Single-step additive manufacturing of silicon carbide through laser-induced phase separation

Omer Karakoc (  omerkarakoc2018@gmail.com ) Oak Ridge National Laboratory https://orcid.org/0000-0001-9512-6156 Keyou Mao Oak Ridge National Laboratory Jianqi Xi University of Wisconsin-Madison Takaaki Koyanagi Oak Ridge National Laboratory https://orcid.org/0000-0001-7272-4049 Jian Liu Polaronyx Company Izabela Szlufarska University of Wisconsin–Madison Yutai Katoh Oak Ridge National Laboratory https://orcid.org/0000-0001-9494-5862


Introduction
Silicon carbide (SiC) has potential as a structural material for use in extreme environments such as space and nuclear applications owing to its strong corrosion resistance, high-temperature strength, excellent damage irradiation tolerance, adequate scattering cross-sections, and low neutron absorption 1,2,3,4 . Unlike metals or alloys 5,6 , however, machining, near net shaping of SiC via-conventional machining are impractical and extremely difficult due to their brittleness and chemical stability 7,8,9 . Conventional machining consumes large amount of energy to shape SiC due to high sintering temperatures above 2000 °C 9 . Additive manufacturing (AM) promises a cost-and energy-effective approach to solving these issues and is a strategy for developing nextgeneration parts for advanced nuclear applications 7 , because it significantly reduces the amount of waste produced in the process 10 and enables rapid prototyping and fabrication of parts with complex geometries. Thus, additive manufacturing of SiC is fast-growing technology for wide variety of applications 11 . AM technology of SiC will be revolutionary, but dense and high purity SiC part by AM have not been realized due to strong covalent nature of SiC, SiC sublimation rather than melting at high temperatures 7 . Full capability of SiC component is only achievable with nuclear-grade SiC, which is highly crystalline and dense and pure 12 . AM of SiC generally involves preforming a green body and densification step 7 . Most widely used processing options for AM of SiC are wet processing (sterolithography, gel casting, and direct ink writing) and dry processing (SLS, laminated object manufacturing, and binder jet printing) 7 . Another AM process is laser-induced chemical vapor deposition (LCVD), which uses reagent gases 13 . Heat of focused laser results in decomposition of reagent gases in which AM part is produced. Technological challenges lie in SiC densification process: part size for LCVD; high densification temperature for liquid-phase sintering; and volume shrinkage for powder sintering and pre-ceramic polymer pyrolysis 7 . AM of SiC by SLS involves reaction sintering of silicon and carbon, which results in formation SiO2 impurities 7 .
Therefore, in the present study, single-step additive manufacturing of SiC has been developed using laser powder bed fusion (LPBF) without use of any sintering additives. Pulsed-LPBF joins materials by consolidating successive layers of powder and selectively sintering them using a highenergy pulsed laser to fabricate final components from 3D model data 14,15,16 . Thus, it is possible to make objects with arbitrary geometries without the need to adapt the conventional production process itself. This approach enables LPBF to fabricate complex 3D parts with high accuracy without extensive tooling and without the geometric limitations inherent in typical subtractive manufacturing processes 17,18 . The capability to process a wide variety of materials with a large range of mechanical and physical properties will enable a broad range of applications in the aerospace, nuclear, biology, and medical industries 17,18,19,20,21,22 . Despite its great advantages, concerns over AM object quality and consistency limit the widespread utilization of LPBF 23 .
Large differences in the mechanical properties of AM objects pose challenges for certification authorities 24 and designers 25,26 . Sintering involves neck formation between adjacent powders to lower the free energy while powder particles grow. These regions can occur many times in a single AM part-typically close to the fusion of powder particles, where the influences on chemical, mechanical, and physical properties are the most pronounced. Thus, providing insight on the lasermatter interaction process in those regions could lead to remarkable outcomes for the quality and consistency of SiC parts fabricated by AM 27 .
Nonetheless, there has not been extensive research to obtain significant surface information such as the microstructural evolution and binding mechanisms of single SiC powder particles under high-energy short-pulse laser irradiation. We conducted detailed microstructural characterization, which led to findings that explains the physical process of SiC AM and important laser-material interactions. Those topics are ideal for investigation by transmission electron microscopy (TEM), transmission Kikuchi diffraction (TKD), Raman spectroscopy, and scanning electron microscopy (SEM). In this study, XRD, TEM, TKD, SEM, and Raman spectroscopy are carried out to explore fiber laser-SiC powder particle interactions during AM processing and elucidate the binding mechanisms that result in the consolidation of SiC powders. Complementary microstructural characterization enables a deeper understanding of general trends in laser-SiC powder interactions and elucidates the binding mechanism and phase separations detailed in the experimental efforts.
A successful mitigation has been implemented to consolidate SiC powders through phase separation of SiC (Fig. 1). The experimental observations, demonstrated herein, significantly improves the reliability of parts made by LPBF. Our work will lead to AM of SiC of unprecedented quality/performance and application of LPBF to refractory ceramics that are difficult to sinter.
Also, this technique is very crucial for advances in the fabrication of SiC-based materials for various structural/thermal/medical applications and the semiconductor industry.

Results
Additive manufacturing of silicon carbide by LPBF. High-energy, short-pulse femtosecond fiber laser 28 (Supplementary Fig. 1) is used to fabricate dimensionally accurate SiC components from computer-aided designs ( Supplementary Fig. 2). The starting SiC powders consisted of polycrystalline 6H-SiC with particle sizes of 20-40 µm and 99% nominal purity, confirmed by xray diffraction (XRD) patterns and Raman spectroscopy (Fig. 2). To identify the appropriate processing parameter set, various laser powers and scan speeds were applied to produce 12 SiC tubes (Supplementary Table 1). The laser-sintered compounds were viewed by SEM to assess the material structure and porosity level ( Supplementary Fig. 2). As seen in Supplementary Fig. 2, the high-power femtosecond fiber laser fuses SiC powders. The structures of the top and surfaces of sintered objects appear very similar. Thus, one SEM image was selected to represent the surface structures of the others. This approach was extended to other figures throughout the paper. The AM objects had a high level of porosity in a random pattern. Buoyancy and caliber methods were applied to measure the porosity and density of laser-sintered objects. The investigated AM objects indicated porosities from 49.8% to 53.2% (Supplementary Table 2). The AM objects investigated in this study had bulk densities from 1.50 g/cm 3 to 1.61 g/cm 3 . All density measurements were performed based on the theoretical density of SiC, 3.21 g/cm 3 , which was assumed in deriving the porosity. To evaluate the effect of laser power and scanning speed, these two parameters were varied and were found to have insignificant effects on the porosity level and density of AM objects implying that the different processing parameters used in the LPBF process likely induced the same effects on the powder surface. Porosity content was ascribed to incomplete sintering in the powder layer. Fig. 2. Powder x-ray diffraction pattern of a feedstock SiC powder and b laser-sintered SiC. c Powder Raman spectroscopy of feedstock SiC and laser-sintered SiC. Four phases-6H-SiC, 3C-SiC, silicon, and carbon-were identified, as marked by symbols. The probe size of the laser was between 500 nm and 1 µm. One representative spectra are shown for laser-sintered to demonstrate characteristic peaks of 3C-SiC and 6H-SiC found separately at two different R1 and R2 regions.

(c)
Raman spectroscopy, were utilized for the phase analysis. XRD provides structural analysisinformation regarding how atoms of molecules are packed in the crystalline structure-while Raman analysis is designed to examine thin structure electronic levels and vibrational modes present in a sample. Hence, the combination of XRD and Raman provided complete information regarding the structural aspects of the samples. The analysis allowed us to identify the phase separation of 6H-SiC into silicon (Si) and carbon (C) and the subsequent nucleation of 3C-SiC and spheroidal graphite (Fig. 2). XRD measurements determined structural changes and phase separation following laser irradiation ( Fig. 2a, b). The as-received SiC powders were mainly identified as hexagonal 6H-SiC crystal structures and much smaller phase fractions of rhombohedral 15R-SiC. There was no detectable Si phase and SiO2 (Fig. 2a). In addition to 6H-SiC peaks, coinciding cubic 3C-SiC and Si diffraction peaks emerged for powders obtained from laser-sintered objects (Fig. 2b). XRD results were the first experimental evidence of phase separation during laser-material interaction.

Fig. 3. Raman mapping of additively manufactured SiC part showing phase separation induced by high-energy short-pulse laser irradiation.
The Raman scanner is capable of carrying out rapid point-topoint mapping of a the laser-irradiated particle surface and b the polished cross-section of the neck region where particles bind. For the area shown by each color, the corresponding Raman spectrum is demonstrated. Univariate images were constructed by bracketing bands of ~520, 780, and 1350 cm -1 with cursors for Si, SiC, and C, respectively. The intensity between those cursors at each data point is demonstrated in the Raman map.
Raman spectroscopy measurements were carried out on as-fabricated SiC powder and lasersintered particles (Table S1) A univariate image was rendered using green brackets to enclose the area around 520 cm -1 , red brackets around the area between 766 and 788 cm -1 , and blue brackets to enclose the area around 1350 cm -1 . The intensity between the bands selected by the cursors at each data point was calculated to construct the Raman image. Thus, the green, red, and blue areas in the Raman image predominantly correspond to Si, 6H-SiC, and C, respectively. Fig. 3a demonstrates a laserirradiated particle surface. Phase separation is clearly distinguishable on the particle surface. The green and blue areas in the Raman image are associated with strong Si and C Raman peaks, respectively. At some locations, the nucleation of the 3C-SiC polytype occurred primarily subsequent to the thermal decomposition of SiC. In Fig. 3b, Raman spectra in the red color area indicate unirradiated 6H-SiC powder with an accompanying intensity peak at 520 cm -1 that is the characteristic peak of Si. The Raman scattering efficiency of crystalline Si is about ten times greater than that of the crystalline SiC peaks 31,33 . Thus, Si content is negligible in the red area.
The blue area predominantly consists of C and small amounts of 3C-SiC and Si. The green area is indicated by two separate Raman spectra. In addition to a Si signal at around 520 cm -1 , one spectrum has the characteristic peak of 3C-SiC and the second contains the Raman spectrum of 6H-SiC. The important takeaway point from the Raman data is that following the laser-driven solid-state phase separation of 6H-SiC, the solidification process favored the reaction of Si and C to form 3C-SiC and 6H-SiC, depending on the equilibrium conditions and temperature. Further characterizations were performed via SEM, TKD, STEM, and HRTEM to provide insight into the occurrence of phase separation and the nucleation of 3C-SiC, 6H-SiC polytypes, and graphite. The microstructural features after laser sintering were investigated to assess the binding mechanism and phases at junctions using backscattered-electron (BSE) imaging and energydispersive x-ray spectroscopy (EDS) elemental distribution mapping. Supplementary Fig. 3 indicates a mirror-polished cross-section of a laser-sintered part. The presence of different phases is clearly distinguishable, particularly at the locations where particles bind. Further composition analysis using SEM-EDS maps of elemental distribution illustrates the C and Si variation across the polished cross-section of the AM object (magnified in Supplementary Fig. 3b, c). It reveals that Si is enriched at the particle-particle interface, where C is depleted. The high-intensity red area is a consequence of using a C polymer during the polishing of the AM parts. Silicon enrichment at some locations was ascribed to laser-heating-induced SiC decomposition. The elemental distribution map revealed that a Si-rich phase played a significant role in the fusion of 6H-SiC powder particles. These regions occurred thousands of times throughout the laser-sintered parts and consolidated the SiC tubes. To obtain a better understanding of the underlying mechanism that governs phase separation and fusion of SiC particles, three TEM lamellas were prepared from three different locations where two SiC particles bind together and the Si phase exists as an interface (Supplementary Fig. 3). Fig.   4 demonstrates the results of TKD mapping and corresponding STEM-EDS analysis. The images at the top, middle and bottom in Fig. 4 represent region 1, 2 and 3, respectively. The XY-plane refers to the layer that is fabricated parallel to the building direction in the laser sintering process.
Identification of cubic and hexagonal crystal structures and undetected region was carried out through TKD mapping, and spatial distribution mapping of Si and C was performed by STEM-EDS. The combination of TKD mapping and STEM-EDS enabled precise phase identification during laser-material interactions. The phase separation at the irradiated area is clearly distinguishable. Region 1 indicates that the reaction layer is composed of two crystal structure, cubic and hexagonal. The phase map exhibits the cubic phase connecting two hexagonal powders.
Moreover, relatively large pockets of hexagonal phase grains are dispersed inside the cubic phase, with grain sizes ranging from 500 nm to 1.5 µm. The TKD phase map of region 2 and 3 shows that the reaction layer consists of cubic, hexagonal nano-precipitates and some undetected areas.
To resolve the undetected region, a corresponding STEM-EDS analysis was performed in the same region where the TKD mapping was obtained. The intensity of silicon phase is relatively uniform across region 1, while carbon is depleted in some parts of the region investigated where two SiC powders are apparently joined by a cubic phase but is enriched in the relatively large pockets of hexagonal phase inside the Si interface and hexagonal powders. STEM-EDS analysis of region 2 and 3 indicated that areas undetected by TKD mapping were C phase. In the C-rich region, there is a lack of Si content. The analysis showed this region almost completely took the form of C. TKD mapping is unable to differentiate the cubic Si and 3C-SiC due to similarity in Kikuchi pattern of these phases, while hexagonal phase is identified as 6H-SiC.
To differentiate cubic phases and provide understanding of the structural order of the C phase, TEM analysis was performed in the modes of STEM, HRTEM, and bright-field TEM (BFTEM).  Our results also indicated that the 6H-SiC nano-precipitates was formed on the order of ~2-20 nm. We called these repetitive nanoscale 6H-SiC patterns "nanobreathing". The TKD phase map shown in Fig. 6a clearly demonstrates red dots (6H-SiC) in the yellow region (Si). Detailed examination of these small features was performed using HRTEM inside the cubic Si phase on the [101] zone axis (Fig. 6b, c). The TKD phase map exhibits densely populated, uniformly dispersed nanoscale 6H-SiC precipitates inside the cubic Si. The distribution of these nanoscale patterns appears homogenous across the Si phase. The white box in the TKD phase map denotes where HRTEM was performed. The 6H-SiC nanoprecipitates in the [101] zone are marked with arrows to indicate the difference in lattice spacing between Si and the incipient of 6H-SiC (Fig. 6c). The lattice spacing in the Si and 6H-SiC a-direction was 5.10 Å (RLS ~5.43 Å ) and 3.01 Å (RLS ~3.08 Å), respectively, which was near the Si and 6H-SiC a-direction lattice parameter 34 . Based on SADPs, simulated crystal models are constructed using CrystalMaker®, as overlays on the HRTEM images of silicon and 6H-SiC for better understanding (Fig. 6d, e).
Femtosecond laser irradiation resulted in the formation of well-ordered and highly oriented PGSs through a solid-solid transformation (Fig. 7e, f, g). The diameter span range was 200-600 nm. The HRTEM images show the graphitic degree of the C materials. The interplanar spacing near the periphery of the PGSs is about 3.43 Å. The fringes in each sector near the periphery are mostly parallel straight lines, exhibiting a high graphitic degree. The PGSs were found to disperse nonuniformly in the region where the phase separation of SiC took place (Fig. 7a, b). In the STEM-HAADF images, the brightest areas correspond to heavy Si atoms and the darkest areas represent light C atoms. The contrast of the SiC grains is between Si and C. The TEM-BF images indicate that the shapes, morphology, and structural order of the PGSs are quite similar (Fig. 7a). HRTEM images indicate that the peripheries of the PGS exhibit a higher degree of graphitization than the central regions (Fig.7f, g). The STEM-EDS results were highly consistent with the STEM-HAADF analysis (Fig. 7c, d). Silicon is absent, whereas C is enriched, which is sign of PGS formation.
These results confirmed the graphite synthesis whose characteristic peaks were also detected through Raman analysis. Even though pyrolytic graphite is commonly obtained through gas-solid transformations like chemical vapor deposition, this study proves that laser-induced solid-state disintegration of SiC can also be used to synthesize spheroidal pyrolytic graphite. Fig. 1 schematically illustrates the laser-induced phase separation of 6H-SiC into Si and C and subsequent formation 6H-SiC nanoprecipitates, spheroidal graphite and small pockets of 3C-and 6H-SiC.

Discussion
In this study, single-step AM of SiC via powder sintering routes is demonstrated, followed by extensive microstructural characterization. Fabrication of AM SiC was achieved without the use of sintering additives or binder elements. In addition, undesired SiO2 formation was not observed during the laser-material interactions (Fig. 2). The mechanism responsible for consolidation of SiC powders may lie in inertial confinement fusion. Very short pulse and high power laser may result in highly localized high pressure state on the process. Thus, highly volatile Si reacts with carbon rather escaping under vacuum environment due to laser confinement. Juodkazis et al.
showed formation of nano-cavities in sapphire by single, 800 nm, 150 fs, 120 nJ pulses 35 . Single laser pulse (100 nJ, 800 nm, 200 fs) produced high temperature (5 x 10 5 K) and pressure ( ~10 TPa) 35 . This study supports the possibility of high pressure state during material-laser interaction. Fig. 8. a Temperature dependence of stability diagram for many different SiC polytypes 36 . b Temperature dependence of the decomposition free energy for the reaction of 6H-SiC → Si + C at 1 atm. The decomposition reaction occurs when the free energy become positive above ~2500 K.
There was a narrow process window that was satisfactory for the fabrication of SiC tubes as designed 7 . In this narrow gap, varying AM parameter sets, such as laser powers and scanning speeds, had insignificant impacts on the properties of the AM parts. Tubes made with different parameters sets were consolidated. Density measurements revealed almost equivalent porosity levels and densities for all AM parts. This equivalence can be ascribed to sufficient laser energy and scanning speed delivered to the powder bed to bind the SiC powder particles via the disintegration of SiC. The laser is only connecting the neighboring powder particles without much changing the particle shapes and how the particle are stack together. There is very little or no displacement of particles.
To adequately reveal the nucleation mechanism of the 6H and 3C polytypes after the  (Fig. 8a) 36,37 . The high-energy short-pulse femtosecond laser fiber induced high nonequilibrium cooling conditions, which yielded a rich variety of microstructures and often preferentially selected nonequilibrium growth modes (Fig. 3) 41 . Thus, the present study results are highly consistent with the phase stability diagram of SiC (Fig. 8a) 36 . Small quantities of impurities and non-stoichiometry also had a great impact on polytype stabilization. The partial pressure of Si vapor was several times higher than that found in C 42 . The multiplicity of low-energy surfaces and the high-symmetry nature of 3C-SiC may account for its occurrence in the initial stages of growth over a broad range of temperature (1400-2000 °C). These factors could have given rise to rapid growth and easy nucleation along several directions, which led to large crystals bounded by low-energy forms 37 . While this kinetic argument can be linked to the occurrence of 3C-SiC over a large temperature range, its high symmetry presumably increased the vibrational entropy contribution to the free energy, hence making a contribution to the equilibrium stability of 3C-SiC at elevated temperatures. Besides that, the temperature-dependent free reaction energy for the decomposition reaction of 6H-SiC → Si + C has been calculated through the density functional theory (DFT), as shown in Fig. 8b. The positive free reaction energy means that the decomposition reaction is energetically favorable.
From Fig. 8b, we can see that the decomposition reaction occurs above ~2500 K. This is consistent with the previous experiments, in which the decomposition of SiC into solid C plus liquid Si begins at ~2840 K 43 . The slight difference of the DFT results with the experimental value could arise from the approximation of the anharmonic effect at high temperatures.
This study establishes a fundamental understanding of the phase separation mechanism of a complex SiC compound material during high-energy short-pulse laser-material interactions.
Extensive microstructural observation by XRD, Raman spectroscopy, SEM, TEM, and TKD revealed the decomposition and surface reconstruction of SiC. Thus, phase separation was confirmed by multiple characterization tools. It was found that femtosecond laser irradiation yielded a rich variety of microstructures and phases-thin Si and C nanomaterials, multiscale 6Hand 3C-SiC pockets, and highly ordered PGSs. The polytype 6H-SiC decomposed into Si and C, and subsequently Si(l)+C(s)→ α or β SiC(s) reactions occurred to form multiscale 6H-and 3C-SiC pockets (Fig. 4). For the first time, densely populated, uniformly dispersed nanoscale (~2-20) 6H-SiC precipitate-nanobreathing was formed inside a Si phase following the phase separation of 6H-SiC by laser irradiation. This remarkable discovery can be exploited in many different ways, To the best of author's knowledge, for the first time, highly oriented PGSs were reported during the phase separation of SiC using high-energy laser irradiation. Fig. 7 shows the HRTEM analysis of the 002 fringes of a PGS. A high degree of graphitization occurring near the periphery of graphite sphere can be deducted from the fringes, which are mostly aligned as parallel straight lines. Elemental mapping through the cross-section of the focused ion beam foil revealed the formation of a C sphere; Si was absent when C was enriched. Spheroidal graphite was produced through solid-to-solid transformation. High-energy short-pulse laser-derived graphite aggregate tended to extend in the c-direction rather than the a-direction, generating spheroidal (nodular) graphite. Graphite spheroids are widely found in spheroidal graphite cast iron 44,45 . Spheroidal graphite acts as a "crack arrester" because its rounded shapes induce fewer stress points- and 22×27 µm 2 at the powder neck region and powder surface, respectively. The Raman image was constructed using green-colored brackets enclosing 520 cm -1 , red brackets enclosing between 766 and 788 cm -1 , and blue brackets enclosing 1350 cm -1 .

SEM.
Laser-sintered components were analyzed by SEM (Tescan Mira3) to gain knowledge of the porosity level and binding mechanism of the SiC powders. The cross-sectional microstructure analysis at powder neck region was carried out using BSE imaging at an accelerating voltage of 10 kV. Elemental distribution mapping was performed using EDS analysis to determine the distribution of Si and C. A thin foil with a high-quality polished surface was prepared using an FEI Quanta focused ion beam with a low accelerating voltage of 5 kV and 2 kV at the final thinning step. TKD maps was generated using an Oxford Instruments Nordlys detector mounted on a Tescan Mira3 with an accelerating voltage of 20 kV in high current mode. TKD mapping was conducted at a working distance of about 4 mm with a tilting angle of -20° and step size of 20 nm.
TEM. Electron-transparent TEM lamella were prepared using an FEI Quanta focused ion beam at here is 3N vibrational modes. is the anharmonic free energy. In order to estimate the anharmonic free energy, we followed the approach of Wallace 47 who showed that the anharmonic part of the free energy can be written as = 2 2 . Experiments for different crystals showed that there is an empirical relation between the average Gruneisen parameter 〈 〉 and 2 , which is given per atom by 2 = 3 Θ ∞ � (0.0078〈 〉 − 0.0154) 48 . The values of Gruneisen parameter for 6H-SiC, Si, and C are 1.23 49 , and 2.28 50 , respectively. Θ ∞ is the high temperature harmonic Debye temperature defined by Θ ∞ = ℏ( 5〈 2 〉 3 � ) 1/2 � 48 . The setting of DFT calculations have been discussed elsewhere 51 .

Data availability
Supporting Information is available in the supplementary materials and more data can be obtained upon reasonable request from the corresponding author.