Next Article in Journal
Managing Surcharge Risk in Strategic Fleet Deployment: A Partial Relaxed MIP Model Framework with a Case Study on China-Built Ships
Previous Article in Journal
Anthropometric Characteristics and Somatotype of Young Slovenian Tennis Players
Previous Article in Special Issue
In Situ Al3BC/Al Composite Fabricated via Solid-Solid Reaction: An Investigation on Microstructure and Mechanical Behavior
 
 
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Effect of Si Addition on Microstructure and Mechanical Properties of SiC Ceramic Fabricated by Direct LPBF with CVI Technology

1
State Key Laboratory of Precision Welding & Joining of Materials and Structures, Harbin Institute of Technology, Harbin 150001, China
2
Henan Key Laboratory of High Performance Carbon Fiber Reinforced Composites, Institute of Carbon Matrix Composites, Henan Academy of Sciences, Zhengzhou 450046, China
3
Zhengzhou Research Institute, Harbin Institute of Technology, Zhengzhou 450003, China
4
Carbon/Carbon Composites Research Center, Shaanxi Key Laboratory of Fiber Reinforced Light Composite Materials, Northwestern Polytechnical University, Xi’an 710072, China
5
School of Materials Science and Engineering, North China University of Water Resources and Electric Power, Zhengzhou 450045, China
6
Mechanics Surfaces and Materials Processing, Arts et Metiers Institute of Technology, F-51006 Châlons-En-Champagne, France
7
Department of Mechanical and Industrial Engineering, Tallinn University of Technology, Ehitajete tee 5, 19086 Tallinn, Estonia
8
Centre for Biomaterials, Cellular and Molecular Theranostics (CBCMT), School of Mechanical Engineering, Vellore Institute of Technology, Vellore 632014, Tamil Nadu, India
*
Author to whom correspondence should be addressed.
Appl. Sci. 2025, 15(15), 8585; https://doi.org/10.3390/app15158585
Submission received: 14 March 2025 / Revised: 9 July 2025 / Accepted: 31 July 2025 / Published: 1 August 2025

Abstract

In this paper, SiC and Si/SiC ceramics were fabricated using direct laser powder bed fusion with chemical vapor infiltration. Their microstructure, mechanical properties and the impacts of silicon addition were analyzed. The incorporation of silicon led to an increase in the relative density of the silicon carbide ceramics from 76.4% to 78.3% and the compression strength increased from 39 ± 13 MPa to 90 ± 8 MPa after laser powder bed fusion with chemical vapor infiltration. The melting and re-solidification of silicon allows the silicon to encapsulate the silicon carbide grains, changing the microstructure and the failure mechanism of the silicon carbide ceramics, resulting in a small amount of silicon residue. In the LPBF-CVI SiC ceramic specimen, the LPBF-formed SiC exhibits a microhardness of 24.2 ± 1.0 GPa. In LPBF-CVI Si/SiC, the spherical dual-phase structure displays a moderately increased hardness (25.9 ± 4.4 GPa), and the CVI-formed SiC exhibits a hardness of 55.3 ± 9.3 GPa.

1. Introduction

Silicon carbide (SiC) ceramic exhibits exceptional properties, including high-temperature resistance, oxidation resistance, wear resistance, chemical stability, a high specific modulus and specific strength, etc. [1,2]. It is useful for various extreme environment applications such as aerospace, electrical power devices, and nuclear fusion and fission environments [3,4]. The covalent nature of Si-C bonds in SiC is the source of its high strength [5]. These same properties also pose significant challenges to their processing. Complex-shaped components require expensive tooling and additional machining. In small-scale production or prototyping, these costs are not cost-effective. Additive manufacturing (AM) techniques are well-suited to the production of complex-shaped parts with small batch sizes or prototypes [6].
SiC can be fabricated utilizing various additive manufacturing technologies, including binder jetting (BJ), direct ink writing (DIW), stereolithography (SL), and selective laser sintering (SLS) [7,8]. However, it should be noted that these methods utilize an organic binder to bond the silicon carbide powder. Before printing, a powder mixing process is necessary, and subsequent debinding and sintering steps are required to remove the binder. This increases the complexity and cost of the process. AM techniques, particularly laser additive manufacturing, offer high material utilization and enable near-net-shape fabrication. It allows the customization of the microstructure, properties, and texturization of the surface, which affects the surface-related properties, such as lubrication [9]. Laser powder bed fusion (LPBF) has been demonstrated to maintain component stability during the forming process, thereby ensuring stable microstructure and properties. Furthermore, significant advantages have been identified in terms of grain refinement and material property enhancement [10,11]. The direct fabrication of silicon carbide ceramics using laser powder bed fusion technology eliminates the need for the pre-mixing of powders and obviates subsequent debinding and sintering processes. This approach streamlines the fabrication process and prevents the sample shrinkage and cracking typically associated with debinding and sintering. However, the direct LPBF process of Carbide ceramics (including SiC-based ceramics) without any binder remains a formidable challenge [12]. The SiC ceramic components fabricated by direct LPBF exhibit low densities and mechanical properties, necessitating posttreatment processes for improving both the density and strength [13]. To address the aforementioned challenges, improvements can be made in terms of the powder feedstock, the preheating platform, the addition of silicon powder, and the densification post-processing. Spherical powders with better flowability and higher packing density are considered to be preferable in the LPBF process, which ultimately results in a near-fully dense final component [14]. The preheating of the powder bed has been shown to reduce the high-temperature gradient, thereby alleviating the residual stress [15]. The formation process can be enhanced by utilizing silicon as a binder, while the residual silicon can be converted to regenerate silicon carbide by post-processing. At present, the widely used post-treatment densification processes include liquid silicon infiltration (LSI), polymer impregnation and pyrolysis (PIP), chemical vapor infiltration (CVI), etc. [16,17,18,19]. CVI is especially suitable for depositing SiC matrix in the preform with interconnected pores compared with PIP and RMI [20]. Furthermore, the ceramic matrix fabricated by CVI is characterized by a dense microstructure, high purity, and excellent properties [21]. Therefore, the subsequent introduction of the SiC matrix into the SiC green parts can be carried out by CVI.
In this paper, the innovative use of spherical silicon carbide powder with a high-temperature (350 °C) preheating platform enables the direct laser powder bed fusion printing of silicon carbide without adding any binder. The impacts of silicon incorporation on the microstructure and mechanical properties of silicon carbide ceramics were investigated using various characterization methods.

2. Materials and Methods

Si/SiC mixed powders were obtained by mixing SiC powders (15–53 μm, purity 99.9%, Shandong Yinuo Rongshen Composite Materials Co., Ltd., Laiwu, China) with Si particles (15–53 μm, Qinghe County Chuangjia Welding Materials Co., Ltd., Xingtai, China) in an inert argon atmosphere. This mixing process was facilitated by a three-dimensional shaker mixer (TURBULA T2GE, Muttenz, Switzerland) at a speed of 72 RPM, and its mass ratio of Si and SiC was 1/9. The SiC powder exhibited a near-spherical morphology (Figure 1a). From Figure 1c,d, it can be seen that there are pores inside the powder, which may adversely affect the relative density of the formed specimens. The size distribution of the powder is shown in Figure 1b. The Si powder exhibited a near-spherical morphology, as shown in Figure 2a, and the corresponding EDS map is shown in Figure 2b.
SiC and Si/SiC cuboid samples with a dimension of 5 mm × 5 mm and a thickness of about 12.5 mm were fabricated using the LPBF equipment HBD 150 (Shanghai Hanbang United 3D Tech Co., Ltd., Shanghai, China) based on optimized LPBF parameters. The LPBF processing parameters were obtained through initial optimization design experiments. The process parameters were set as shown in Table 1.
The SiC ceramic specimens and the Si/SiC ceramic specimens fabricated via direct LPBF will hereafter be referred to as ‘LPBF SiC’ and ‘LPBF Si/SiC’, respectively. The LPBF specimens underwent chemical vapor infiltration treatment to obtain the final SiC ceramic samples. The process parameters were set as shown in Table 2. The LPBF SiC ceramic specimens and the LPBF Si/SiC ceramic specimens subjected to the CVI densification will hereafter be denoted as ‘LPBF-CVI SiC’ and ‘LPBF-CVI Si/SiC’, respectively.
The phase identification of the samples was carried out through X-ray diffraction (XRD, Empyrean, Malvern, UK) using Cu Kα radiation. The measurement range of XRD was set from 20° to 110°, the step size was set to 0.0131303°, and the total number of steps was 6854. X-ray computed tomography (Micro-CT, ZEISS Xradia 515 Versa, Oberkochen and Jena, Germany) was conducted to identify the pore distribution in the specimens. The resulting 3D images with a voxel size of 1.5 μm were analyzed using Dragonfly V2024.1 software. A scanning electron microscope (SEM, Gemini SEM 360, Oberkochen, Germany) was used for microstructure and fracture surface characterization. The SEM device was equipped with an electron backscatter diffraction device (EBSD, OXFORD Symmetry S3, Abingdon, UK). The microstructure was observed using a transmission electron microscope (TEM, Talos F200S, Hillsboro, OR, USA) equipped with an energy-dispersive X-ray (EDS, OXFORD X-MaxN 80T IE250, Abingdon, UK) detector. The TEM specimens were precisely extracted from designated areas using a focused ion beam (FEI Strata 400s Dual Beam FIB, Hillsboro, OR, USA). Room-temperature compression tests were performed following the GB/T 8489-2006 standard [22] using a universal testing machine (LD26.504, Lishi Instruments Co., Ltd., Shanghai, China). For each set of parameters, five specimens were tested for compressive strength. A Vickers hardness tester (QHV-1000SPTA, Laizhou weiyi Experimental Machine Manufacturing Co., Ltd., Yantai, China) was used to test the microhardness of the LPBF specimens. The test force was 500 gf, the holding time was 15 s, and the measured Vickers hardness values were uniformly represented by the symbol HV0.5. Nanoindentation testing was performed using a Nanoindenter (CV500, NANOVEA Inc., Irvine, CA, USA) with a maximum load of 50 mN, a 100 mN/min loading rate, and a 5 s dwell time at the peak load; the distance between each indentation was set to 5 μm.

3. Results and Discussion

The X-ray diffraction patterns presented in Figure 3a illustrate that the diffraction peaks of silicon were detected after the direct LPBF process, indicating that a small amount of SiC decomposed to liquid Si and solid C. After the CVI process, the diffraction peaks of silicon disappeared, and the positions of the peaks were consistent with the diffraction peaks of the SiC powder, indicating that the high-temperature process of CVI converted free silicon and carbon to silicon carbide again, and ultimately the high-purity and highly crystalline silicon carbide ceramics were obtained. After the addition of 10 wt% silicon powder, the diffraction peaks of the silicon phase were detected from the direct LPBF specimen, which is consistent with the diffraction peaks of the Si powder. After the CVI processing, it was found that the surface of the LPBF-CVI specimen is β-sic, which is a typical crystalline structure of CVI-SiC [21]. The diffraction peaks of silicon in the interior were reduced and the diffraction intensity was weakened. The silicon diffraction peaks at (311), (422), and (440) disappeared, while the intensities of the (111), (220), (311), and (400) peaks were significantly reduced. This result may be due to an augmentation in the silicon carbide content and a decrease in the relative content of silicon after CVI, so the diffraction intensity is reduced. It also indicates that not all the silicon was reacted to yield silicon carbide during the CVI process. There is a small amount of silicon residue in the ceramics. Compared with previously published results, the SiC ceramics fabricated in this work exhibit a very low residual silicon content, and the CVI processing further reduces the relative silicon fraction. In contrast, densification methods such as PIP and LSI tend to increase the amount of residual silicon [23,24].
Figure 4a–d shows the three-dimensional morphology of LPBF SiC, LPBF-CVI SiC, LPBF Si/SiC, and LPBF-CVI Si/SiC samples. It is essential to highlight that the grey region is the SiC ceramic, while the dark region is porous. Columnar structures along the BD direction are present inside the LPBF SiC sample, and between the columnar structures are large-volume connecting pores. The distance between the columnar structures is consistent with the scanning spacing. After CVI treatment, the surface of the specimen forms a shell layer with a thickness of a few hundred micrometers, and the relative density of the specimen was increased from 65.7% to 76.4%. Figure 4c demonstrates that adding silicon improves the specimens’ relative densities from 65.7% to 70.5%. Moreover, the addition of silicon resulted in the refinement of the columnar structure morphology of the specimen, with a concomitant decrease in the prevalence of small-sized synapse-like morphologies. After the CVI processing, the relative density of the Si/SiC ceramic specimens was increased from 70.5% to 78.3%, as plotted in Figure 4d.
As demonstrated in Figure 5a,b, the SEM and EBSD results after CVI densification reveal that the microstructure consists of original large equiaxed grains surrounded by the typical small, columnar grains formed by CVI densification. The average grain size of the original grains is 2.45 μm. The small, columnar grains in the region of the CVI-SiC matrix were undefined in the EBSD maps, and it can be deduced that the time step limits of EBSD made the identification of any nanocrystallites within the CVI-SiC matrix difficult. Thus, lamellas were sliced using the FIB technique in each of the LPBF and CVI formed regions for TEM analysis. The results are shown in Figure 5c,d. Regions of interest are highlighted by boxes in Figure 5a. The SiC matrix formed by CVI has nanoscale needle-like grains, and the diffraction pattern results in polycrystalline rings without amorphous features. The diffraction patterns obtained from both the original grain and the Chemical Vapor Infiltration (CVI) matrix confirmed the presence of highly crystalline 3C-SiC, with no observable amorphous regions or pockets. The formation of 3C-SiC indicates that the SiC powder undergoes a partial transformation after laser sintering. The binding of SiC particles is considered a result of the formation of a metastable Si/C mixture as a result of the incongruent melting of SiC. The relative density of the LPBF SiC sample indicates that the samples were still in the initial stages of sintering, where material transport is primarily interparticle neck growth. Thus, a melting and recrystallization mechanism would explain the formation of 3C-SiC within particular regions, differing from the feedstock 6H-SiC powder [25].
After the silicon addition, it was observed from Figure 6a,b that the original LPBF microstructure preserved the spherical shape, similar to sic powders, since the resolidified Si can cover the SiC particles and bind the SiC particles together. The EBSD result shows that the average grain size of the original grains slightly increased to 3.23 μm. The microstructure of the CVI-SiC matrix area is similar to LPBF-CVI-SiC specimens. Figure 6c depicts the HRTEM image of the microstructure, which comprises large SiC particles embedded within a continuous Si matrix. EDS maps of C and Si are shown in Figure 6d. Selected area diffraction patterns obtained from both the SiC particles and the Si matrix confirmed the presence of Si and 6H-SiC, as depicted in Figure 6c. Residual silicon may adversely affect the chemical stability and high-temperature performance of silicon carbide ceramic parts.
As illustrated in Figure 7(a1), the Vickers hardness values of SiC in the LPBF-CVI SiC specimens exhibit variation across different regions. In the LPBF-formed SiC region, the hardness value is recorded as 1297 HV0.5, whereas in the CVI-formed SiC region, it increases to 2651 HV0.5. Corresponding results are observed in LPBF-CVI Si/SiC specimens, as shown in Figure 7(b1), with the hardness value in the LPBF-formed Si/SiC region recorded as 1798 HV0.5, and 3315 HV0.5 in the CVI-formed SiC region. To further investigate the variation in hardness from the LPBF-formed SiC region to the CVI-formed SiC matrix region, a series of nanoindentation tests were performed across the transition zone between the two regions. The nano-hardness of the LPBF-formed SiC region in the LPBF-CVI SiC ceramic specimen is lower (24.2 ± 1.0 GPa) due to the larger grain size and the presence of pores. The average nano-hardness of the CVI-formed nanoneedle-like SiC region is 39.2 ± 8.4 GPa Figure 7(a4). The hardness of the LPBF-CVI Si/SiC ceramic specimen with a spherical dual-phase structure is slightly higher (25.9 ± 4.4 GPa), and the average nano-hardness of the CVI-formed SiC region is 55.3 ± 9.3 GPa.
The stress–strain curve obtained from the compressive strength test is depicted in Figure 8(a1). The results indicate that the compressive strength of the LPBF SiC specimen is 5.89 ± 1.21 MPa, while its compressive strength is greatly increased after the CVI densification post-treatment, reaching 38.59 ± 12.99 MPa. Figure 8(a2) illustrates that the interior is not completely densified after CVI treatment and still retains a columnar structure. There are also porosity defects inside the columnar structure. After compression, a fracture occurs at the columnar structure. The fracture morphology of the LPBF-CVI SiC specimen is illustrated in Figure 8(a3), showcasing a typical brittle fracture. The LPBF grain particles exhibited planar and smooth cleavage planes, indicative of trans-granular fracture features; tiny pores around grain particles were observed, hindering the compressive strength of the sample.
In Figure 8(b1), it is illustrated that the LPBF-fabricated Si/SiC specimens exhibit an ultimate compressive strength of 19 ± 8 MPa. When compared to the LPBF Si/SiC samples, the LPBF-CVI Si/SiC specimens display a significantly higher ultimate compressive strength, achieving a value of 90 ± 8 MPa. As can be seen in Figure 8(b2), after Si addition, the specimen is similarly incompletely densified internally, and the fracture penetrates through the original spheroidal microstructure and is deep inside the microstructure, significantly enhancing its mechanical properties. With the addition of Si powders, resolidified Si could fill in the pores and bind the SiC grain particles together, as shown in Figure 8(b3), with SiC formed by CVI wrapping around the original spheroidal microstructure, resulting in a highly dense structure; the multiple river-like patterns indicate trans-granular fracture. Combined with the relative density improvement, the addition of Si significantly increased the load-bearing area, thereby improving the compressive strength. Figure 9a,b schematically explain the fracture mechanisms of the two distinct ceramic specimens.
In summary, the relative densities and compressive strengths of the SiC ceramics and Si/SiC composite ceramics fabricated by LPBF and LPBF-CVI are summarized in Table 3.

4. Conclusions

In this work, the SiC ceramic and Si/SiC ceramic were manufactured via direct LPBF and CVI. A comprehensive investigation of their microstructure characterization and strengthening mechanisms was conducted. The principal findings are as follows:
(1)
The densities of LPBF-SiC specimens were 65.7%; the compressive strength was 6 ± 1 MPa. After CVI, the density increased to 76.4% and the compressive strength increased to 39 ± 13 MPa.
(2)
A 70.5% density and 19 ± 8 MPa compressive strength were obtained for LPBF-Si/SiC specimens after adding Si. After CVI, the density of the LPBF-CVI Si/SiC specimen was increased to 78.3% and the compressive strength was increased to 90 ± 8 MPa.
(3)
During the LPBF-forming process, silicon carbide undergoes slight decomposition, and the subsequent CVI densification treatment enables the decomposed silicon to re-react and regenerate silicon carbide.
(4)
The nano-hardness values of the LPBF-CVI and LPBF-CVI SiC ceramic specimens differ in different regions. The nano-hardness values of the LPBF-formed SiC and Si/SiC regions are 24.2 ± 1.0 GPa and 25.9 ± 4.4 GPa, respectively, whereas the nano-hardness values of the CVI-formed SiC region can be as high as 55.3 ± 9.3 GPa.
High-purity and crystalline silicon carbide can be fabricated using direct LPBF combined with CVI. The addition of silicon will increase the relative density, change the microstructure morphology, and improve the compressive strength, but there will be a small amount of silicon residue. This research is anticipated to serve as a foundation for the advancement of SiC ceramics with refined microstructures and enhanced strength through direct LPBF.

Author Contributions

Methodology, N.K.; Formal analysis, J.P.; Investigation, J.Z. and L.L.; Writing—original draft, Y.W. and K.G.P.; Writing—review and editing, P.W., Y.Z., K.G.P., X.W. and M.E.M. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by National Key Research and Development Program of China (Grant No. 2024YFB4608600), National Science and Technology Major Project (Grant No. J2022-VI-0011-0042), the Joint Fund of Henan Province Science and Technology R&D Program (Grant No. 225200810002 and 235200810030), the program for overseas high-level talents introduction of Henan Province (HNGD2025040) and the Startup Research Fund of Henan Academy of Sciences (No. 242021011).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflict of interest.

References

  1. Zhang, J.; Zhang, Y.; Fu, Y.; Chen, R.; Li, T.; Hou, X.; Li, H. Research progress in chemical vapor deposition for high-temperature anti-oxidation/ablation coatings on thermal structural composites. Compos. Part B Eng. 2025, 291, 112015. [Google Scholar] [CrossRef]
  2. Fu, Y.; Zhang, Y.; Li, T.; Han, L.; Miao, Q. Effect of SiC on the anti-ablation resistance and flexural strength of (Hf-Ta-Zr)C-C/C composites. J. Eur. Ceram. Soc. 2024, 44, 107–118. [Google Scholar] [CrossRef]
  3. Lv, J.S.; Li, W.; Li, Z.L.; Fu, Y.Q.; Ma, Y.W.; Guo, L.X.; Li, J.C.; Li, T.; Zhang, Y.L. Ablation-resistant (Hf,Zr)B2-SiC composite coating with alternating lamellar architecture by one-step atmospheric plasma spraying. Compos. Part B Eng. 2025, 297, 112302. [Google Scholar] [CrossRef]
  4. Xu, M.; Girish, Y.R.; Rakesh, K.P.; Wu, P.; Manukumar, H.M.; Byrappa, S.M.; Udayabhanu; Byrappa, K. Recent advances and challenges in silicon carbide (SiC) ceramic nanoarchitectures and their applications. Mater. Today Commun. 2021, 28, 102533. [Google Scholar] [CrossRef]
  5. Tanaka, H. Silicon Carbide Ceramics-1: Fundamental and Solid Reaction. In Sintering of Silicon Carbide; Sömiya, S., Inomata, Y., Eds.; Springer: Dordrecht, The Netherlands, 1991; pp. 213–238. [Google Scholar]
  6. Gibson, I.; Rosen, D.; Stucker, B.; Khorasani, M. Additive Manufacturing Technologies; Springer: Berlin/Heidelberg, Germany, 2021. [Google Scholar]
  7. Chen, Z.W.; Li, Z.Y.; Li, J.J.; Liu, C.B.; Lao, C.S.; Fu, Y.L.; Liu, C.Y.; Li, Y.; Wang, P. He, Y. 3D printing of ceramics: A review. J. Eur. Ceram. Soc. 2019, 39, 661–687. [Google Scholar] [CrossRef]
  8. Sun, J.; Ye, D.; Zou, J.; Chen, X.; Wang, Y.; Yuan, J.; Liang, H.; Qu, H.; Binner, J.; Bai, J. A review on additive manufacturing of ceramic matrix composites. J. Mater. Sci. Technol. 2023, 138, 1–16. [Google Scholar] [CrossRef]
  9. Bartkowiak, T.; Peta, K.; Królczyk, J.B.; Niesłony, P.; Bogdan-Chudy, M.; Przeszłowski, Ł.; Trych-Wildner, A.; Wojciechowska, N.; Królczyk, G.M.; Wieczorowski, M. Wetting properties of polymer additively manufactured surfaces-Multiscale and multi-technique study into the surface-measurement-function interactions. Tribol. Int. 2025, 202, 110394. [Google Scholar] [CrossRef]
  10. Wang, P.; Eckert, J.; Prashanth, K.-G.; Wu, M.-W.; Kaban, I.; Xi, L.-X.; Scudino, S. A review of particulate-reinforced aluminum matrix composites fabricated by selective laser melting. Trans. Nonferrous Met. Soc. China 2020, 30, 2001–2034. [Google Scholar] [CrossRef]
  11. Shen, Z.; Su, H.; Yu, M.; Cao, Y.; Guo, Y.; Jiang, H.; Liu, Y.; Li, X.; Dong, D.; Yang, P.; et al. Unveiling exotic multi-scale microstructure transformation and crack formation mechanisms in eutectic ceramic composite by laser powder bed fusion. Compos. Part B Eng. 2025, 288, 111883. [Google Scholar] [CrossRef]
  12. Xu, T.; Cheng, S.; Jin, L.; Zhang, K.; Zeng, T. High-temperature flexural strength of SiC ceramics prepared by additive manufacturing. Int. J. Appl. Ceram. Technol. 2020, 17, 438–448. [Google Scholar] [CrossRef]
  13. Deckers, J.; Vleugels, J.; Kruth, J.P. Additive Manufacturing of Ceramics: A Review. J. Ceram. Sci. Technol. 2014, 5, 245–260. [Google Scholar]
  14. Nan, W.; Pasha, M.; Bonakdar, T.; Lopez, A.; Zafar, U.; Nadimi, S.; Ghadiri, M. Jamming during particle spreading in additive manufacturing. Powder Technol. 2018, 338, 253–262. [Google Scholar] [CrossRef]
  15. Wang, K.; Yin, J.; Chen, X.; Wang, L.; Xiao, H.; Liu, X.; Huang, Z. Advances on direct selective laser printing of ceramics: An overview. J. Alloy. Compd. 2024, 975, 172821. [Google Scholar] [CrossRef]
  16. Fleisher, A.; Zolotaryov, D.; Kovalevsky, A.; Muller-Kamskii, G.; Eshed, E.; Kazakin, M.; Popov, V. Reaction bonding of silicon carbides by Binder Jet 3D-Printing, phenolic resin binder impregnation and capillary liquid silicon infiltration. Ceram. Int. 2019, 45, 18023–18029. [Google Scholar] [CrossRef]
  17. Du, W.; Ma, B.; Thomas, J.; Singh, D. Concurrent reaction-bonded joining and densification of additively manufactured silicon carbide by liquid silicon infiltration. J. Eur. Ceram. Soc. 2023, 43, 2345–2353. [Google Scholar] [CrossRef]
  18. Lv, X.; Ye, F.; Cheng, L.; Fan, S.; Liu, Y. Fabrication of SiC whisker-reinforced SiC ceramic matrix composites based on 3D printing and chemical vapor infiltration technology. J. Eur. Ceram. Soc. 2019, 39, 3380–3386. [Google Scholar] [CrossRef]
  19. Gai, W.H.; Zhang, Y.L.; Zhang, J.; Chen, H.; Chen, G.H.; Kong, J.A. Oxidation behavior of supersonic air plasma sprayed Yb2SiO5/Si and Yb2Si2O7/Si coating for CVD-SiC coated C/C composites in wet oxygen at 1773 K: Experimental and first-principle calculation. Ceram. Int. 2025, 51, 11088–11096. [Google Scholar] [CrossRef]
  20. Zhang, H.; Yang, Y.; Liu, B.; Huang, Z. The preparation of SiC-based ceramics by one novel strategy combined 3D printing technology and liquid silicon infiltration process. Ceram. Int. 2019, 45, 10800–10804. [Google Scholar] [CrossRef]
  21. Terrani, K.; Jolly, B.; Trammell, M. 3D printing of high-purity silicon carbide. J. Am. Ceram. Soc. 2020, 103, 1575–1581. [Google Scholar] [CrossRef]
  22. GB/T 8489-2006; Test Method for Compressive Strength of Fine Ceramics (Advanced Ceramics, Advanced Technical Ceramics). China Standards Press: Beijing, China, 2006.
  23. Chaugule, P.S.; Du, W.; Kamath, R.R.; Barua, B.; Messner, M.C.; Singh, D. Reliability comparisons between additively manufactured and conventional SiC–Si ceramic composites. J. Am. Ceram. Soc. 2024, 107, 3117–3133. [Google Scholar] [CrossRef]
  24. Meyers, S.; De Leersnijder, L.; Vleugels, J.; Kruth, J.-P. Direct laser sintering of reaction bonded silicon carbide with low residual silicon content. J. Eur. Ceram. Soc. 2018, 38, 3709–3717. [Google Scholar] [CrossRef]
  25. Lamm, B.W.; Karakoc, O.; Mao, K.; Koyanagi, T.; Liu, J.; Katoh, Y. Phase separation during the direct powder bed fusion of SiC. Int. J. Appl. Ceram. Technol. 2024, 21, 1722–1734. [Google Scholar] [CrossRef]
Figure 1. SiC powders of (a) SEM micrographs; (b) Powder particle size distribution; (c) Pores distribution; (d) Micro-CT result.
Figure 1. SiC powders of (a) SEM micrographs; (b) Powder particle size distribution; (c) Pores distribution; (d) Micro-CT result.
Applsci 15 08585 g001
Figure 2. Si powders of (a) SEM micrographs for Si powders; (b) corresponding EDS maps.
Figure 2. Si powders of (a) SEM micrographs for Si powders; (b) corresponding EDS maps.
Applsci 15 08585 g002
Figure 3. (a) XRD patterns of SiC powder, LPBF SiC parts, LPBF-CVI SiC parts; (b) XRD patterns of Si powder, LPBF Si/SiC parts, and LPBF-CVI Si/SiC parts.
Figure 3. (a) XRD patterns of SiC powder, LPBF SiC parts, LPBF-CVI SiC parts; (b) XRD patterns of Si powder, LPBF Si/SiC parts, and LPBF-CVI Si/SiC parts.
Applsci 15 08585 g003
Figure 4. (a,b) Three-dimensional morphology of the LPBF SiC specimen and the LPBF-CVI SiC specimen; (c,d) Three-dimensional morphology of the LPBF Si/SiC specimen and the LPBF-CVI Si/SiC specimen.
Figure 4. (a,b) Three-dimensional morphology of the LPBF SiC specimen and the LPBF-CVI SiC specimen; (c,d) Three-dimensional morphology of the LPBF Si/SiC specimen and the LPBF-CVI Si/SiC specimen.
Applsci 15 08585 g004
Figure 5. (a) SEM images of LPBF-CVI SiC ceramic. (b) Inverse pole figure map (inset: corresponding grain size diameter in IPF maps). (c,d) TEM images and diffraction patterns of the large SiC grain and the CVI-SiC matrix.
Figure 5. (a) SEM images of LPBF-CVI SiC ceramic. (b) Inverse pole figure map (inset: corresponding grain size diameter in IPF maps). (c,d) TEM images and diffraction patterns of the large SiC grain and the CVI-SiC matrix.
Applsci 15 08585 g005
Figure 6. (a) SEM images of LPBF-CVI Si/SiC ceramic. (b) Inverse pole figure map (inset: corresponding grain size diameter in IPF maps). (c) The distribution state of SiC particles within the Si matrix (inset: Diffraction patterns of Si and SiC from (c), regions of interest are highlighted by boxes). (d) The elemental surface distribution results within (c) region.
Figure 6. (a) SEM images of LPBF-CVI Si/SiC ceramic. (b) Inverse pole figure map (inset: corresponding grain size diameter in IPF maps). (c) The distribution state of SiC particles within the Si matrix (inset: Diffraction patterns of Si and SiC from (c), regions of interest are highlighted by boxes). (d) The elemental surface distribution results within (c) region.
Applsci 15 08585 g006
Figure 7. LPBF-CVI SiC specimens of (a1) Vickers Microhardness map; (a2) Variation in nano-hardness from LPBF-forming SiC region to CVI-forming SiC region; (a3,a4) nano-hardness of the LPBF-forming SiC region and CVI-forming SiC region; LPBF-CVI Si/SiC specimens of (b1) Vickers Microhardness map; (b2) Variation in nano-hardness from LPBF-forming SiC region to CVI-forming SiC region; (b3,b4) nano-hardness of the LPBF-forming Si/SiC region and CVI-forming SiC region.
Figure 7. LPBF-CVI SiC specimens of (a1) Vickers Microhardness map; (a2) Variation in nano-hardness from LPBF-forming SiC region to CVI-forming SiC region; (a3,a4) nano-hardness of the LPBF-forming SiC region and CVI-forming SiC region; LPBF-CVI Si/SiC specimens of (b1) Vickers Microhardness map; (b2) Variation in nano-hardness from LPBF-forming SiC region to CVI-forming SiC region; (b3,b4) nano-hardness of the LPBF-forming Si/SiC region and CVI-forming SiC region.
Applsci 15 08585 g007
Figure 8. LPBF-CVI SiC specimens of (a1) Compressive stress-strain curves; (a2) Fracture surface in low magnification; (a3) Fracture surface in high magnification; LPBF-CVI Si/SiC specimens of (b1) Compressive stress-strain curves; (b2) Fracture surface in low magnification; (b3) Fracture surface in high magnification.
Figure 8. LPBF-CVI SiC specimens of (a1) Compressive stress-strain curves; (a2) Fracture surface in low magnification; (a3) Fracture surface in high magnification; LPBF-CVI Si/SiC specimens of (b1) Compressive stress-strain curves; (b2) Fracture surface in low magnification; (b3) Fracture surface in high magnification.
Applsci 15 08585 g008
Figure 9. Schematic illustration of the fracture of (a) LPBF-CVI SiC specimens; (b) LPBF-CVI Si/SiC specimens.
Figure 9. Schematic illustration of the fracture of (a) LPBF-CVI SiC specimens; (b) LPBF-CVI Si/SiC specimens.
Applsci 15 08585 g009
Table 1. The LPBF processing parameters.
Table 1. The LPBF processing parameters.
Laser PowerScanning SpeedLayer ThicknessScanning SpacingPreheating Temperature
3005006070350
Table 2. The CVI process parameters.
Table 2. The CVI process parameters.
Pressure (kPa)Temperature (°C)CH3Cl3Si (mL/min)H2 (mL/min)Ar (mL/min)Duration (Hours)
5–15900–120050–200500–1500400–8002–8
Table 3. Comparison of relative density and compressive strength in LPBF and LPBF-CVI ceramics.
Table 3. Comparison of relative density and compressive strength in LPBF and LPBF-CVI ceramics.
CompositeRelative Density (%)Compressive Strength (MPa)
LPBF SiC65.76 ± 1
LPBF-CVI SiC76.439 ± 13
LPBF Si/SiC70.519 ± 8
LPBF-CVI Si/SiC78.390 ± 8
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

Wang, Y.; Wang, P.; Li, L.; Zhang, J.; Zhang, Y.; Peng, J.; Wang, X.; Kang, N.; El Mansori, M.; Prashanth, K.G. Effect of Si Addition on Microstructure and Mechanical Properties of SiC Ceramic Fabricated by Direct LPBF with CVI Technology. Appl. Sci. 2025, 15, 8585. https://doi.org/10.3390/app15158585

AMA Style

Wang Y, Wang P, Li L, Zhang J, Zhang Y, Peng J, Wang X, Kang N, El Mansori M, Prashanth KG. Effect of Si Addition on Microstructure and Mechanical Properties of SiC Ceramic Fabricated by Direct LPBF with CVI Technology. Applied Sciences. 2025; 15(15):8585. https://doi.org/10.3390/app15158585

Chicago/Turabian Style

Wang, Yipu, Pei Wang, Liqun Li, Jian Zhang, Yulei Zhang, Jin Peng, Xingxing Wang, Nan Kang, Mohamed El Mansori, and Konda Gokuldoss Prashanth. 2025. "Effect of Si Addition on Microstructure and Mechanical Properties of SiC Ceramic Fabricated by Direct LPBF with CVI Technology" Applied Sciences 15, no. 15: 8585. https://doi.org/10.3390/app15158585

APA Style

Wang, Y., Wang, P., Li, L., Zhang, J., Zhang, Y., Peng, J., Wang, X., Kang, N., El Mansori, M., & Prashanth, K. G. (2025). Effect of Si Addition on Microstructure and Mechanical Properties of SiC Ceramic Fabricated by Direct LPBF with CVI Technology. Applied Sciences, 15(15), 8585. https://doi.org/10.3390/app15158585

Note that from the first issue of 2016, this journal uses article numbers instead of page numbers. See further details here.

Article Metrics

Back to TopTop