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Al-Si composites
Journal Pre-proof In-situ SiC reinforced Si-SiC 3D skeletons in SiC/Al-Si composites Chongchong Wu, Tong Gao, Xiangfa Liu PII:
S0925-8388(19)32963-9
DOI:
https://doi.org/10.1016/j.jallcom.2019.151730
Reference:
JALCOM 151730
To appear in:
Journal of Alloys and Compounds
Received Date: 5 June 2019 Revised Date:
29 July 2019
Accepted Date: 5 August 2019
Please cite this article as: C. Wu, T. Gao, X. Liu, In-situ SiC reinforced Si-SiC 3D skeletons in SiC/Al-Si composites, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/ j.jallcom.2019.151730. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
In-situ SiC reinforced Si-SiC 3D skeletons in SiC/Al-Si composites
Chongchong Wu1, Tong Gao1, Xiangfa Liu1* 1 Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials, Ministry of Education, Shandong University, 17923 Jingshi Road, Jinan 250061, PR China
Abstract: For Al-Si alloys, the three-dimensional (3D) network structure of Si plays a key role on their mechanical properties. In order to improve the microstructure and elevate the strength of the network, insitu SiC was introduced into Al-Si alloys and four kinds of composites and their Si-SiC 3D skeletons were obtained. The microstructures especially the interfacial characteristics were studied in detail. Due to the heterogeneous nucleation role of in-situ SiC, effective interface was formed between in-situ SiC and Si. Moreover, in-situ SiC can grow epitaxially at the surface of external SiC and form a stable Si/SiCin/SiCex sandwich structure with in-situ SiC as the interlayer. As a result, the Si-SiC 3D skeletons became more intact and stable, and their compressive strength was remarkably elevated.
Keywords: Three-dimensional network; Skeleton; Al-Si alloy; In-situ SiC; Interfacial bonding; Strengthening mechanism
1. Introduction Al-Si alloys are the most common used aluminum casting alloys due to their good fluidity and properties. Al-Si alloys can be treated as composites in which ductile α-Al is the matrix and stiff Si is the
*Corresponding author. E-mail: [email protected] (X. Liu).
reinforcement [1-3]. Thus, the morphology and structure of the reinforcement, Si, plays a key role on the mechanical properties of Al-Si alloys. In general, Si phases can form a three-dimensional (3D) network structure during solidification, and it has been proved that more intact of the network, much higher of Al-Si alloys strength [4-6]. The morphologies of the Si 3D network can be changed by many different methods. Heat treatment is the easiest way to change the 3D network, but this method is always detrimental to the network connectivity when used alone [7-9]. With solution tr eatment, eutectic Si with interconnected lamellar structure in as-cast condition will coarsen and spheroidize, resulting in the loss of network connectivity [7]. Modifiers like Sr, Na and TiB2 can also affect the network morphology by affecting the growth of Si during solidification [9-16]. For instance, the addition of about 0.02% wt% Sr can induce the modification of the Si network from lamellar to coral morphology, and thus elevate the ductility and thermal fatigue of the alloys [9, 10]. TiB2 particles can also refine eutectic Si by obstructing solute redistribution and thus improve the Si network [15]. Transition elements like Cu, Fe and Ni can improve the connectivity of Si network by forming intermetallics and supplementing the Si network [17-21]. Asghar [17] found that the addition of Ni can improve the high temperature strength of Al-Si alloys by participating in the formation of a highly interconnected Si network. In reference 20, it is proved that Cu, Ni intermetallics can not only develop the 3D network but also strengthen the Al matrix. In order to achieve desired networks, above mentioned methods can be employed together [9, 11, 14]. For instance, Sr modified Al-12Si alloy is more sensitive to solution treatment than unmodified alloy and thus spheroidize and coarsen more quickly [9]. With the combined addition of modifier Sr and element Cu to Al-7Si, networks consist of fine globular-fibrous eutectic Si and CuAl2 can be obtained, and the mechanical properties of the alloy are much better than the single-addition ones [14]. Al-Si matrix composites are also widely used nowadays and they have much higher strength than Al-Si alloys. For these composites, the Si 3D network also plays an important role on their properties [22-27]. With the addition of Al2O3 short fiber preforms into Al-Si alloys, an interconnected Al2O3-Si network was formed and the elastic modulus and the strength of the alloys was elevated [22]. SiC
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particles can also enhance the formation of a continuous Si-SiC network, resulting in the considerable improvement of the modulus [23]. Besides, the hybrid addition of Al2O3 short fiber and SiC particles has also been studied [24, 25]. Compared with the heat treatment, modifiers and transition elements methods, less attention has been paid on the 3D network of Al-Si matrix composites. In order to study the 3D morphology of the Si network, tomographic imaging techniques are the most useful methods [9, 12, 17, 19, 26]. For instance, focused ion beam/scanning electron microscopy (FIB/SEM) tomography was used to reveal the 3D structure of eutectic Si in reference 9, while energy dispersive synchrotron X-ray diffraction was employed in reference 26. Methods like in situ neutron diffraction can also be used to analysis the Si 3D network [5]. However, these techniques are costly and time-consuming. In our previous works, in-situ SiC was successfully synthesized in Al-Si alloys and it can obviously affect the mechanical properties of Al-Si alloys [28-31]. In this research, the effect of in-situ SiC on the Si 3D networks was studied. At first, four kinds of Al-Si-SiC composites, Al-33Si (primary Si particles reinforced Al-Si composites), Al-13Si-20SiCex (externally added SiC reinforced Al-Si composites), Al13Si-20SiCin (in-situ SiC reinforced Al-Si composites) and Al-13Si-10SiCex-SiCin (externally added SiC and in-situ SiC reinforced Al-Si composites) were obtained. In order to observe the 3D networks, the AlSi-SiC composites were fully etched by hydrochloric acid, and the networks were completely extracted as skeletons. Thus, the skeletons can be studied in detail using routine methods. Results show the impressive effect of in-situ SiC on the formation of intact Si-SiC skeletons.
2. Experiment Commercial Al (99.7 wt% purity), Si (99.9 wt% purity), graphite (99.2 wt% purity, 1.3 µm, Huatai Graphite, China) and green SiC powders (98.5 wt% purity, 75 µm, Sinopharm Chemical Reagent, China) were used in this research. The morphologies of the raw materials are shown in Fig. 1. At first, four kinds of raw powders, with compositions (wt%) of Al-33Si, Al-13Si-20SiCex, Al-27Si-6C and Al20Si-3C-10SiCex were prepared. The raw powders were then mixed in blender mixer for 8 h. After that,
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they were pressed into cylindrical mould at 270 MPa with cold isostatic press (Taiyuan East Dragon Machinery Co. Ltd, LDJ200/500-380YS, China), with size of Φ60×50 mm. Finally, the compacted samples were sintered in vacuum resistance furnace at 800˚C for 2 h with vacuum degree higher than 103
Pa and cooled to room temperature in the furnace, under which condition the samples can react
completely [28]. The overall reaction during synthesis is shown in equation 1 [28]. According to the compositions of the raw powders, the as-sintered composites were Al-33Si, Al-13Si-20SiCex, Al-13Si20SiCin and Al-13Si-10SiCex-10SiCin, respectively [28]. In these composites, to guarantee the same proportion of the skeletons in the composites, the weight percent of the strengthening particles (Si and SiC) is equal. Si(l) + C(s) = SiC(s)
(1)
To observe the skeletons in these composites, samples were cut from the center of the as-sintered materials with size of 10×10×10 mm, and grinded in standard routines. The samples were then etched by 4 vol% hydrochloric acid sufficiently (about 96 h) until the pH value of the etched solution remain constant. Thus, four kinds of Si-SiC skeletons were prepared. The macrostructures of the composites and their skeletons are shown in Fig. 2. To check the heterogeneous nucleation role of in-situ SiC, Al-13Si-20SiCin was introduced into an Al-12Si-4Cu-2Ni alloy at 800˚C in resistance furnace. Al, Si, Cu and Ni were melted at first, and Al13Si-20SiCin alloy was added after. Refined by high purity argon, the melt was finally cast into iron mould at 800˚C. Metallographic specimens were cut from the center of the cast, and prepared in standard routines. To analyze the phase composition, the skeletons were grinded into powders and tested with X-ray diffraction (XRD, Rigaku D/max-rB, Cu Kα radiation). The microstructures of the skeletons were observed with field emission scanning electron microscopy (FESEM, Hitachi SU8010) equipped with an energy dispersive X-ray spectroscopy (EDS) detector. To conveniently observe the microstructures of the skeletons under scanning electron microscopy (SEM), samples with size of Φ15×10 mm were cut from the composites and were deeply etched (about 10 min) by 4 vol% hydrochloric acid. The fine
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microstructures of the interfaces were observed with atomic resolution scanning transmission electron microscope (S/TEM, FEI Titan G2 60-300). The S/TEM samples were prepared using focused ion beam (FIB) in regular process. The compressive strength of the skeletons was obtained according to GB/T8489-2008 by universal testing machine (Shimadzu, AG-IC100KN), and at least three samples were tested for each skeleton. To test the pore size distribution of the skeletons, mercury intrusion method (Micromeritics, AutoPore Iv9500), with the stem volume of 1.131 mL, was used. During mercury intrusion, non-wetting mercury was pressed into porous medium, and the pressure is inversely proportional to pore size. The pressure was converted to the equivalent pore size using Washburn equation [32, 33], with the surface tension of 0.485 N/m and the contact angle of 130°. Compared with other methods like gas absorption, mercury intrusion method can test pores with a large range of size, from micropores to macropores [34, 35], and therefore this method is more appropriate to the skeletons in this work which have near-sphere pores range from several nm to several µm.
3. Results and discussion Fig. 2 is the macrostructures of Al-Si-SiC composites and their skeletons. As shown by Fig. 2a to 2d, all the composites, with or without SiC, are compact. Among them, as shown in Fig. 2b, only composites with external SiC show inhomogeneous structure: a strip bright area can be observed in the front face. By contrast, as shown in Fig. 2e to 2h, the skeletons of these composites show much different behaviors. As shown in Fig. 2e, Si skeletons, the skeletons of hypereutectic Al-Si alloys, can almost remain as a whole, but some detached particles can also be observed. Fig. 3 shows the microstructures of Si skeletons. From Fig. 3a and 3b, it is obvious that the skeletons are composed of primary Si with irregular blocky shape and eutectic Si with irregular rod-like shape. From Fig. 3c, the combination between primary Si particles is realized by eutectic Si. However, from Fig. 3a and 3b, this combination is not always valid. Therefore, some Si particles cannot firmly connect to the skeletons, just as the detaching
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particles shown in Fig. 2e. Besides, from Fig. 3a, the shapes of the pores in the skeletons are not regular and the size of them vary widely, which reflect the nonuniform microstructure of the skeletons. In contrast, as shown in Fig. 2f, when external SiC was added into Al-Si alloys, the original skeletons are destroyed and fell apart. Fig. 4 shows the microstructures of Si-SiCex skeletons, the skeletons of Al-13Si-20SiCex composites. As shown, external SiC distribute among Si phases, but no effective connection exists between them. Because of the mechanical crushing process during the production of externally added green SiC, the surfaces of them are irregular and have no certain crystal orientation, as shown in Fig. 1d [31, 36]. This make it hard for Si to nucleate and grow at the surface of external SiC during the solidification of Al-Si alloy. Therefore, no effective connection exists between external SiC and Si. In other words, the introduction of external SiC breaks the original interfaces between Si phases. Unlike external SiC, when in-situ SiC was introduced into Al-Si alloys, the skeletons in the alloys became more intact without detached particles, which can be seen in Fig. 2g. Moreover, when in-situ SiC was introduced into Al-Si-SiCex composites, the destroyed skeletons in Fig. 2f changed into unbroken ones again, as shown in Fig. 2h. This means that the introduction of in-situ SiC can enhance the Si/Si and SiCex/Si interfacial bonding. To study the enhancing mechanism of in-situ SiC on Si/Si and Si/SiCex interfaces, further research was carried out. Fig. 5 are the XRD patterns and the microstructures of the skeletons in Al-13Si-20SiCin (in-situ SiC reinforced Al-Si composites) and Al-13Si-10SiCex-10SiCin (in-situ SiC reinforced Al-Si-SiCex composites). From Fig. 5a and 5b, only diffraction peaks of Si and SiC can be observed while no peaks of Al were detected. Therefore, the skeletons are composed of Si and SiC. From Fig. 5c and 5d, intercommunicating pores with micrometer size distribute homogenously in both skeletons. As shown in Fig. 5d, external SiC can be observed clearly in the skeletons of Al-13Si-10SiCex-10SiCin. According to the compositions of the composites, the skeletons of Al-13Si-20SiCin and Al-13Si-10SiCex-10SiCin are Si-60SiCin and Si-30SiCex-30SiCin, respectively.
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Fig. 6 shows the microstructures of Si-SiCin skeletons (the skeletons of Al-13Si-20SiCin composites) in Fig. 2g. From the overall structure in Fig. 6a and the pore size distribution in Fig. 6b, bimodal pore distribution was shown: two kinds of pores, macropores with size between 5 µm and 10 µm and mesopores with size between 30 nm and 100 nm, can be observed. As shown in Fig. 6a and 6c, macropores have circular shape in two-dimension and distribute uniformly among the skeletons. By contrast, as shown in Fig. 6d and 6e, mesopores have no certain shape and scatter randomly between the particles. Compared with the Si skeletons in Fig. 3, the shape of the pores becomes more regular and the distribution of them becomes more uniform. Fig. 6f and 6g is the EDS analysis of Si and in-situ SiC in the skeletons. As shown, the light grey phases with regular morphology, hexagonal plates mostly, are in-situ SiC, while the dark grey phases with irregular morphology are Si. No eutectic Si as the one in Fig. 3d can be observed. From Fig. 6d and 6e, in-situ SiC particles with different directions interweave with each other and distribute irregularly among Si phases. Moreover, in Fig. 6e, Si phases attach to in-situ SiC particles with certain angles. The relationship between in-situ SiC and Si phases is determined by their precipitation process. During the sintering of Al-Si-SiCin composites, in-situ SiC with preferred growth orientations began to formed when temperature is higher than 750˚C [28]. Growing SiC particles with different directions may contact with each other, but still can remain their growth orientations, resulting in the interweaved SiC structures in Fig. 6e. During solidification, in-situ SiC can serve as the heterogeneous nucleation sites for Si. Thus, eutectic Si cannot grow freely in Al-Si eutectic structure but grow around in-situ SiC particles as shown in Fig. 6d. Therefore, no dendritic Si can be observed. In addition, owing to twins and multiple twins, the growth orientations of Si can change with certain degrees [37-39]. Thus, Si was nucleated and grew among interweaved in-situ SiC with certain orientation relationships as shown in Fig. 6d and 6e. To check the heterogeneous nucleation role of in-situ SiC, Al-Si-SiCin was added to an Al-12Si4Cu-2Ni alloy with proportion of 1%. The microstructures of the obtained materials are shown in Fig. 7. In Fig. 7a, micron particles can be found in the center of Si phases. Fig. 7b shows one of these particles. EDS analysis of this particle in Fig. 7b and 7c prove that this particle is SiC. Therefore, in-situ SiC can
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act as the heterogeneous nucleus of Si during Al-Si solidification process. According to the Si-SiC interweaved structure in Fig. 6d and the analysis in Fig. 7, Si should heterogeneously nucleate on specific surface of SiC which has certain orientations or surface energy. From Fig. 6 and reference 29, in-situ synthesized SiC has certain exposed faces, while from Fig. 1d, the exposed faces of external SiC are random. Therefore, in-situ SiC is more appropriate for the heterogeneous nucleation of Si. Compared with the Si skeletons in Fig. 3, two main changes can be observed with in-situ SiC reinforced Si skeletons. For one thing, more effective connection is realized between Si particles with insitu SiC as the attachment. For another, due to the heterogeneous nucleation effect of in-situ SiC, dendritic Si disappears and the pores in the skeletons become more uniform and regular. Similarly, when in-situ SiC was introduced into Si-SiCex skeletons, obvious changes can be observed. Fig. 8 shows the microstructures of Si-SiCex-SiCin skeletons (the skeletons of Al-13Si-10SiCex10SiCin composites). As shown in Fig. 8a, external SiC particles are inlaid in the skeletons with effective connection. From Fig. 8c and 8d, a Si/SiCin/SiCex sandwich structure is formed with in-situ SiC as the interlayer. In addition, as shown in Fig. 8b, due to the introduction of external SiC, pores with trimodal distribution are observed—three sizes of pores with diameter around 150nm, 1µm and 7 µm are present— which is different from the bimodal distribution in Fig. 6b. The interface between in-situ SiC and external SiC is further studied in Fig. 9. From Fig. 9a and 9b, the interface is close and intact. As shown by the HRTEM results in Fig. 9c to 9e, in-situ SiC and external SiC have same crystal structures and parallel directions. This indicates that in-situ SiC interlayer was formed by the epitaxial growth of external SiC. The epitaxial growth of in-situ SiC can also refer to reference 31. From above analysis, the formation of in-situ SiC reinforced Si-SiCin skeletons is illustrated in Fig. 10. During sintering of Al-Si-SiCex-SiCin composites, in-situ SiC interlayers began to grow epitaxially at the surface of external SiC at about 750˚C, while at the same time, in-situ SiC particles were also formed in the melt. During solidification, in-situ SiC interlayers and particles can both act as the heterogeneous
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nucleus for Si. As a result, effective connection was obtained between external SiC and Si with in-situ SiC as the interlayer, just as the Si/SiCin/SiCex sandwich structure shown in Fig. 8d. Moreover, because in-situ SiC interlayers have same crystal structure with in-situ SiC particles, they can also participate in the formation of the interweaved structures in Fig. 2e. Therefore, as shown in Fig. 10, firm and consistent skeletons were formed. Compared with the Si-SiCex skeletons in Fig. 4, by introducing in-situ SiC, obvious differences were brought, especially the behavior of external SiC particles. In Fig. 4, external SiC particles were isolated from the skeletons, and barely no effective Si/SiCex bonding exists. By contrast, in in-situ SiC reinforced skeletons, owing to the formation of in-situ SiC interlayer, a Si/SiCin/SiCex sandwich structure was obtained, and thus external SiC particles were embedded in the skeletons closely. Fig. 11 shows the microstructures of four kinds of skeletons discussed above. Compared with Si and Si-SiCex skeletons in Fig. 11a and 11b, it is apparently that more effective and intact interfacial bonding is obtained with in-situ SiC reinforced Si skeletons (Fig. 11c) and in-situ SiC reinforced SiSiCex skeletons (Fig. 11d). Besides, the microstructure of the skeletons became more consistent and uniform. From above analysis, in-situ SiC is key for the formation of effective interfacial bonding in Si-SiC skeletons. On the one hand, it can serve as heterogeneous nucleation centers for Si. On the other, it can grow epitaxially at external SiC and form interlayers between external SiC and Si. Due to the formation of enhanced interfacial bonding, the strength of the skeletons has been elevated obviously. As shown in Fig. 12, after introducing in-situ SiC, the compressive strength of the skeletons has been elevated, from 0.39 MPa to 2.12/1.89 MPa, with an elevation of about 5 times. Above results show that introducing in-situ SiC into Si skeletons can guarantee the firm interfacial bonding between SiC and Si, and impressively improve the strength of Si-SiC skeletons. The introduction of in-situ SiC offers a new method to strengthen the 3D network structure in Al-Si alloys. More research, such as the optimization of the skeleton structures by controlling the content and size of in-situ SiC, still need to be carried out in future.
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4 Conclusions In-situ SiC reinforced Si-SiC skeletons in Al-Si-SiC composites were studied in this research. Their microstructures, interfacial properties and strengthening mechanisms were discussed in detail. The following conclusions can be drawn: (1) With the introduction of in-situ SiC, the skeletons in Al-Si-SiC composites became more intact and stable, the interfacial bonding in the skeletons became more effective, and the microstructure of the skeletons became more consistent and uniform. (2) In-situ SiC can serve as the heterogeneous nucleation centers for Si during the solidification of Al-Si alloys, resulting in a firm Si/SiCex interface. Moreover, in-situ SiC can grow epitaxially at the surface of external SiC and act as the interlayers between external SiC and Si, forming a stable Si/SiCin/SiCex sandwich structure. (3) With the introduction of in-situ SiC, the compressive strength of the Si skeletons increased obviously. The compressive strength of Si skeletons has been increased from 0.39 MPa to 2.12MPa, and the broken Si-SiCex skeletons have been changed to intact ones with compressive strength of 1.89 MPa.
Acknowledgement This work was supported by the National Natural Science Foundation of China (No. 51731007) and the key foundation of Shandong Province (No. ZR2016QZ005).
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incorporation of micron-sized SiC particles into molten pure aluminum utilizing a Co coating, Metall. Mater. Trans B 46 (2015) 12-19. [37] K. F. Kobayashi, L. M. Hogan, The crystal growth of silicon in Al-Si alloys, J. Mater. Sci. 20 (1985) 1961-1975. [38] K. Fujiwara, K. Maeda, N. Usami, K. Nakajima, Growth mechanism of Si-faceted dendrites, Phys. Rev. let. 101 (2008) 055503. [39] H. K. Lin, C. W. Lan, Revisiting the twinning mechanism in directional solidification of multicrystalline silicon sheet, Acta Mater. 131 (2017) 1-10.
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Figure captions: •
Fig. 1. Morphologies of raw materials: (a) Al, (b) Si, (c) graphite and (d) green SiC powders.
•
Fig. 2. Macrostructures of (a-d) Al-Si-SiC composites and (e-h) their corresponding skeletons: (a) Al-33Si and (e) the Si skeleton; (b) Al-13Si-20SiCex and (f) the Si-SiCex skeleton; (c) Al-13Si20SiCin and (g) the Si-SiCin skeleton; (d) Al-13Si-10SiCex-10SiCin and (h) the Si-SiCex-SiCin skeleton.
•
Fig. 3. Microstructures of Si skeletons in Al-33Si alloys: (a) (b) overall microstructure, (c) interface betweeen priamry Si particles and (d) eutectic Si. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (d).
•
Fig. 4. Microstructures of Si-SiCex skeletons in Al-13Si-20SiCex composites: (a) (b) overall microstructure and (c) interface between external SiC and Si. As shown by the white lines, (b) is the enlargement of the rectangle area in (a) and (c) is the enlargement of the rectangle area in (b).
•
Fig. 5. XRD patterns and microstructures of the skeletons in (a, c) Al-13Si-20SiCin and (b, d) Al13Si-10SiCex-10SiCin.
•
Fig. 6. Microstructures and pore size distribution of Si-SiCin skeletons: (a) overall microstructure, (b) pore size distribution, (c) macropores, (d) mesopores, (e) Si/SiCex interface and (f) (g) EDS analysis of the blue and green points in e.
•
Fig. 7. (a) Microstructures and EDS (b) point and (c) line analysis of in-situ SiC particles acting as heterogeneous nucleus of Si phases in Al-12Si-4Cu-2Ni alloys. As shown by the white lines, the Si phase in (b) is the enlargement of the one in the rectangle area of (a).
•
Fig. 8. Microstructures and pore size distribution of Si-SiCex-SiCin skeletons: (a) overall microstructure, (b) pore size distribution and (c) (d) the Si/SiCin/SiCex sandwich structure. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (c).
•
Fig. 9. Microstructures of the interface between in-situ SiC and external SiC in Al-SiCex-SiCin
skeletons: (a) (b) SEM analysis, (c) (d) HRTEM analysis and (e) the corresponding FFT results. As shown by the white lines, (b) is the enlargement of the rectangle area in (a). •
Fig. 10. (a) Microstructures and (b) the schematic diagram of Al-SiCex-SiCin skeletons.
•
Fig. 11. Comparison of the microstructures of (a) Si skeletons, (b) Si-SiCex skeletons, (c) Si-SiCin skeletons and (d) Si-SiCex-SiCin skeletons.
•
Fig. 12. Compressive strength of Si, Si-SiCin and Si-SiCex-SiCin skeletons.
Fig. 1. Morphologies of raw materials: (a) Al, (b) Si, (c) graphite and (d) green SiC powders.
Fig. 2. Macrostructures of (a-d) Al-Si-SiC composites and (e-h) their corresponding skeletons: (a) Al33Si and (e) the Si skeleton; (b) Al-13Si-20SiCex and (f) the Si-SiCex skeleton; (c) Al-13Si-20SiCin and (g) the Si-SiCin skeleton; (d) Al-13Si-10SiCex-10SiCin and (h) the Si-SiCex-SiCin skeleton.
Fig. 3. Microstructures of Si skeletons in Al-33Si alloys: (a) (b) overall microstructure, (c) interface betweeen priamry Si particles and (d) eutectic Si. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (d).
Fig. 4. Microstructures of Si-SiCex skeletons in Al-13Si-20SiCex composites: (a) (b) overall microstructure and (c) interface between external SiC and Si. As shown by the white lines, (b) is the enlargement of the rectangle area in (a) and (c) is the enlargement of the rectangle area in (b).
Fig. 5. XRD patterns and microstructures of the skeletons in (a, c) Al-13Si-20SiCin and (b, d) Al-13Si10SiCex-10SiCin.
Fig. 6. Microstructures and pore size distribution of Si-SiCin skeletons: (a) overall microstructure, (b) pore size distribution, (c) macropores, (d) mesopores, (e) Si/SiCex interface and (f) (g) EDS analysis of the blue and green points in e.
Fig. 7. (a) Microstructures and EDS (b) point and (c) line analysis of in-situ SiC particles acting as heterogeneous nucleus of Si phases in Al-12Si-4Cu-2Ni alloys. As shown by the white lines, the Si phase in (b) is the enlargement of the one in the rectangle area of (a).
Fig. 8. Microstructures and pore size distribution of Si-SiCex-SiCin skeletons: (a) overall microstructure, (b) pore size distribution and (c) (d) the Si/SiCin/SiCex sandwich structure. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (c).
Fig. 9. Microstructures of the interface between in-situ SiC and external SiC in Al-SiCex-SiCin skeletons: (a) (b) SEM analysis, (c) (d) HRTEM analysis and (e) the corresponding FFT results. As shown by the white lines, (b) is the enlargement of the rectangle area in (a).
Fig. 10. (a) Microstructures and (b) the schematic diagram of Al-SiCex-SiCin skeletons.
Fig. 11. Comparison of the microstructures of (a) Si skeletons, (b) Si-SiCex skeletons, (c) Si-SiCin skeletons and (d) Si-SiCex-SiCin skeletons.
Fig. 12. Compressive strength of Si, Si-SiCin and Si-SiCex-SiCin skeletons.
Highlights: •
In-situ SiC reinforced Si-SiC skeletons were obtained.
•
Si can nucleate at in-situ SiC resulting in firm interfacial bonding between them.
•
A firm Si/SiCin/SiCex sandwich structure was formed with in-situ SiC as interlayer.
•
The compressive strength of Si-SiC skeletons was elevated obviously.
S0925-8388(19)32963-9
DOI:
https://doi.org/10.1016/j.jallcom.2019.151730
Reference:
JALCOM 151730
To appear in:
Journal of Alloys and Compounds
Received Date: 5 June 2019 Revised Date:
29 July 2019
Accepted Date: 5 August 2019
Please cite this article as: C. Wu, T. Gao, X. Liu, In-situ SiC reinforced Si-SiC 3D skeletons in SiC/Al-Si composites, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/ j.jallcom.2019.151730. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
In-situ SiC reinforced Si-SiC 3D skeletons in SiC/Al-Si composites
Chongchong Wu1, Tong Gao1, Xiangfa Liu1* 1 Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials, Ministry of Education, Shandong University, 17923 Jingshi Road, Jinan 250061, PR China
Abstract: For Al-Si alloys, the three-dimensional (3D) network structure of Si plays a key role on their mechanical properties. In order to improve the microstructure and elevate the strength of the network, insitu SiC was introduced into Al-Si alloys and four kinds of composites and their Si-SiC 3D skeletons were obtained. The microstructures especially the interfacial characteristics were studied in detail. Due to the heterogeneous nucleation role of in-situ SiC, effective interface was formed between in-situ SiC and Si. Moreover, in-situ SiC can grow epitaxially at the surface of external SiC and form a stable Si/SiCin/SiCex sandwich structure with in-situ SiC as the interlayer. As a result, the Si-SiC 3D skeletons became more intact and stable, and their compressive strength was remarkably elevated.
Keywords: Three-dimensional network; Skeleton; Al-Si alloy; In-situ SiC; Interfacial bonding; Strengthening mechanism
1. Introduction Al-Si alloys are the most common used aluminum casting alloys due to their good fluidity and properties. Al-Si alloys can be treated as composites in which ductile α-Al is the matrix and stiff Si is the
*Corresponding author. E-mail: [email protected] (X. Liu).
reinforcement [1-3]. Thus, the morphology and structure of the reinforcement, Si, plays a key role on the mechanical properties of Al-Si alloys. In general, Si phases can form a three-dimensional (3D) network structure during solidification, and it has been proved that more intact of the network, much higher of Al-Si alloys strength [4-6]. The morphologies of the Si 3D network can be changed by many different methods. Heat treatment is the easiest way to change the 3D network, but this method is always detrimental to the network connectivity when used alone [7-9]. With solution tr eatment, eutectic Si with interconnected lamellar structure in as-cast condition will coarsen and spheroidize, resulting in the loss of network connectivity [7]. Modifiers like Sr, Na and TiB2 can also affect the network morphology by affecting the growth of Si during solidification [9-16]. For instance, the addition of about 0.02% wt% Sr can induce the modification of the Si network from lamellar to coral morphology, and thus elevate the ductility and thermal fatigue of the alloys [9, 10]. TiB2 particles can also refine eutectic Si by obstructing solute redistribution and thus improve the Si network [15]. Transition elements like Cu, Fe and Ni can improve the connectivity of Si network by forming intermetallics and supplementing the Si network [17-21]. Asghar [17] found that the addition of Ni can improve the high temperature strength of Al-Si alloys by participating in the formation of a highly interconnected Si network. In reference 20, it is proved that Cu, Ni intermetallics can not only develop the 3D network but also strengthen the Al matrix. In order to achieve desired networks, above mentioned methods can be employed together [9, 11, 14]. For instance, Sr modified Al-12Si alloy is more sensitive to solution treatment than unmodified alloy and thus spheroidize and coarsen more quickly [9]. With the combined addition of modifier Sr and element Cu to Al-7Si, networks consist of fine globular-fibrous eutectic Si and CuAl2 can be obtained, and the mechanical properties of the alloy are much better than the single-addition ones [14]. Al-Si matrix composites are also widely used nowadays and they have much higher strength than Al-Si alloys. For these composites, the Si 3D network also plays an important role on their properties [22-27]. With the addition of Al2O3 short fiber preforms into Al-Si alloys, an interconnected Al2O3-Si network was formed and the elastic modulus and the strength of the alloys was elevated [22]. SiC
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particles can also enhance the formation of a continuous Si-SiC network, resulting in the considerable improvement of the modulus [23]. Besides, the hybrid addition of Al2O3 short fiber and SiC particles has also been studied [24, 25]. Compared with the heat treatment, modifiers and transition elements methods, less attention has been paid on the 3D network of Al-Si matrix composites. In order to study the 3D morphology of the Si network, tomographic imaging techniques are the most useful methods [9, 12, 17, 19, 26]. For instance, focused ion beam/scanning electron microscopy (FIB/SEM) tomography was used to reveal the 3D structure of eutectic Si in reference 9, while energy dispersive synchrotron X-ray diffraction was employed in reference 26. Methods like in situ neutron diffraction can also be used to analysis the Si 3D network [5]. However, these techniques are costly and time-consuming. In our previous works, in-situ SiC was successfully synthesized in Al-Si alloys and it can obviously affect the mechanical properties of Al-Si alloys [28-31]. In this research, the effect of in-situ SiC on the Si 3D networks was studied. At first, four kinds of Al-Si-SiC composites, Al-33Si (primary Si particles reinforced Al-Si composites), Al-13Si-20SiCex (externally added SiC reinforced Al-Si composites), Al13Si-20SiCin (in-situ SiC reinforced Al-Si composites) and Al-13Si-10SiCex-SiCin (externally added SiC and in-situ SiC reinforced Al-Si composites) were obtained. In order to observe the 3D networks, the AlSi-SiC composites were fully etched by hydrochloric acid, and the networks were completely extracted as skeletons. Thus, the skeletons can be studied in detail using routine methods. Results show the impressive effect of in-situ SiC on the formation of intact Si-SiC skeletons.
2. Experiment Commercial Al (99.7 wt% purity), Si (99.9 wt% purity), graphite (99.2 wt% purity, 1.3 µm, Huatai Graphite, China) and green SiC powders (98.5 wt% purity, 75 µm, Sinopharm Chemical Reagent, China) were used in this research. The morphologies of the raw materials are shown in Fig. 1. At first, four kinds of raw powders, with compositions (wt%) of Al-33Si, Al-13Si-20SiCex, Al-27Si-6C and Al20Si-3C-10SiCex were prepared. The raw powders were then mixed in blender mixer for 8 h. After that,
3
they were pressed into cylindrical mould at 270 MPa with cold isostatic press (Taiyuan East Dragon Machinery Co. Ltd, LDJ200/500-380YS, China), with size of Φ60×50 mm. Finally, the compacted samples were sintered in vacuum resistance furnace at 800˚C for 2 h with vacuum degree higher than 103
Pa and cooled to room temperature in the furnace, under which condition the samples can react
completely [28]. The overall reaction during synthesis is shown in equation 1 [28]. According to the compositions of the raw powders, the as-sintered composites were Al-33Si, Al-13Si-20SiCex, Al-13Si20SiCin and Al-13Si-10SiCex-10SiCin, respectively [28]. In these composites, to guarantee the same proportion of the skeletons in the composites, the weight percent of the strengthening particles (Si and SiC) is equal. Si(l) + C(s) = SiC(s)
(1)
To observe the skeletons in these composites, samples were cut from the center of the as-sintered materials with size of 10×10×10 mm, and grinded in standard routines. The samples were then etched by 4 vol% hydrochloric acid sufficiently (about 96 h) until the pH value of the etched solution remain constant. Thus, four kinds of Si-SiC skeletons were prepared. The macrostructures of the composites and their skeletons are shown in Fig. 2. To check the heterogeneous nucleation role of in-situ SiC, Al-13Si-20SiCin was introduced into an Al-12Si-4Cu-2Ni alloy at 800˚C in resistance furnace. Al, Si, Cu and Ni were melted at first, and Al13Si-20SiCin alloy was added after. Refined by high purity argon, the melt was finally cast into iron mould at 800˚C. Metallographic specimens were cut from the center of the cast, and prepared in standard routines. To analyze the phase composition, the skeletons were grinded into powders and tested with X-ray diffraction (XRD, Rigaku D/max-rB, Cu Kα radiation). The microstructures of the skeletons were observed with field emission scanning electron microscopy (FESEM, Hitachi SU8010) equipped with an energy dispersive X-ray spectroscopy (EDS) detector. To conveniently observe the microstructures of the skeletons under scanning electron microscopy (SEM), samples with size of Φ15×10 mm were cut from the composites and were deeply etched (about 10 min) by 4 vol% hydrochloric acid. The fine
4
microstructures of the interfaces were observed with atomic resolution scanning transmission electron microscope (S/TEM, FEI Titan G2 60-300). The S/TEM samples were prepared using focused ion beam (FIB) in regular process. The compressive strength of the skeletons was obtained according to GB/T8489-2008 by universal testing machine (Shimadzu, AG-IC100KN), and at least three samples were tested for each skeleton. To test the pore size distribution of the skeletons, mercury intrusion method (Micromeritics, AutoPore Iv9500), with the stem volume of 1.131 mL, was used. During mercury intrusion, non-wetting mercury was pressed into porous medium, and the pressure is inversely proportional to pore size. The pressure was converted to the equivalent pore size using Washburn equation [32, 33], with the surface tension of 0.485 N/m and the contact angle of 130°. Compared with other methods like gas absorption, mercury intrusion method can test pores with a large range of size, from micropores to macropores [34, 35], and therefore this method is more appropriate to the skeletons in this work which have near-sphere pores range from several nm to several µm.
3. Results and discussion Fig. 2 is the macrostructures of Al-Si-SiC composites and their skeletons. As shown by Fig. 2a to 2d, all the composites, with or without SiC, are compact. Among them, as shown in Fig. 2b, only composites with external SiC show inhomogeneous structure: a strip bright area can be observed in the front face. By contrast, as shown in Fig. 2e to 2h, the skeletons of these composites show much different behaviors. As shown in Fig. 2e, Si skeletons, the skeletons of hypereutectic Al-Si alloys, can almost remain as a whole, but some detached particles can also be observed. Fig. 3 shows the microstructures of Si skeletons. From Fig. 3a and 3b, it is obvious that the skeletons are composed of primary Si with irregular blocky shape and eutectic Si with irregular rod-like shape. From Fig. 3c, the combination between primary Si particles is realized by eutectic Si. However, from Fig. 3a and 3b, this combination is not always valid. Therefore, some Si particles cannot firmly connect to the skeletons, just as the detaching
5
particles shown in Fig. 2e. Besides, from Fig. 3a, the shapes of the pores in the skeletons are not regular and the size of them vary widely, which reflect the nonuniform microstructure of the skeletons. In contrast, as shown in Fig. 2f, when external SiC was added into Al-Si alloys, the original skeletons are destroyed and fell apart. Fig. 4 shows the microstructures of Si-SiCex skeletons, the skeletons of Al-13Si-20SiCex composites. As shown, external SiC distribute among Si phases, but no effective connection exists between them. Because of the mechanical crushing process during the production of externally added green SiC, the surfaces of them are irregular and have no certain crystal orientation, as shown in Fig. 1d [31, 36]. This make it hard for Si to nucleate and grow at the surface of external SiC during the solidification of Al-Si alloy. Therefore, no effective connection exists between external SiC and Si. In other words, the introduction of external SiC breaks the original interfaces between Si phases. Unlike external SiC, when in-situ SiC was introduced into Al-Si alloys, the skeletons in the alloys became more intact without detached particles, which can be seen in Fig. 2g. Moreover, when in-situ SiC was introduced into Al-Si-SiCex composites, the destroyed skeletons in Fig. 2f changed into unbroken ones again, as shown in Fig. 2h. This means that the introduction of in-situ SiC can enhance the Si/Si and SiCex/Si interfacial bonding. To study the enhancing mechanism of in-situ SiC on Si/Si and Si/SiCex interfaces, further research was carried out. Fig. 5 are the XRD patterns and the microstructures of the skeletons in Al-13Si-20SiCin (in-situ SiC reinforced Al-Si composites) and Al-13Si-10SiCex-10SiCin (in-situ SiC reinforced Al-Si-SiCex composites). From Fig. 5a and 5b, only diffraction peaks of Si and SiC can be observed while no peaks of Al were detected. Therefore, the skeletons are composed of Si and SiC. From Fig. 5c and 5d, intercommunicating pores with micrometer size distribute homogenously in both skeletons. As shown in Fig. 5d, external SiC can be observed clearly in the skeletons of Al-13Si-10SiCex-10SiCin. According to the compositions of the composites, the skeletons of Al-13Si-20SiCin and Al-13Si-10SiCex-10SiCin are Si-60SiCin and Si-30SiCex-30SiCin, respectively.
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Fig. 6 shows the microstructures of Si-SiCin skeletons (the skeletons of Al-13Si-20SiCin composites) in Fig. 2g. From the overall structure in Fig. 6a and the pore size distribution in Fig. 6b, bimodal pore distribution was shown: two kinds of pores, macropores with size between 5 µm and 10 µm and mesopores with size between 30 nm and 100 nm, can be observed. As shown in Fig. 6a and 6c, macropores have circular shape in two-dimension and distribute uniformly among the skeletons. By contrast, as shown in Fig. 6d and 6e, mesopores have no certain shape and scatter randomly between the particles. Compared with the Si skeletons in Fig. 3, the shape of the pores becomes more regular and the distribution of them becomes more uniform. Fig. 6f and 6g is the EDS analysis of Si and in-situ SiC in the skeletons. As shown, the light grey phases with regular morphology, hexagonal plates mostly, are in-situ SiC, while the dark grey phases with irregular morphology are Si. No eutectic Si as the one in Fig. 3d can be observed. From Fig. 6d and 6e, in-situ SiC particles with different directions interweave with each other and distribute irregularly among Si phases. Moreover, in Fig. 6e, Si phases attach to in-situ SiC particles with certain angles. The relationship between in-situ SiC and Si phases is determined by their precipitation process. During the sintering of Al-Si-SiCin composites, in-situ SiC with preferred growth orientations began to formed when temperature is higher than 750˚C [28]. Growing SiC particles with different directions may contact with each other, but still can remain their growth orientations, resulting in the interweaved SiC structures in Fig. 6e. During solidification, in-situ SiC can serve as the heterogeneous nucleation sites for Si. Thus, eutectic Si cannot grow freely in Al-Si eutectic structure but grow around in-situ SiC particles as shown in Fig. 6d. Therefore, no dendritic Si can be observed. In addition, owing to twins and multiple twins, the growth orientations of Si can change with certain degrees [37-39]. Thus, Si was nucleated and grew among interweaved in-situ SiC with certain orientation relationships as shown in Fig. 6d and 6e. To check the heterogeneous nucleation role of in-situ SiC, Al-Si-SiCin was added to an Al-12Si4Cu-2Ni alloy with proportion of 1%. The microstructures of the obtained materials are shown in Fig. 7. In Fig. 7a, micron particles can be found in the center of Si phases. Fig. 7b shows one of these particles. EDS analysis of this particle in Fig. 7b and 7c prove that this particle is SiC. Therefore, in-situ SiC can
7
act as the heterogeneous nucleus of Si during Al-Si solidification process. According to the Si-SiC interweaved structure in Fig. 6d and the analysis in Fig. 7, Si should heterogeneously nucleate on specific surface of SiC which has certain orientations or surface energy. From Fig. 6 and reference 29, in-situ synthesized SiC has certain exposed faces, while from Fig. 1d, the exposed faces of external SiC are random. Therefore, in-situ SiC is more appropriate for the heterogeneous nucleation of Si. Compared with the Si skeletons in Fig. 3, two main changes can be observed with in-situ SiC reinforced Si skeletons. For one thing, more effective connection is realized between Si particles with insitu SiC as the attachment. For another, due to the heterogeneous nucleation effect of in-situ SiC, dendritic Si disappears and the pores in the skeletons become more uniform and regular. Similarly, when in-situ SiC was introduced into Si-SiCex skeletons, obvious changes can be observed. Fig. 8 shows the microstructures of Si-SiCex-SiCin skeletons (the skeletons of Al-13Si-10SiCex10SiCin composites). As shown in Fig. 8a, external SiC particles are inlaid in the skeletons with effective connection. From Fig. 8c and 8d, a Si/SiCin/SiCex sandwich structure is formed with in-situ SiC as the interlayer. In addition, as shown in Fig. 8b, due to the introduction of external SiC, pores with trimodal distribution are observed—three sizes of pores with diameter around 150nm, 1µm and 7 µm are present— which is different from the bimodal distribution in Fig. 6b. The interface between in-situ SiC and external SiC is further studied in Fig. 9. From Fig. 9a and 9b, the interface is close and intact. As shown by the HRTEM results in Fig. 9c to 9e, in-situ SiC and external SiC have same crystal structures and parallel directions. This indicates that in-situ SiC interlayer was formed by the epitaxial growth of external SiC. The epitaxial growth of in-situ SiC can also refer to reference 31. From above analysis, the formation of in-situ SiC reinforced Si-SiCin skeletons is illustrated in Fig. 10. During sintering of Al-Si-SiCex-SiCin composites, in-situ SiC interlayers began to grow epitaxially at the surface of external SiC at about 750˚C, while at the same time, in-situ SiC particles were also formed in the melt. During solidification, in-situ SiC interlayers and particles can both act as the heterogeneous
8
nucleus for Si. As a result, effective connection was obtained between external SiC and Si with in-situ SiC as the interlayer, just as the Si/SiCin/SiCex sandwich structure shown in Fig. 8d. Moreover, because in-situ SiC interlayers have same crystal structure with in-situ SiC particles, they can also participate in the formation of the interweaved structures in Fig. 2e. Therefore, as shown in Fig. 10, firm and consistent skeletons were formed. Compared with the Si-SiCex skeletons in Fig. 4, by introducing in-situ SiC, obvious differences were brought, especially the behavior of external SiC particles. In Fig. 4, external SiC particles were isolated from the skeletons, and barely no effective Si/SiCex bonding exists. By contrast, in in-situ SiC reinforced skeletons, owing to the formation of in-situ SiC interlayer, a Si/SiCin/SiCex sandwich structure was obtained, and thus external SiC particles were embedded in the skeletons closely. Fig. 11 shows the microstructures of four kinds of skeletons discussed above. Compared with Si and Si-SiCex skeletons in Fig. 11a and 11b, it is apparently that more effective and intact interfacial bonding is obtained with in-situ SiC reinforced Si skeletons (Fig. 11c) and in-situ SiC reinforced SiSiCex skeletons (Fig. 11d). Besides, the microstructure of the skeletons became more consistent and uniform. From above analysis, in-situ SiC is key for the formation of effective interfacial bonding in Si-SiC skeletons. On the one hand, it can serve as heterogeneous nucleation centers for Si. On the other, it can grow epitaxially at external SiC and form interlayers between external SiC and Si. Due to the formation of enhanced interfacial bonding, the strength of the skeletons has been elevated obviously. As shown in Fig. 12, after introducing in-situ SiC, the compressive strength of the skeletons has been elevated, from 0.39 MPa to 2.12/1.89 MPa, with an elevation of about 5 times. Above results show that introducing in-situ SiC into Si skeletons can guarantee the firm interfacial bonding between SiC and Si, and impressively improve the strength of Si-SiC skeletons. The introduction of in-situ SiC offers a new method to strengthen the 3D network structure in Al-Si alloys. More research, such as the optimization of the skeleton structures by controlling the content and size of in-situ SiC, still need to be carried out in future.
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4 Conclusions In-situ SiC reinforced Si-SiC skeletons in Al-Si-SiC composites were studied in this research. Their microstructures, interfacial properties and strengthening mechanisms were discussed in detail. The following conclusions can be drawn: (1) With the introduction of in-situ SiC, the skeletons in Al-Si-SiC composites became more intact and stable, the interfacial bonding in the skeletons became more effective, and the microstructure of the skeletons became more consistent and uniform. (2) In-situ SiC can serve as the heterogeneous nucleation centers for Si during the solidification of Al-Si alloys, resulting in a firm Si/SiCex interface. Moreover, in-situ SiC can grow epitaxially at the surface of external SiC and act as the interlayers between external SiC and Si, forming a stable Si/SiCin/SiCex sandwich structure. (3) With the introduction of in-situ SiC, the compressive strength of the Si skeletons increased obviously. The compressive strength of Si skeletons has been increased from 0.39 MPa to 2.12MPa, and the broken Si-SiCex skeletons have been changed to intact ones with compressive strength of 1.89 MPa.
Acknowledgement This work was supported by the National Natural Science Foundation of China (No. 51731007) and the key foundation of Shandong Province (No. ZR2016QZ005).
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Figure captions: •
Fig. 1. Morphologies of raw materials: (a) Al, (b) Si, (c) graphite and (d) green SiC powders.
•
Fig. 2. Macrostructures of (a-d) Al-Si-SiC composites and (e-h) their corresponding skeletons: (a) Al-33Si and (e) the Si skeleton; (b) Al-13Si-20SiCex and (f) the Si-SiCex skeleton; (c) Al-13Si20SiCin and (g) the Si-SiCin skeleton; (d) Al-13Si-10SiCex-10SiCin and (h) the Si-SiCex-SiCin skeleton.
•
Fig. 3. Microstructures of Si skeletons in Al-33Si alloys: (a) (b) overall microstructure, (c) interface betweeen priamry Si particles and (d) eutectic Si. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (d).
•
Fig. 4. Microstructures of Si-SiCex skeletons in Al-13Si-20SiCex composites: (a) (b) overall microstructure and (c) interface between external SiC and Si. As shown by the white lines, (b) is the enlargement of the rectangle area in (a) and (c) is the enlargement of the rectangle area in (b).
•
Fig. 5. XRD patterns and microstructures of the skeletons in (a, c) Al-13Si-20SiCin and (b, d) Al13Si-10SiCex-10SiCin.
•
Fig. 6. Microstructures and pore size distribution of Si-SiCin skeletons: (a) overall microstructure, (b) pore size distribution, (c) macropores, (d) mesopores, (e) Si/SiCex interface and (f) (g) EDS analysis of the blue and green points in e.
•
Fig. 7. (a) Microstructures and EDS (b) point and (c) line analysis of in-situ SiC particles acting as heterogeneous nucleus of Si phases in Al-12Si-4Cu-2Ni alloys. As shown by the white lines, the Si phase in (b) is the enlargement of the one in the rectangle area of (a).
•
Fig. 8. Microstructures and pore size distribution of Si-SiCex-SiCin skeletons: (a) overall microstructure, (b) pore size distribution and (c) (d) the Si/SiCin/SiCex sandwich structure. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (c).
•
Fig. 9. Microstructures of the interface between in-situ SiC and external SiC in Al-SiCex-SiCin
skeletons: (a) (b) SEM analysis, (c) (d) HRTEM analysis and (e) the corresponding FFT results. As shown by the white lines, (b) is the enlargement of the rectangle area in (a). •
Fig. 10. (a) Microstructures and (b) the schematic diagram of Al-SiCex-SiCin skeletons.
•
Fig. 11. Comparison of the microstructures of (a) Si skeletons, (b) Si-SiCex skeletons, (c) Si-SiCin skeletons and (d) Si-SiCex-SiCin skeletons.
•
Fig. 12. Compressive strength of Si, Si-SiCin and Si-SiCex-SiCin skeletons.
Fig. 1. Morphologies of raw materials: (a) Al, (b) Si, (c) graphite and (d) green SiC powders.
Fig. 2. Macrostructures of (a-d) Al-Si-SiC composites and (e-h) their corresponding skeletons: (a) Al33Si and (e) the Si skeleton; (b) Al-13Si-20SiCex and (f) the Si-SiCex skeleton; (c) Al-13Si-20SiCin and (g) the Si-SiCin skeleton; (d) Al-13Si-10SiCex-10SiCin and (h) the Si-SiCex-SiCin skeleton.
Fig. 3. Microstructures of Si skeletons in Al-33Si alloys: (a) (b) overall microstructure, (c) interface betweeen priamry Si particles and (d) eutectic Si. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (d).
Fig. 4. Microstructures of Si-SiCex skeletons in Al-13Si-20SiCex composites: (a) (b) overall microstructure and (c) interface between external SiC and Si. As shown by the white lines, (b) is the enlargement of the rectangle area in (a) and (c) is the enlargement of the rectangle area in (b).
Fig. 5. XRD patterns and microstructures of the skeletons in (a, c) Al-13Si-20SiCin and (b, d) Al-13Si10SiCex-10SiCin.
Fig. 6. Microstructures and pore size distribution of Si-SiCin skeletons: (a) overall microstructure, (b) pore size distribution, (c) macropores, (d) mesopores, (e) Si/SiCex interface and (f) (g) EDS analysis of the blue and green points in e.
Fig. 7. (a) Microstructures and EDS (b) point and (c) line analysis of in-situ SiC particles acting as heterogeneous nucleus of Si phases in Al-12Si-4Cu-2Ni alloys. As shown by the white lines, the Si phase in (b) is the enlargement of the one in the rectangle area of (a).
Fig. 8. Microstructures and pore size distribution of Si-SiCex-SiCin skeletons: (a) overall microstructure, (b) pore size distribution and (c) (d) the Si/SiCin/SiCex sandwich structure. As shown by the white lines, (c) is the enlargement of the rectangle area in (a) and (d) is the enlargement of the rectangle area in (c).
Fig. 9. Microstructures of the interface between in-situ SiC and external SiC in Al-SiCex-SiCin skeletons: (a) (b) SEM analysis, (c) (d) HRTEM analysis and (e) the corresponding FFT results. As shown by the white lines, (b) is the enlargement of the rectangle area in (a).
Fig. 10. (a) Microstructures and (b) the schematic diagram of Al-SiCex-SiCin skeletons.
Fig. 11. Comparison of the microstructures of (a) Si skeletons, (b) Si-SiCex skeletons, (c) Si-SiCin skeletons and (d) Si-SiCex-SiCin skeletons.
Fig. 12. Compressive strength of Si, Si-SiCin and Si-SiCex-SiCin skeletons.
Highlights: •
In-situ SiC reinforced Si-SiC skeletons were obtained.
•
Si can nucleate at in-situ SiC resulting in firm interfacial bonding between them.
•
A firm Si/SiCin/SiCex sandwich structure was formed with in-situ SiC as interlayer.
•
The compressive strength of Si-SiC skeletons was elevated obviously.