder mixture was obtained by filtration, dried in a vacuum drying oven
Fig. 1. (a) SEM image of the S136 powder, (b) the distribution of the S136 particle size, (c) SEM image showing the morphology of RGO and (d) the RGO under high magnification.
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Table 1 Chemical composition of the S136 steel powder (wt%). Si
Mn
Cr
V
C
O
P
S
Fe
0.96
0.98
13.55
0.4
0.29
0.078
0.01
/
Bal.
(Model DZF-1, Shanghai, China) at 90 °C for 24 h and then sieved through a 325-mesh sieve for further use. 2.2. SLM process An SLM Solutions instrument comprising a 400 W IPG fibre laser and an inert argon gas protection system was used in this work. The 0.2 wt% RGO/S136 mixed powder was chosen, and 20 groups of SLM parameters were set to form the samples. A diagram of the relationship between the density and line energy density is shown in Fig. 3 and can be divided into region 1, region 2 and region 3 according to the different line energy densities. A typical optical micrograph of the corresponding sample was inset in each region. There were numerous voids on the surface of the sample in region 1 and an abundance of cracks on the surface of the sample in region 3. However, the sample in region 2 had fewer surface defects, which determined that the optimum combination of process parameters had been obtained in region 2. Therefore, in this study, the laser scanning power and speed were taken as 200 W and 900 mm/s, respectively, in an argon atmosphere with pre-heating to 200 °C. The laser scanning strategy and scanning direction of the samples are depicted in Fig. 4a. The cube-shaped samples, with dimensions of 10 mm × 10 mm × 10 mm and 5 mm × 5 mm × 5 mm, and tensile strength samples (mm) are shown in Fig. 4b. The samples were cut from the stainless steel substrate after fabrication by a wire electrical discharge machine. For comparison, pure S136 samples were also prepared by the same parameters. For simplicity, the 0, 0.1, 0.2 and 0.5 wt % RGO/S136 samples were labelled A1, A2, A3 and A4, respectively. 2.3. Characterization The densities of the samples were determined by the Archimedes method. Each set of samples was measured three times, and the results
Fig. 3. Relationship between density and the line energy density.
were averaged as the final density. The samples were ground with sandpaper from 240 mesh to 2000 mesh, followed by polishing with the diamond suspension (Lab Testing Technology Co.,Ltd. Shanghai) for metallographic testing. The defects such as cracks of the polished samples were observed under an optical microscopy (OM, Olympus BX60). The corrosive solution consisted of HCl (15 ml), H2O (60 ml) and FeCl3 (5 g). The microstructure of the samples and the fracture of the tensile specimens were observed under a FEI Sirion Quanta 200 scanning electron microscope (SEM). An X-ray diffraction pattern having a diffraction angle (2θ) ranging from 30° to 100° and a step size of 10°/min was obtained using an XRD-7000s (Shimadzu, Japan) X-ray diffractometer (CuKα radiation, a tube voltage of 40 KV and a tube current of 40 mA) in the RGO/S136 composites. The electron backscatter diffraction images were obtained by an HKL Nordlys orientation imaging microscope system (Oxford, Oxford Instruments, United Kingdom) with a step size of 1 μm. The acquired data is processed by HKL Channel 5 software to obtain images such as grain size and grain boundary on the surface of the samples. Five-point sampling method was used to select test points and the Vickers
Fig. 2. Schematic illustration of the procedure for preparing the mixed powder.
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Fig. 4. Schematic of (a) laser scanning strategy and (b) tensile strength sample.
3. Results and discussion 3.1. Phase composition and microstructure Typical XRD spectra of the SLM RGO/S136 composites obtained within a wide 2θ range (30–100°) are depicted in Fig. 5. The main phases in the composites were the α-Fe and γ-Fe phases. The composites with different RGO content have the same phase, and no obvious second phase is detected, which indicates the dominant martensitic transformation and austenitic transformation during the high temperature molten pool reaction of the SLM process. The α-Fe diffraction peak of the SLM samples slightly shifted to the left with increasing RGO content. According to Bragg's law [28]: 2d sinθ ¼ nλ ðn ¼ 1; 2; 3; …Þ Fig. 5. XRD analysis of the A1, A2, A3 and A4 samples.
hardness was measured at room temperature using a Vickers hardness tester (Wolpert Wilson Instruments 432 SVD, China), and the average value was taken as the final hardness. The fixed loading force was 3 kg and the loading time was 15 s. The Zwick/Roell Z020 (Germany) was used to completely pull off the sample at a load of 90 KN and a strain rate of 1 mm/min at room temperature. Each group of tensile samples was tested 5 times, and finally, the average yield strength and ultimate tensile strength were calculated in this study.
ð1Þ
the observed decrease of the 2θ value indicated an increase in the lattice plane distance d at higher RGO contents. This increase in lattice plane distance was primarily based on lattice distortion caused by martensitic transformation, as well as microscopic volume expansion and residual stress at grain boundaries [29,30]. Additionally, a portion of the C atoms in the alloy entered the Fe lattice as a solid solution and formed ferrite in the Fe matrix [31,32], leading to the change in the lattice parameters. However, there were no C peaks detected in and of the RGO/S136 composites due to the sensitivity of the XRD apparatus. Fig. 6a–d shows the fcc/bcc phase distribution obtained by the EBSD test. The results of the fcc/bcc phase ratios are displayed in Fig. 6e, which were calculated to be 31.8%, 1.3%, 42.9% and 58.5% for the A1, A2, A3 and
Fig. 6. (a)–(d) The phase distribution of the A1, A2, A3 and A4 samples and (e) the ratios of the fcc phase to the bcc phase corresponding to the A1, A2, A3 and A4 samples.
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A4 samples, respectively. These results were consistent with the XRD data shown in Fig. 5. Generally, the fcc phase is the austenite phase, and the bcc phase is the martensite phase in iron-based composites [4]. Martensite dominated the phase composition, reaching a maximum of 98.75% in the A2 sample, which contained 0.1 wt% RGO. Most of the fcc grains exhibited a high aspect ratio, and they were located preferentially along the boundaries between the bcc grains [33]. SEM images of the RGO/S136 samples with different RGO contents are presented in Fig. 7. The grains of the pure S136 sample mainly included equiaxed fine grains and cellular dendritic grains, as shown in Fig. 7a, which was consistent with the results reported by Zhou et al. [34]. The RGO/S136 samples were dominated by equiaxed fine grains, and a few columnar grains were found, as shown in Fig. 7b–d, which was different from that of the pure S136 alloy. The different microstructures indicated an obvious transition of the grains from cellular dendritic grains to equiaxed fine grains. The transformation to equiaxed fine grains in the matrix can be attributed to (i) the extremely fast cooling rate of the SLM process provides a shorter growth time; (ii) the RGO dispersed at the grain boundaries provides a large number of nanoscale nucleation sites; (iii) RGO pinned at the grain boundaries retards further grain growth. As Martin et al. [35] described, although the extremely high cooling rate during the SLM process favors equiaxed growth of the grain, heterogeneous nucleation induced by nano-sized RGO particles as nucleation sites is the key to promoting a columnarto-equiaxed transformation [36]. 3.2. Crystal orientation and crystallographic texture To study the effect of the RGO content on the grain size, orientation and texture of the RGO/S136 MMCs, EBSD measurements were carried out on the top surface of the samples (A1, A2, A3 and A4). The
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representative orientation maps and inversed pole figure (IPF) of the RGO/S136 composites are shown in Fig. 8. All the IPF maps in this paper use the same colour coding, with the red, green, and blue colour levels being proportional to the three basic orientations of (001), (101), and (111), respectively. The orientation map of pure S136 (A1) shows that the grains mainly possessed a red-green-blue colour, as shown in Fig. 8a. Combined with the IPF, it could be determined that the pure S136 sample possessed a combination of the (001), (101) and (111) orientations. There were no significant differences between the areas corresponding to the three colours, which indicated that the pure S136 alloy had no obvious grain orientation. When 0.1 wt% RGO was added to the S136 alloy (A2), the grains of the composite exhibited a decrease in the (100) orientation and increased in the (101) and (111) orientations, as shown in Fig. 8b. When the RGO content was further increased to 0.2 wt% and 0.5 wt%, the orientation maps (Fig. 8c–d) were dominated by a green-blue colour, which indicated a strong (101) and (111) orientation in the corresponding composites. Based on the above observations, it could be concluded that the RGO content had a significant effect on the grain orientation of the RGO/S136 composites and promoted a transition of the grains towards the (101) and (111) orientations as the RGO content increased. Additionally, inhibition of the (100) grain orientation occurred when RGO was added. The orientation maps also provided meaningful information regarding the changes in the average grain size. The Channel 5 software was utilized to quantitatively analyse the average grain size [37]. The average grain size of the pure S136 alloy was calculated to be 0.86 μm, and the grain sizes were distributed over the range of 0.5 μm to 6.5 μm. When 0.1 wt% RGO was added to the alloy, the grains were refined and the average grain size was calculated to be 0.75 μm. However, some of the grains tended to become larger as the maximum grain size increased from 6.5 μm to 9.5 μm. Furthermore, the average grain
Fig. 7. SEM images of the A1, A2, A3 and A4 samples.
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Fig. 8. (a)–(d) EBSD orientation maps of the A1, A2, A3 and A4 samples and (e) the inverse pole figure (IPF).
sizes were 1.46 μm and 1.91 μm for the 0.2 wt% and 0.5 wt% RGO/S136 samples, respectively. The results showed that the grains gradually coarsened with increasing RGO content. Fig. 9 illustrates the grain boundary misorientation angle distribution obtained from the top surface of the A1, A2, A3 and A4 samples. Generally, grain boundaries for which the misorientation angle of the crystal is greater than 15° are defined as high-angle grain boundaries (HAGBs) and they are defined as low-angle grain boundaries (LAGBs) when the misorientation angle is less than 15°. The boundaries are colour-coded with LAGBs in green and HAGBs in black [38]. Fig. 10 provided a more intuitive regular distribution pattern of the grain boundary misorientation angles. The linear fractions of HAGBs in this region were calculated to be 69.86%, 72.07%, 73.92% and 34.98% for the A1, A2, A3 and A4 samples, respectively. HAGBs accounted for the majority of misorientation angles when the mass fraction of RGO in the composites was less than 0.2%. LAGBs dominated as the mass fraction of RGO increased to 0.5%.
Generally, the toughness of a material is closely related to the proportion of HAGBs in its structure. HAGBs can hinder crack propagation, i.e., for a higher HAGB fraction, the toughness of the material is better. In addition, the higher the misorientation angle of adjacent grain boundaries is, the stronger the ability of the material to resist crack propagation. Therefore, it can be predicted that the 0.5 wt% RGO/S136 sample may have the lowest toughness value among the four samples, by only considering the effect of grain boundary misorientation. The transformation mechanism between LAGBs and HAGBs is closely related to the uniqueness of the SLM process. The RGO/S136 MMC powders were added layer by layer to fabricate the MMCs by a high-energy laser beam. Along with the rapid melting and solidification of the SLM process, a large number of sub-grains nucleate and grow. The grain boundaries of these sub-grains tend to form LAGBs first. Additionally, the re-melting phenomenon in the SLM process is similar to the temper heat treatment in the conventional casting process, which results in recrystallization and an increase in the amount of HAGBs. Due
Fig. 9. Grain boundary misorientation angles of the A1, A2, A3 and A4 samples.
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Fig. 10. Grain boundary misorientation angle distributions of the A1, A2, A3 and A4 samples.
to the extremely short cooling time of the SLM process, the LAGBs are not fully converted into HAGBs and become residual LAGBs, consequently remaining in the sample [39,40]. As shown in Fig. 8, there were large differences in the grain orientations between the SLM samples with different mass fractions of RGO.
Therefore, the content of RGO had a significant influence on the local crystallographic texture of the SLM RGO/S136 composites. The most common (001) pole figures (PFs) from the top view of the SLM RGO/ S136 samples are shown in Fig. 11. Fig. 11a–d illustrates the (001) pole figures of the basic bcc phase in the MMC matrix, the
Fig. 11. (001) pole figures (PFs) of the A1, A2, A3 and A4samples. (a)–(d) The bcc phase PFs and (e)–(h) the fcc phase PFs.
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accompanying data corresponding to the fcc phase are shown in Fig. 11e–f. For the SLM of the pure S136 sample, a weaker (001) texture along the scanning direction was observed, as shown in Fig. 11a. The relatively low intensities in the PF suggested the presence of a fibrous texture in the tissue [41]. The texture index (TI) is usually induced to quantify the fibrous texture intensity. TI is obtained from the orientation function (ODF) f(g) by Eq. (2) [42]: Z Texture Index ¼ TI ¼
2
eulerspace
ð f ðgÞÞ dg
ð2Þ
where f is the orientation distribution as a function of the Euler space coordinates g. The texture indexes of the bcc phase corresponding to the S136 alloy and the 0.1 wt%, 0.2 wt% and 0.5 wt% RGO/S136 composites, were determined to be 7.17, 5.68, 9.79 and 12.4, respectively. The texture indexes of the fcc phase were determined to be 27.99, 20.02, 30.98 and 33.24, respectively. Therefore, the fcc phase showed a stronger (001) orientation than that of the bcc phase. Considering the results of the texture indexes of the fcc and the bcc phases comprehensively, it can be seen that they were in good agreement with the EBSD orientation maps shown in Fig. 8. Fig. 12 shows the proportions of the Schmidt factors of different values, corresponding to the A1, A2, A3 and A4 samples. The inset in each figure indicates the different regions of the sample surface corresponding to the different Schmidt factor values. The Schmidt factors of the A1, A2, A3 and A4 samples were calculated to be 0.4442, 0.4573, 0.4039 and 0.4484, respectively. It is well known that the plastic deformation ability of a material can be characterized by its Schmidt factor [43]. Hence, the Schmidt factor is often utilized for evaluating the yield strength (σs) of a material, which can be calculated as follows: σ s ¼ τs
1 cosλ cosφ
ð3Þ
where λ is the angle between the applied load F and the 〈110〉 orientation, φ is the angle between the applied load F and normal direction of {111}, and τs is the critical shear stress. The product of cos λ and cos φ is called the Schmidt factor [44]. Based on Eq. (3), a lower Schmidt factor usually leads to a higher σs value. The Schmidt factor of A3 was the smallest of all samples, as seen in Fig. 12, indicating that A3 may have the highest yield strength. 3.3. Mechanical properties of the SLM RGO/S136 composites The Vickers micro-hardness values of the A1, A2, A3 and A4 samples are shown in Fig. 13. The Vickers micro-hardness of the pure S136 sample was approximately 544.1 HV, which was similar to the hardness value reported by Wen et al. [4]. The maximum Vickers microhardness value of the RGO/S136 composites was 580.6 HV, reaching the peak value after the addition of 0.1 wt% RGO, which was 6.3% higher than that of the pure S136 sample. Fig. 13 also indicates that the Vickers micro-hardness decreased with the further addition of RGO, for which the Vickers micro-hardness values of A3 and A4 were measured to be 534.6 HV and 468.7 HV, respectively. Nevertheless, the SLM RGO/S136 MMCs, for all the given RGO contents, demonstrated superior Vickers micro-hardness values compared to that of commercially used S136, which has a typical hardness of 380–400 HV. In general, residual stress generated during the SLM process is unavoidable, and excessive residual stress can cause local failure of components. However, it should be pointed out that in a high-density metal matrix composite with few cracks and voids (Fig. 3), a reasonable level of residual stress is conducive to the improvement of hardness. Moreover, the significant grain refinement effect (Fig. 8) due to the inclusion of nanoscale RGO favoured a further increase in the obtained microhardness. Fig. 14a shows the stress-strain curves of the RGO/S136 composites, and the trends in the yield strength and ultimate tensile strength for different RGO contents is described in Fig. 14b. Both the yield strength and ultimate tensile strength showed an initially increasing trend and then
Fig. 12. Distribution histograms of the Schmidt factors of the A1, A2, A3 and A4 samples.
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Table 2 Mechanical properties of samples A1, A2, A3 and A4. Sample number A1 A2 A3 A4
Fig. 13. Vickers hardness values of the surfaces of the A1, A2, A3 and A4 samples.
decreased with increasing RGO content. The maximum yield strength and ultimate tensile strength appeared in 0.1 wt% RGO/S136, increasing by 4.1% and 30.4%, respectively, compared with those of the pure S136 sample. Tensile properties of the composites decreased as the RGO content increased to 0.2 wt% and 0.5 wt%. Importantly, the elongation remained essentially unchanged. The σs of A1, A2, A3 and A4 are shown in Table 2 and were measured to be 359.2 ± 12.7 MPa, 515.8 ± 15.3 MPa, 259.2 ± 19.1 MPa and 275.9 ± 15.8 MPa, respectively. The results indicated that the effects of the Schmidt factor on the yield strength were no longer the main factor for the SLM RGO/S136 samples. The strengthening mechanism of RGO in the S136 matrix can be explained by the following aspects: stress transfer, dislocation strengthening and grain refinement [45–47]. The stress transfer mainly depended on the interfacial bonding between the S136 matrix and RGO [48]. The large mismatch in the thermal expansion coefficients of S136 and RGO formed dislocations [49]. The hard RGO phase impeded the movement of these dislocations, leading to an increase in the dislocation density and resulting in dislocation strengthening [50]. The variation in average grain size is reflected by the EBSD results in Fig. 8, and the average grain size decreased from 0.86 μm to 0.75 μm in the pure S136 and 0.1 wt% RGO/S136 samples, respectively. The pinning effect of RGO was remarkably beneficial to the dispersion of the grain boundaries, which prevented grain growth and reduced the crack propagation properties to some extent. Therefore, RGO as a reinforcement could improve the yield strength and ultimate tensile strength of the S136 matrix. However, the tensile properties decreased significantly as the RGO content
Yield strength σs (MPa) 359.2 ± 12.7 515.8 ± 15.3 259.2 ± 19.1 275.9 ± 15.8
Tensile strength σb (MPa)
Elongation (%)
513.6 ± 20.3 535.3 ± 18.5 339.4 ± 17.2 311.5 ± 21.4
4.4 ± 0.3 4.3 ± 0.8 4.4 ± 1.2 3.7 ± 0.5
increased further. There were two main reasons accounting for these results. On the one hand, agglomeration inevitably occurred with increasing RGO content, which resulted in the degradation of the pinning effect and resulted in the grain sizes becoming coarser [51]. On the other hand, RGO may partially react with the S136 matrix under the high temperature of the laser beam and lose its unique strengthening properties. The fractured surfaces of the RGO/S136 composites are displayed in Fig. 15. The fractured surface of pure S136 (Fig. 15a) showed a brittle fracture, but a small number of dimples could be found, as marked by the yellow arrow in Fig. 15b. For the 0.1 wt% RGO/S136 composite, a very small amount of RGO can be observed in Fig. 15c, and numerous dimples were observed at the interface between the RGO and S136 matrix in Fig. 15e. In Fig. 15d, it can also be seen that RGO possessed a significant number of fractures and holes. The fractured surface of the 0.2 wt% RGO/S136 composite is shown in Fig. 15f–g. The areas marked by the arrows in Fig. 15f had typical curving and wrinkling structures from the removal of RGO. The unique honeycombed structure of RGO was observed in Fig. 15g, further proving the presence of RGO. At the same time, several regular hexagonal components underwent significant deformation and elongation under stress. There were cracks at the junction of the RGO and S136 matrix, as shown in Fig. 15h, indicating that the interfaces between the RGO and S136 matrix had not been well combined. Different shrinkages along the interfaces during the cooling process may be the reason for this phenomenon [52]. 4. Conclusions RGO/S136 composites were successfully fabricated by a liquid deposition method and SLM process. The effects of the RGO content (0, 0.1 wt %, 0.2 wt% and 0.5 wt%) on the microstructures, phase compositions and mechanical properties of the RGO/S136 composites were investigated. The results obtained can be summarized as follows: 1. The RGO content showed a significant effect on the grain orientation of RGO/S136 composites, promoting a transition of the grain orientation from a combination of the (001), (101), and (111) orientations to the (101) and (111) orientations as the RGO content increased.
Fig. 14. (a) Stress-strain curve of the RGO/S136 composites and (b) relationships of the ultimate tensile strength and yield strength with different RGO contents.
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Fig. 15. Fractured surfaces of the RGO/S136 composites. (a)–(b) Pure S136, (c)–(e) 0.1 wt% RGO/S136, (f)–(g) 0.2 wt% RGO/S136 and (h) 0.5 wt% RGO/S136.
2. The RGO content was confirmed to have a refinement effect on the average grain size, as the grain size was 0.75 μm with 0.1 wt% RGO, which was smaller than the grain size of 0.85 μm in the pure S136 alloy. 3. The phases in the RGO/S136 composites mainly consisted of the α-Fe and γ-Fe phases. The addition of RGO induced a transition of the grains from cellular dendritic grains to equiaxed fine grains.
4. With increasing RGO content, the hardness, ultimate tensile strength and yield strength of the composites initially increased and then decreased. The microhardness, tensile strength and yield strength reached the maximum of 580.6 ± 15.4 HV, 535.3 ± 18.5 MPa and 515.8 ± 15.3 MPa, respectively, when the RGO content was 0.1 wt %. The strengthening mechanism in the RGO/S136 composites was mainly attributed to stress transfer, dislocation strengthening and grain refinement.
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CRediT authorship contribution statement Shifeng Wen: Writing - original draft. Keyu Chen: Data curation, Investigation. Wei Li: Methodology. Yan Zhou: Funding acquisition, Formal analysis. Qingsong Wei: Supervision. Yusheng Shi: Writing - review & editing. Acknowledgements This study was supported by National Hi-tech R&D Program of China (863 Program, Grant No. 2015AA042501), the National Natural Science Foundation of China (Grant No. 51605176), Hubei Provincial Natural Science Foundation of China (Grant No. 2018CFB502), Guangdong Provincial Technology Major Project of China (Grant No. 2017B090911007), State Key Laboratory of Materials Processing and Die and Mould Technology, Huazhong University of Science and Technology (Grant No. P2019-006), and Engineering Research Center of Rock-Soil Drilling & Excavation and Protection, Ministry of Education (Grant No. 201804). References [1] Y. Wang, J. Shi, S. Lu, Y. Wang, Selective laser melting of graphene-reinforced Inconel 718 superalloy: evaluation of microstructure and tensile performance, J. Manuf. Sci. Eng. 139 (2016) (041005–041005-6). [2] T. Kimura, T. Nakamoto, Microstructures and mechanical properties of A356 (AlSi7Mg0.3) aluminum alloy fabricated by selective laser melting, Mater. Des. 89 (2016) 1294–1301. [3] P. Han, Additive design and manufacturing of jet engine parts, Engineering 3 (2017) 648–652. [4] S. Wen, H. Hu, Y. Zhou, Z. Chen, Q. Wei, Y. Shi, Enhanced hardness and wear property of S136 mould steel with nano-TiB2 composites fabricated by selective laser melting method, Appl. Surf. Sci. 457 (2018) 11–20. [5] S. Wen, X. Ji, Y. Zhou, C. Han, Q. Wei, Y. Shi, Corrosion behavior of the S136 mold steel fabricated by selective laser melting, Chinese J. Mech. Eng. 31 (2018) 108. [6] B. He, Reseach on the failure and material selection of plastic mold, Procedia Eng 23 (2011) 46–52. [7] X. Li, L. Duan, J. Li, X. Wu, Experimental study and numerical simulation of dynamic recrystallization behavior of a micro-alloyed plastic mold steel, Mater. Des. 66 (2015) 309–320. [8] W. Li, H. Yuan, Z. Sun, Z. Zhang, Surface vs. interior failure behaviors in a structural steel under gigacycle fatigue: failure analysis and life prediction, Int. J. Fatigue 64 (2014) 42–53. [9] Z. Chen, L. Dong, D. Yang, H. Lu, Superhydrophobic graphene-based materials: surface construction and functional applications, Adv. Mater. 25 (2013) 5352–5359. [10] H. Bi, I. Chen, T. Lin, F. Huang, A new tubular graphene form of a tetrahedrally connected cellular structure, Adv. Mater. 27 (2015) 5943–5949. [11] A.K. Geim, K.S. Novoselov, The rise of graphene, Nat. Mater. 6 (2007) 183–191. [12] C. Lee, X. Wei, J.W. Kysar, H. James, Measurement of the elastic properties and intrinsic strength of monolayer graphene, Science 321 (2008) 385–388. [13] S.C. Tjong, Recent progress in the development and properties of novel metal matrix nanocomposites reinforced with carbon nanotubes and graphene nanosheets, Mater. Sci. Eng. R Reports. 74 (2013) 281–350. [14] A. Nieto, A. Bisht, D. Lahiri, C. Zhang, A. Agarwal, Graphene reinforced metal and ceramic matrix composites: a review, Metall. Rev. 62 (2016) 241–302. [15] S. Feng, G. Qiang, L. Zan, G. Fan, Z. Li, D. Xiong, Y. Su, Z. Tan, Z. Jie, Z. Di, Strengthening and toughening mechanisms in graphene-Al nanolaminated composite micropillars, Acta Mater. 125 (2017) 98–108. [16] Y. Ming, W. Lin, H. Zhu, T. Fan, Z. Di, Simultaneously enhancing the strength, ductility and conductivity of copper matrix composites with graphene nanoribbons, Carbon N. Y. 118 (2017) 250–260. [17] W.J. Kim, T.J. Lee, S.H. Han, Multi-layer graphene/copper composites: preparation using high-ratio differential speed rolling, microstructure and mechanical properties, Carbon N. Y. 69 (2014) 55–65. [18] L. Chen, J. Xu, H. Choi, M. Pozuelo, X. Ma, S. Bhowmick, J. Yang, S. Mathaudhu, X. Li, Processing and properties of magnesium containing a dense uniform dispersion of nanoparticles, Nature 528 (2015) 539. [19] K. Nie, X. Wang, K. Wu, L. Xu, M. Zheng, X. Hu, Processing, microstructure and mechanical properties of magnesium matrix nanocomposites fabricated by semisolid stirring assisted ultrasonic vibration, J. Alloy. Compd. 509 (2011) 8664–8669. [20] Y. Huang, P. Bazarnik, D. Wan, D. Luo, P.H.R. Pereira, M. Lewandowska, J. Yao, B.E. Hayden, T.G. Langdon, The fabrication of graphene-reinforced Al-based nanocomposites using high-pressure torsion, Acta Mater. 164 (2019) 499–511. [21] L. Wang, C. Tian, S. Wang, Microstructural characteristics and mechanical properties of carbon nanotube reinforced AlSi10Mg composites fabricated by selective laser melting, Opt. - Int. J. Light Electron Opt. 143 (2017) 173–179. [22] Z. Li, Q. Guo, Z. Li, G. Fan, D. Xiong, Y. Su, J. Zhang, D. Zhang, Enhanced mechanical properties of graphene (reduced graphene oxide)/aluminum composites with a bioinspired nanolaminated structure, Nano Lett. 15 (2015) 8077–8083.
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Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Selective laser melting of reduced graphene oxide/S136 metal matrix composites with tailored microstructures and mechanical properties Shifeng Wen a, Keyu Chen a, Wei Li a,c, Yan Zhou a,b,⁎, Qingsong Wei a, Yusheng Shi a a b c
State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China Faculty of Engineering, China University of Geosciences, Wuhan 430074, China School of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The fully dense RGO/S136 composites via molecular-level mixing were firstly fabricated by selective laser melting. • EBSD patterns revealed the refined grain and tailored microstructure for SLM RGO/S136 composites. • SLM RGO/S136 composites with 0.1 wt% RGO exhibited maximum tensile properties and hardness.
a r t i c l e
i n f o
Article history: Received 17 March 2019 Received in revised form 7 April 2019 Accepted 19 April 2019 Available online 28 April 2019 Keywords: Selective laser melting S136 steel Reduced graphene oxide (RGO) Microstructure Mechanical properties
a b s t r a c t In this work, a novel approach combining liquid deposition with selective laser melting (SLM) is used for fabricating reduced graphene oxide (RGO)/S136 metal matrix composites (MMCs). The grain sizes, crystallographic textures, phase compositions and mechanical properties can be tailored by controlling the RGO content in the RGO/S136 MMCs. The results show that the average grain size reaches its smallest size of 0.75 μm when 0.1 wt % RGO was added to the RGO/S136 MMCs. As the RGO content is increased from 0 wt% to 0.5 wt%, a continuous transition of the grains from the (001) orientation to the (101) and (111) orientations is observed. In addition, the cellular dendritic grains transform into equiaxed fine grains with increasing RGO content. The SLMprepared RGO/S136 MMCs are dominated by high-angle grain boundaries (˃15°) and the martensite (bcc) phase. The hardness, ultimate tensile strength and yield strength of the SLM RGO/S136 MMCs exhibit trends that initially increase and then decrease, with maximum values of 580.6 HV, 535.3 MPa and 515.8 MPa, respectively. This paper highlights the possibility of controlling the RGO content to achieve the desired microstructural characteristics and mechanical properties of RGO/S136 MMCs fabricated by the SLM process. © 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).
1. Introduction
⁎ Corresponding author at: State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China. E-mail address: [email protected] (Y. Zhou).
Recently, selective laser melting (SLM) technique has attracted great attention from researchers because of its high material utilization rate, short manufacturing cycle, and the ability to form metal parts with complex shapes [1-3]. A typical process of SLM technique is to design the component using a three-dimensional computer-assisted design
https://doi.org/10.1016/j.matdes.2019.107811 0264-1275/© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).
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method, follow by melting the material through a high-energy laser beam and accumulate it layer-by-layer to form a desired threedimensional entity. S136 is a kind of martensitic carbon steel widely used in the plastic mould industry. It is highly valued owing to its excellent corrosion resistance, which can tolerate the harsh corrosion environment of plastic mould processing [4,5]. However, extremely high mould pressures and impact loads, as well as frequent work times, lead to common failure modes in the plastic moulds, such as surface wear, deformation and fracturing. To improve the mould life, mould materials are required to have high tensile strength, hardness and wear resistance [6–8]. RGO, which exhibits a very high tensile strength (130 GPa), Young's modulus (0.5–1 TPa) and thermal conductivity (5000 Wm−1 k−1) [9] while being chemically inert (not easily reacted with the metal matrix) [10], has been considered an ideal reinforcement material for producing metal matrix composites (MMCs) [11–14]. Several works have been reported on the fabrication of RGOreinforced MMCs (such as Al, Cu and Mg) by powder metallurgy, spark plasma sintering, laser sintering, hot extrusion, selective laser melting and so on [15–21]. Li et al. fabricated RGO/Al MMCs via a composite powder assembly approach, after which the stiffness and tensile strength were significantly improved (~50% enhancement over the Al matrix) [22]. Hwang et al. innovatively adopted a molecular-level mixing process and spark plasma sintering (SPS) to fabricate RGO/Cu MMCs. The elastic modulus and yield strength of this material were 131 GPa and 284 MPa, respectively, which were 1.3 and 1.8 times higher than those of pure Cu [23]. Xiang et al. prepared graphene nanoplatelet (GNP)-reinforced Mg6Zn MMCs and the mechanical properties of the Young's modulus, YS and UTS were improved by 20%, 166% and 35%, respectively, compared to those of the base material [24]. However, to the best of our knowledge, no research regarding SLM RGO/S136 MMCs has been reported. Therefore, the present study aims to investigate the feasibility of RGO-reinforced S136 stainless steel prepared by SLM for the first time.
However, due to the unique monoatomic layered structure of RGO, aggregation and entanglement are likely to occur in RGO MMCs [25,26]. Herein, a novel process combining molecular-level mixing techniques based on liquid deposition methods [27] and SLM techniques is used to investigate the effect of RGO content on the microstructures and mechanical properties of RGO/S136 MMCs and is presented in this study. In this paper, the RGO content in RGO/S136 MMCs is optimized; that is, the desired microstructural features and mechanical properties are achieved by tailoring the RGO content. 2. Experimental 2.1. Fabrication of RGO/S136 mixed powders The gas-atomized S136 steel powder (Changsha Hualiu Metallurgy Powder Co., Ltd., China) exhibited a spherical or nearly spherical shape with an average particle size of 15 μm, as shown in Fig. 1a and b. Its chemical composition is shown in Table 1. The quality and morphology of the RGO (Arke Inc.) are shown in Fig. 1c and d. The RGO was a single layer of graphene approximately 6–8 nm in thickness and 5 μm across. Sodium dodecyl benzene sulfonate (SDBS) with a purity of no less than 90% was obtained from Tianjin Kaitong Chemical Reagent Co., Ltd. Several steps were required for preparing the RGO/S136 mixed powders, and a schematic diagram of the process is shown in Fig. 2. In step 1, 1 g of RGO and 0.5 g of SDBS were placed into a beaker, which was then filled with 5000 ml of deionized water followed by mechanical stirring and ultrasonication for 40 min. Then, S136 powder was added to the suspension prepared above, and the container was placed onto a thermostatic heating platform at a constant temperature of 80 °C. During the entirety of step 2, ultrasonic stirring was maintained until the deionized water had completely evaporated. Finally, a RGO/S136 powder mixture was obtained by filtration, dried in a vacuum drying oven
Fig. 1. (a) SEM image of the S136 powder, (b) the distribution of the S136 particle size, (c) SEM image showing the morphology of RGO and (d) the RGO under high magnification.
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Table 1 Chemical composition of the S136 steel powder (wt%). Si
Mn
Cr
V
C
O
P
S
Fe
0.96
0.98
13.55
0.4
0.29
0.078
0.01
/
Bal.
(Model DZF-1, Shanghai, China) at 90 °C for 24 h and then sieved through a 325-mesh sieve for further use. 2.2. SLM process An SLM Solutions instrument comprising a 400 W IPG fibre laser and an inert argon gas protection system was used in this work. The 0.2 wt% RGO/S136 mixed powder was chosen, and 20 groups of SLM parameters were set to form the samples. A diagram of the relationship between the density and line energy density is shown in Fig. 3 and can be divided into region 1, region 2 and region 3 according to the different line energy densities. A typical optical micrograph of the corresponding sample was inset in each region. There were numerous voids on the surface of the sample in region 1 and an abundance of cracks on the surface of the sample in region 3. However, the sample in region 2 had fewer surface defects, which determined that the optimum combination of process parameters had been obtained in region 2. Therefore, in this study, the laser scanning power and speed were taken as 200 W and 900 mm/s, respectively, in an argon atmosphere with pre-heating to 200 °C. The laser scanning strategy and scanning direction of the samples are depicted in Fig. 4a. The cube-shaped samples, with dimensions of 10 mm × 10 mm × 10 mm and 5 mm × 5 mm × 5 mm, and tensile strength samples (mm) are shown in Fig. 4b. The samples were cut from the stainless steel substrate after fabrication by a wire electrical discharge machine. For comparison, pure S136 samples were also prepared by the same parameters. For simplicity, the 0, 0.1, 0.2 and 0.5 wt % RGO/S136 samples were labelled A1, A2, A3 and A4, respectively. 2.3. Characterization The densities of the samples were determined by the Archimedes method. Each set of samples was measured three times, and the results
Fig. 3. Relationship between density and the line energy density.
were averaged as the final density. The samples were ground with sandpaper from 240 mesh to 2000 mesh, followed by polishing with the diamond suspension (Lab Testing Technology Co.,Ltd. Shanghai) for metallographic testing. The defects such as cracks of the polished samples were observed under an optical microscopy (OM, Olympus BX60). The corrosive solution consisted of HCl (15 ml), H2O (60 ml) and FeCl3 (5 g). The microstructure of the samples and the fracture of the tensile specimens were observed under a FEI Sirion Quanta 200 scanning electron microscope (SEM). An X-ray diffraction pattern having a diffraction angle (2θ) ranging from 30° to 100° and a step size of 10°/min was obtained using an XRD-7000s (Shimadzu, Japan) X-ray diffractometer (CuKα radiation, a tube voltage of 40 KV and a tube current of 40 mA) in the RGO/S136 composites. The electron backscatter diffraction images were obtained by an HKL Nordlys orientation imaging microscope system (Oxford, Oxford Instruments, United Kingdom) with a step size of 1 μm. The acquired data is processed by HKL Channel 5 software to obtain images such as grain size and grain boundary on the surface of the samples. Five-point sampling method was used to select test points and the Vickers
Fig. 2. Schematic illustration of the procedure for preparing the mixed powder.
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Fig. 4. Schematic of (a) laser scanning strategy and (b) tensile strength sample.
3. Results and discussion 3.1. Phase composition and microstructure Typical XRD spectra of the SLM RGO/S136 composites obtained within a wide 2θ range (30–100°) are depicted in Fig. 5. The main phases in the composites were the α-Fe and γ-Fe phases. The composites with different RGO content have the same phase, and no obvious second phase is detected, which indicates the dominant martensitic transformation and austenitic transformation during the high temperature molten pool reaction of the SLM process. The α-Fe diffraction peak of the SLM samples slightly shifted to the left with increasing RGO content. According to Bragg's law [28]: 2d sinθ ¼ nλ ðn ¼ 1; 2; 3; …Þ Fig. 5. XRD analysis of the A1, A2, A3 and A4 samples.
hardness was measured at room temperature using a Vickers hardness tester (Wolpert Wilson Instruments 432 SVD, China), and the average value was taken as the final hardness. The fixed loading force was 3 kg and the loading time was 15 s. The Zwick/Roell Z020 (Germany) was used to completely pull off the sample at a load of 90 KN and a strain rate of 1 mm/min at room temperature. Each group of tensile samples was tested 5 times, and finally, the average yield strength and ultimate tensile strength were calculated in this study.
ð1Þ
the observed decrease of the 2θ value indicated an increase in the lattice plane distance d at higher RGO contents. This increase in lattice plane distance was primarily based on lattice distortion caused by martensitic transformation, as well as microscopic volume expansion and residual stress at grain boundaries [29,30]. Additionally, a portion of the C atoms in the alloy entered the Fe lattice as a solid solution and formed ferrite in the Fe matrix [31,32], leading to the change in the lattice parameters. However, there were no C peaks detected in and of the RGO/S136 composites due to the sensitivity of the XRD apparatus. Fig. 6a–d shows the fcc/bcc phase distribution obtained by the EBSD test. The results of the fcc/bcc phase ratios are displayed in Fig. 6e, which were calculated to be 31.8%, 1.3%, 42.9% and 58.5% for the A1, A2, A3 and
Fig. 6. (a)–(d) The phase distribution of the A1, A2, A3 and A4 samples and (e) the ratios of the fcc phase to the bcc phase corresponding to the A1, A2, A3 and A4 samples.
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A4 samples, respectively. These results were consistent with the XRD data shown in Fig. 5. Generally, the fcc phase is the austenite phase, and the bcc phase is the martensite phase in iron-based composites [4]. Martensite dominated the phase composition, reaching a maximum of 98.75% in the A2 sample, which contained 0.1 wt% RGO. Most of the fcc grains exhibited a high aspect ratio, and they were located preferentially along the boundaries between the bcc grains [33]. SEM images of the RGO/S136 samples with different RGO contents are presented in Fig. 7. The grains of the pure S136 sample mainly included equiaxed fine grains and cellular dendritic grains, as shown in Fig. 7a, which was consistent with the results reported by Zhou et al. [34]. The RGO/S136 samples were dominated by equiaxed fine grains, and a few columnar grains were found, as shown in Fig. 7b–d, which was different from that of the pure S136 alloy. The different microstructures indicated an obvious transition of the grains from cellular dendritic grains to equiaxed fine grains. The transformation to equiaxed fine grains in the matrix can be attributed to (i) the extremely fast cooling rate of the SLM process provides a shorter growth time; (ii) the RGO dispersed at the grain boundaries provides a large number of nanoscale nucleation sites; (iii) RGO pinned at the grain boundaries retards further grain growth. As Martin et al. [35] described, although the extremely high cooling rate during the SLM process favors equiaxed growth of the grain, heterogeneous nucleation induced by nano-sized RGO particles as nucleation sites is the key to promoting a columnarto-equiaxed transformation [36]. 3.2. Crystal orientation and crystallographic texture To study the effect of the RGO content on the grain size, orientation and texture of the RGO/S136 MMCs, EBSD measurements were carried out on the top surface of the samples (A1, A2, A3 and A4). The
5
representative orientation maps and inversed pole figure (IPF) of the RGO/S136 composites are shown in Fig. 8. All the IPF maps in this paper use the same colour coding, with the red, green, and blue colour levels being proportional to the three basic orientations of (001), (101), and (111), respectively. The orientation map of pure S136 (A1) shows that the grains mainly possessed a red-green-blue colour, as shown in Fig. 8a. Combined with the IPF, it could be determined that the pure S136 sample possessed a combination of the (001), (101) and (111) orientations. There were no significant differences between the areas corresponding to the three colours, which indicated that the pure S136 alloy had no obvious grain orientation. When 0.1 wt% RGO was added to the S136 alloy (A2), the grains of the composite exhibited a decrease in the (100) orientation and increased in the (101) and (111) orientations, as shown in Fig. 8b. When the RGO content was further increased to 0.2 wt% and 0.5 wt%, the orientation maps (Fig. 8c–d) were dominated by a green-blue colour, which indicated a strong (101) and (111) orientation in the corresponding composites. Based on the above observations, it could be concluded that the RGO content had a significant effect on the grain orientation of the RGO/S136 composites and promoted a transition of the grains towards the (101) and (111) orientations as the RGO content increased. Additionally, inhibition of the (100) grain orientation occurred when RGO was added. The orientation maps also provided meaningful information regarding the changes in the average grain size. The Channel 5 software was utilized to quantitatively analyse the average grain size [37]. The average grain size of the pure S136 alloy was calculated to be 0.86 μm, and the grain sizes were distributed over the range of 0.5 μm to 6.5 μm. When 0.1 wt% RGO was added to the alloy, the grains were refined and the average grain size was calculated to be 0.75 μm. However, some of the grains tended to become larger as the maximum grain size increased from 6.5 μm to 9.5 μm. Furthermore, the average grain
Fig. 7. SEM images of the A1, A2, A3 and A4 samples.
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Fig. 8. (a)–(d) EBSD orientation maps of the A1, A2, A3 and A4 samples and (e) the inverse pole figure (IPF).
sizes were 1.46 μm and 1.91 μm for the 0.2 wt% and 0.5 wt% RGO/S136 samples, respectively. The results showed that the grains gradually coarsened with increasing RGO content. Fig. 9 illustrates the grain boundary misorientation angle distribution obtained from the top surface of the A1, A2, A3 and A4 samples. Generally, grain boundaries for which the misorientation angle of the crystal is greater than 15° are defined as high-angle grain boundaries (HAGBs) and they are defined as low-angle grain boundaries (LAGBs) when the misorientation angle is less than 15°. The boundaries are colour-coded with LAGBs in green and HAGBs in black [38]. Fig. 10 provided a more intuitive regular distribution pattern of the grain boundary misorientation angles. The linear fractions of HAGBs in this region were calculated to be 69.86%, 72.07%, 73.92% and 34.98% for the A1, A2, A3 and A4 samples, respectively. HAGBs accounted for the majority of misorientation angles when the mass fraction of RGO in the composites was less than 0.2%. LAGBs dominated as the mass fraction of RGO increased to 0.5%.
Generally, the toughness of a material is closely related to the proportion of HAGBs in its structure. HAGBs can hinder crack propagation, i.e., for a higher HAGB fraction, the toughness of the material is better. In addition, the higher the misorientation angle of adjacent grain boundaries is, the stronger the ability of the material to resist crack propagation. Therefore, it can be predicted that the 0.5 wt% RGO/S136 sample may have the lowest toughness value among the four samples, by only considering the effect of grain boundary misorientation. The transformation mechanism between LAGBs and HAGBs is closely related to the uniqueness of the SLM process. The RGO/S136 MMC powders were added layer by layer to fabricate the MMCs by a high-energy laser beam. Along with the rapid melting and solidification of the SLM process, a large number of sub-grains nucleate and grow. The grain boundaries of these sub-grains tend to form LAGBs first. Additionally, the re-melting phenomenon in the SLM process is similar to the temper heat treatment in the conventional casting process, which results in recrystallization and an increase in the amount of HAGBs. Due
Fig. 9. Grain boundary misorientation angles of the A1, A2, A3 and A4 samples.
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Fig. 10. Grain boundary misorientation angle distributions of the A1, A2, A3 and A4 samples.
to the extremely short cooling time of the SLM process, the LAGBs are not fully converted into HAGBs and become residual LAGBs, consequently remaining in the sample [39,40]. As shown in Fig. 8, there were large differences in the grain orientations between the SLM samples with different mass fractions of RGO.
Therefore, the content of RGO had a significant influence on the local crystallographic texture of the SLM RGO/S136 composites. The most common (001) pole figures (PFs) from the top view of the SLM RGO/ S136 samples are shown in Fig. 11. Fig. 11a–d illustrates the (001) pole figures of the basic bcc phase in the MMC matrix, the
Fig. 11. (001) pole figures (PFs) of the A1, A2, A3 and A4samples. (a)–(d) The bcc phase PFs and (e)–(h) the fcc phase PFs.
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accompanying data corresponding to the fcc phase are shown in Fig. 11e–f. For the SLM of the pure S136 sample, a weaker (001) texture along the scanning direction was observed, as shown in Fig. 11a. The relatively low intensities in the PF suggested the presence of a fibrous texture in the tissue [41]. The texture index (TI) is usually induced to quantify the fibrous texture intensity. TI is obtained from the orientation function (ODF) f(g) by Eq. (2) [42]: Z Texture Index ¼ TI ¼
2
eulerspace
ð f ðgÞÞ dg
ð2Þ
where f is the orientation distribution as a function of the Euler space coordinates g. The texture indexes of the bcc phase corresponding to the S136 alloy and the 0.1 wt%, 0.2 wt% and 0.5 wt% RGO/S136 composites, were determined to be 7.17, 5.68, 9.79 and 12.4, respectively. The texture indexes of the fcc phase were determined to be 27.99, 20.02, 30.98 and 33.24, respectively. Therefore, the fcc phase showed a stronger (001) orientation than that of the bcc phase. Considering the results of the texture indexes of the fcc and the bcc phases comprehensively, it can be seen that they were in good agreement with the EBSD orientation maps shown in Fig. 8. Fig. 12 shows the proportions of the Schmidt factors of different values, corresponding to the A1, A2, A3 and A4 samples. The inset in each figure indicates the different regions of the sample surface corresponding to the different Schmidt factor values. The Schmidt factors of the A1, A2, A3 and A4 samples were calculated to be 0.4442, 0.4573, 0.4039 and 0.4484, respectively. It is well known that the plastic deformation ability of a material can be characterized by its Schmidt factor [43]. Hence, the Schmidt factor is often utilized for evaluating the yield strength (σs) of a material, which can be calculated as follows: σ s ¼ τs
1 cosλ cosφ
ð3Þ
where λ is the angle between the applied load F and the 〈110〉 orientation, φ is the angle between the applied load F and normal direction of {111}, and τs is the critical shear stress. The product of cos λ and cos φ is called the Schmidt factor [44]. Based on Eq. (3), a lower Schmidt factor usually leads to a higher σs value. The Schmidt factor of A3 was the smallest of all samples, as seen in Fig. 12, indicating that A3 may have the highest yield strength. 3.3. Mechanical properties of the SLM RGO/S136 composites The Vickers micro-hardness values of the A1, A2, A3 and A4 samples are shown in Fig. 13. The Vickers micro-hardness of the pure S136 sample was approximately 544.1 HV, which was similar to the hardness value reported by Wen et al. [4]. The maximum Vickers microhardness value of the RGO/S136 composites was 580.6 HV, reaching the peak value after the addition of 0.1 wt% RGO, which was 6.3% higher than that of the pure S136 sample. Fig. 13 also indicates that the Vickers micro-hardness decreased with the further addition of RGO, for which the Vickers micro-hardness values of A3 and A4 were measured to be 534.6 HV and 468.7 HV, respectively. Nevertheless, the SLM RGO/S136 MMCs, for all the given RGO contents, demonstrated superior Vickers micro-hardness values compared to that of commercially used S136, which has a typical hardness of 380–400 HV. In general, residual stress generated during the SLM process is unavoidable, and excessive residual stress can cause local failure of components. However, it should be pointed out that in a high-density metal matrix composite with few cracks and voids (Fig. 3), a reasonable level of residual stress is conducive to the improvement of hardness. Moreover, the significant grain refinement effect (Fig. 8) due to the inclusion of nanoscale RGO favoured a further increase in the obtained microhardness. Fig. 14a shows the stress-strain curves of the RGO/S136 composites, and the trends in the yield strength and ultimate tensile strength for different RGO contents is described in Fig. 14b. Both the yield strength and ultimate tensile strength showed an initially increasing trend and then
Fig. 12. Distribution histograms of the Schmidt factors of the A1, A2, A3 and A4 samples.
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Table 2 Mechanical properties of samples A1, A2, A3 and A4. Sample number A1 A2 A3 A4
Fig. 13. Vickers hardness values of the surfaces of the A1, A2, A3 and A4 samples.
decreased with increasing RGO content. The maximum yield strength and ultimate tensile strength appeared in 0.1 wt% RGO/S136, increasing by 4.1% and 30.4%, respectively, compared with those of the pure S136 sample. Tensile properties of the composites decreased as the RGO content increased to 0.2 wt% and 0.5 wt%. Importantly, the elongation remained essentially unchanged. The σs of A1, A2, A3 and A4 are shown in Table 2 and were measured to be 359.2 ± 12.7 MPa, 515.8 ± 15.3 MPa, 259.2 ± 19.1 MPa and 275.9 ± 15.8 MPa, respectively. The results indicated that the effects of the Schmidt factor on the yield strength were no longer the main factor for the SLM RGO/S136 samples. The strengthening mechanism of RGO in the S136 matrix can be explained by the following aspects: stress transfer, dislocation strengthening and grain refinement [45–47]. The stress transfer mainly depended on the interfacial bonding between the S136 matrix and RGO [48]. The large mismatch in the thermal expansion coefficients of S136 and RGO formed dislocations [49]. The hard RGO phase impeded the movement of these dislocations, leading to an increase in the dislocation density and resulting in dislocation strengthening [50]. The variation in average grain size is reflected by the EBSD results in Fig. 8, and the average grain size decreased from 0.86 μm to 0.75 μm in the pure S136 and 0.1 wt% RGO/S136 samples, respectively. The pinning effect of RGO was remarkably beneficial to the dispersion of the grain boundaries, which prevented grain growth and reduced the crack propagation properties to some extent. Therefore, RGO as a reinforcement could improve the yield strength and ultimate tensile strength of the S136 matrix. However, the tensile properties decreased significantly as the RGO content
Yield strength σs (MPa) 359.2 ± 12.7 515.8 ± 15.3 259.2 ± 19.1 275.9 ± 15.8
Tensile strength σb (MPa)
Elongation (%)
513.6 ± 20.3 535.3 ± 18.5 339.4 ± 17.2 311.5 ± 21.4
4.4 ± 0.3 4.3 ± 0.8 4.4 ± 1.2 3.7 ± 0.5
increased further. There were two main reasons accounting for these results. On the one hand, agglomeration inevitably occurred with increasing RGO content, which resulted in the degradation of the pinning effect and resulted in the grain sizes becoming coarser [51]. On the other hand, RGO may partially react with the S136 matrix under the high temperature of the laser beam and lose its unique strengthening properties. The fractured surfaces of the RGO/S136 composites are displayed in Fig. 15. The fractured surface of pure S136 (Fig. 15a) showed a brittle fracture, but a small number of dimples could be found, as marked by the yellow arrow in Fig. 15b. For the 0.1 wt% RGO/S136 composite, a very small amount of RGO can be observed in Fig. 15c, and numerous dimples were observed at the interface between the RGO and S136 matrix in Fig. 15e. In Fig. 15d, it can also be seen that RGO possessed a significant number of fractures and holes. The fractured surface of the 0.2 wt% RGO/S136 composite is shown in Fig. 15f–g. The areas marked by the arrows in Fig. 15f had typical curving and wrinkling structures from the removal of RGO. The unique honeycombed structure of RGO was observed in Fig. 15g, further proving the presence of RGO. At the same time, several regular hexagonal components underwent significant deformation and elongation under stress. There were cracks at the junction of the RGO and S136 matrix, as shown in Fig. 15h, indicating that the interfaces between the RGO and S136 matrix had not been well combined. Different shrinkages along the interfaces during the cooling process may be the reason for this phenomenon [52]. 4. Conclusions RGO/S136 composites were successfully fabricated by a liquid deposition method and SLM process. The effects of the RGO content (0, 0.1 wt %, 0.2 wt% and 0.5 wt%) on the microstructures, phase compositions and mechanical properties of the RGO/S136 composites were investigated. The results obtained can be summarized as follows: 1. The RGO content showed a significant effect on the grain orientation of RGO/S136 composites, promoting a transition of the grain orientation from a combination of the (001), (101), and (111) orientations to the (101) and (111) orientations as the RGO content increased.
Fig. 14. (a) Stress-strain curve of the RGO/S136 composites and (b) relationships of the ultimate tensile strength and yield strength with different RGO contents.
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Fig. 15. Fractured surfaces of the RGO/S136 composites. (a)–(b) Pure S136, (c)–(e) 0.1 wt% RGO/S136, (f)–(g) 0.2 wt% RGO/S136 and (h) 0.5 wt% RGO/S136.
2. The RGO content was confirmed to have a refinement effect on the average grain size, as the grain size was 0.75 μm with 0.1 wt% RGO, which was smaller than the grain size of 0.85 μm in the pure S136 alloy. 3. The phases in the RGO/S136 composites mainly consisted of the α-Fe and γ-Fe phases. The addition of RGO induced a transition of the grains from cellular dendritic grains to equiaxed fine grains.
4. With increasing RGO content, the hardness, ultimate tensile strength and yield strength of the composites initially increased and then decreased. The microhardness, tensile strength and yield strength reached the maximum of 580.6 ± 15.4 HV, 535.3 ± 18.5 MPa and 515.8 ± 15.3 MPa, respectively, when the RGO content was 0.1 wt %. The strengthening mechanism in the RGO/S136 composites was mainly attributed to stress transfer, dislocation strengthening and grain refinement.
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CRediT authorship contribution statement Shifeng Wen: Writing - original draft. Keyu Chen: Data curation, Investigation. Wei Li: Methodology. Yan Zhou: Funding acquisition, Formal analysis. Qingsong Wei: Supervision. Yusheng Shi: Writing - review & editing. Acknowledgements This study was supported by National Hi-tech R&D Program of China (863 Program, Grant No. 2015AA042501), the National Natural Science Foundation of China (Grant No. 51605176), Hubei Provincial Natural Science Foundation of China (Grant No. 2018CFB502), Guangdong Provincial Technology Major Project of China (Grant No. 2017B090911007), State Key Laboratory of Materials Processing and Die and Mould Technology, Huazhong University of Science and Technology (Grant No. P2019-006), and Engineering Research Center of Rock-Soil Drilling & Excavation and Protection, Ministry of Education (Grant No. 201804). References [1] Y. Wang, J. Shi, S. Lu, Y. Wang, Selective laser melting of graphene-reinforced Inconel 718 superalloy: evaluation of microstructure and tensile performance, J. Manuf. Sci. Eng. 139 (2016) (041005–041005-6). [2] T. Kimura, T. Nakamoto, Microstructures and mechanical properties of A356 (AlSi7Mg0.3) aluminum alloy fabricated by selective laser melting, Mater. Des. 89 (2016) 1294–1301. [3] P. Han, Additive design and manufacturing of jet engine parts, Engineering 3 (2017) 648–652. [4] S. Wen, H. Hu, Y. Zhou, Z. Chen, Q. Wei, Y. Shi, Enhanced hardness and wear property of S136 mould steel with nano-TiB2 composites fabricated by selective laser melting method, Appl. Surf. Sci. 457 (2018) 11–20. [5] S. Wen, X. Ji, Y. Zhou, C. Han, Q. Wei, Y. Shi, Corrosion behavior of the S136 mold steel fabricated by selective laser melting, Chinese J. Mech. Eng. 31 (2018) 108. [6] B. He, Reseach on the failure and material selection of plastic mold, Procedia Eng 23 (2011) 46–52. [7] X. Li, L. Duan, J. Li, X. Wu, Experimental study and numerical simulation of dynamic recrystallization behavior of a micro-alloyed plastic mold steel, Mater. Des. 66 (2015) 309–320. [8] W. Li, H. Yuan, Z. Sun, Z. Zhang, Surface vs. interior failure behaviors in a structural steel under gigacycle fatigue: failure analysis and life prediction, Int. J. Fatigue 64 (2014) 42–53. [9] Z. Chen, L. Dong, D. Yang, H. Lu, Superhydrophobic graphene-based materials: surface construction and functional applications, Adv. Mater. 25 (2013) 5352–5359. [10] H. Bi, I. Chen, T. Lin, F. Huang, A new tubular graphene form of a tetrahedrally connected cellular structure, Adv. Mater. 27 (2015) 5943–5949. [11] A.K. Geim, K.S. Novoselov, The rise of graphene, Nat. Mater. 6 (2007) 183–191. [12] C. Lee, X. Wei, J.W. Kysar, H. James, Measurement of the elastic properties and intrinsic strength of monolayer graphene, Science 321 (2008) 385–388. [13] S.C. Tjong, Recent progress in the development and properties of novel metal matrix nanocomposites reinforced with carbon nanotubes and graphene nanosheets, Mater. Sci. Eng. R Reports. 74 (2013) 281–350. [14] A. Nieto, A. Bisht, D. Lahiri, C. Zhang, A. Agarwal, Graphene reinforced metal and ceramic matrix composites: a review, Metall. Rev. 62 (2016) 241–302. [15] S. Feng, G. Qiang, L. Zan, G. Fan, Z. Li, D. Xiong, Y. Su, Z. Tan, Z. Jie, Z. Di, Strengthening and toughening mechanisms in graphene-Al nanolaminated composite micropillars, Acta Mater. 125 (2017) 98–108. [16] Y. Ming, W. Lin, H. Zhu, T. Fan, Z. Di, Simultaneously enhancing the strength, ductility and conductivity of copper matrix composites with graphene nanoribbons, Carbon N. Y. 118 (2017) 250–260. [17] W.J. Kim, T.J. Lee, S.H. Han, Multi-layer graphene/copper composites: preparation using high-ratio differential speed rolling, microstructure and mechanical properties, Carbon N. Y. 69 (2014) 55–65. [18] L. Chen, J. Xu, H. Choi, M. Pozuelo, X. Ma, S. Bhowmick, J. Yang, S. Mathaudhu, X. Li, Processing and properties of magnesium containing a dense uniform dispersion of nanoparticles, Nature 528 (2015) 539. [19] K. Nie, X. Wang, K. Wu, L. Xu, M. Zheng, X. Hu, Processing, microstructure and mechanical properties of magnesium matrix nanocomposites fabricated by semisolid stirring assisted ultrasonic vibration, J. Alloy. Compd. 509 (2011) 8664–8669. [20] Y. Huang, P. Bazarnik, D. Wan, D. Luo, P.H.R. Pereira, M. Lewandowska, J. Yao, B.E. Hayden, T.G. Langdon, The fabrication of graphene-reinforced Al-based nanocomposites using high-pressure torsion, Acta Mater. 164 (2019) 499–511. [21] L. Wang, C. Tian, S. Wang, Microstructural characteristics and mechanical properties of carbon nanotube reinforced AlSi10Mg composites fabricated by selective laser melting, Opt. - Int. J. Light Electron Opt. 143 (2017) 173–179. [22] Z. Li, Q. Guo, Z. Li, G. Fan, D. Xiong, Y. Su, J. Zhang, D. Zhang, Enhanced mechanical properties of graphene (reduced graphene oxide)/aluminum composites with a bioinspired nanolaminated structure, Nano Lett. 15 (2015) 8077–8083.
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