# Pressureless sintering of 3D printed SiC filaments at 2200 °C: impact of coarsening, debinding atmosphere, residual carbon and anisotropy

Fatima Hammoud, Jérôme Lecourt, Christelle Bilot, Charles Manière

► To cite this version:

Fatima Hammoud, Jérôme Lecourt, Christelle Bilot, Charles Manière. Pressureless sintering of 3D printed SiC filaments at 2200 °C: impact of coarsening, debinding atmosphere, residual carbon and anisotropy. Progress in Additive Manufacturing, 2025, 2025, 10.1007/s40964-025-01230-1. hal-05156043

HAL Id: hal-05156043

https://hal.science/hal-05156043v1

Submitted on 10 Jul 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
# Pressureless sintering of 3D printed SiC filaments at 2200°C: impact of coarsening, debinding atmosphere, residual carbon and anisotropy

Fatima Hammoud<sup>a*</sup>, Jérôme Lecourt<sup>a</sup>, Christelle Bilot<sup>a</sup> and Charles Manière<sup>a*</sup>

<sup>a</sup> Université de Caen Normandie, ENSICAEN, CNRS, CRISMAT, Normandie Univ, 14000 Caen, France

* Corresponding authors: [fatima.hammoud@ensicaen.fr](mailto:fatima.hammoud@ensicaen.fr), [charles.maniere@ensicaen.fr](mailto:charles.maniere@ensicaen.fr)

## Abstract

Silicon carbide (SiC) is a material of choice for high-temperature, aerospace, nuclear, and wear-resistant applications due to its exceptional mechanical strength, oxidation resistance, and chemical stability. However, full densification of SiC remains challenging because of its strong covalent bonding, which necessitates high sintering temperatures and the use of sintering additives. In this study, 3D-printed SiC parts using Nanoe® filaments were sintered at 2200 °C following debinding under either nitrogen atmosphere or vacuum. Dilatometry was employed to investigate the densification behavior. A significant variation in final density and sintering anisotropy was observed depending on the debinding conditions. The findings underscore the critical role of carbon additives in the intermediate and final stages of sintering, where they inhibit grain coarsening and enable the achievement of high densification. In this system, coarsening plays a dominant role and completely reverses the expected Master Sintering Curve (MSC) behavior when carbon content is insufficient and oxide phases are present. This study demonstrates that both a high heating rate (10 °C/min) and the presence of carbon (originating from the binder) are essential to achieve full densification in 3D-printed SiC specimens.

**Key words:** Silicon carbide; anisotropy; additive manufacturing; sintering aid; coarsening.

## 1. Introduction

Advanced ceramic materials are increasingly utilized in extreme environments where high-temperature resistance, mechanical strength, and oxidation resistance are critical. Their superior performance makes them essential in aerospace, nuclear applications, energy systems, and the automotive industry[1] [2] [3]. Among these materials, silicon carbide (SiC) has emerged as a leading candidate due to its remarkable combination of mechanical and thermal properties. SiC exhibits high hardness, excellent thermal conductivity, and outstanding corrosion resistance,
making it particularly suitable for applications such as gas turbines, heat exchangers, and
spacecraft components [4] [5] [6]. In addition to these applications, SiC and its composites are
gaining attention in nuclear technology due to their low neutron absorption cross-section and
radiation resistance. Their stability under high neutron flux conditions makes them potential
materials for advanced fission and future fusion reactors. Moreover, SiC is widely used in
protective armor systems and cutting tools due to its exceptional wear resistance[7] [8] [9].

Despite these advantages, achieving fully dense SiC components remains a significant
challenge. The intrinsic properties of SiC, such as strong covalent bonding and low self-
diffusivity, make its densification particularly difficult. Pure SiC requires sintering
temperatures above 2100°C, and conventional pressureless sintering often results in incomplete
densification unless sintering additives such as boron carbide (B₄C) or yttrium oxide (Y₂O₃) are
incorporated to enhance the densification and suppress coarsening [10] [11]. Alternative
approaches, such as spark plasma sintering (SPS) and hot pressing, have been developed to
improve densification, but these methods require expensive equipment and are not always
suitable for large-scale production, unlike pressureless sintering [12] [13]. The use of sintering
additives, along with careful monitoring of the final stage of sintering, is crucial for achieving
full densification. In particular, the presence of oxide phases is known to significantly enhance
grain-coarsening mechanisms. Additives such as carbon play a key role in removing surface
oxides, thereby limiting coarsening and enabling full densification by pressureless sintering.
Typically, carbon is added at around 3 wt%, often in combination with 0.2 to 0.5 wt% of B₄C,
to effectively promote densification [14] [15] [16].

Recent advancements in additive manufacturing have provided new opportunities for producing
complex SiC components. Among the different additive manufacturing techniques, fused
filament fabrication (FFF) has been widely explored for its ability to fabricate intricate ceramic
structures efficiently [17]. Nanoe® has developed SiC-based filaments that can be processed
using standard FFF techniques, offering a novel route for manufacturing high-performance SiC
parts. This technique enables the production of customized geometries with reduced material
waste, making it an attractive alternative to conventional shaping methods [18].

However, controlling the sintering behavior of these printed filaments has not been extensively
studied, and optimizing their processing conditions is necessary to ensure high densification
and desirable mechanical properties. A key challenge is to control the amount of carbon after
debinding in order to ensure full densification at 2200°C. The influence of the debinding
atmosphere (e.g., nitrogen, or vacuum) is of primary importance for the success of sintering.
This study aims to address these challenges by investigating the sintering behavior of Nanoe® SiC 3D printed parts under different debinding conditions, including nitrogen atmosphere and vacuum. In addition, a high temperature dilatometer will be used to precisely study the sintering behavior and the influence of debinding on sintering. By addressing these challenges, this study contributes to the advancement of SiC additive manufacturing, paving the way for its integration into high-performance industrial applications.

## 2. Experiment part

### 2.1. Sample Preparation

Zetamix porcelain filaments, specifically formulated for ceramic FFF printing, were utilized. These filaments consist of fine ceramic powder (< 1.0 µm) bound by a polyolefin-based thermoplastic binder system. The printing was performed using a Prusa i3 MK3S printer with PrusaSlicer 2.5.0 equipped with a 0.6 mm nozzle, used here for the Zetamix filaments. The printing parameters were set as follows: a printing temperature of 120°C, 50°C of bed temperature, a layer thickness of 0.2 mm, and a printing speed of 20 mm/s. The printing process adhered to Zetamix guidelines to ensure proper filament handling and extrusion without filament grinding.

After printing, the green parts underwent a two-stage debinding process to remove the thermoplastic binder. The first stage involved solvent debinding in an acetone bath at 40°C, where at least 12% of the part's weight was removed. The duration of this step depended on the part's wall thickness and geometry but required a minimum of two hours. Following solvent debinding, the parts were dried in ambient air for at least two hours, ensuring complete removal of the solvent and allowing for accurate mass measurements of the removed binder. The second stage consisted of thermal debinding, where the parts were placed in a crucible on a refractory powder bed to accommodate shrinkage. A controlled heating cycle was applied, starting from room temperature (RT) to 700°C at a rate of 10°C/h, with a dwell time of 1 hour. Subsequently, the temperature was decreased back to RT at a rate of 50°C/h, completing the debinding process in approximately 3.5 days. The thermal debinding was conducted under nitrogen to convert the remaining organics into graphite, which is mandatory for the sintering of SiC powder that already containing 0.5 wt% of B₄C. Thermal debinding under vacuum was also studied using pressureless Spark Plasma Sintering (P-SPS). Since very low heating rates cannot be applied with this device, the cycle was set to 20 °C/min up to 700 °C, with a 10 min dwell time, followed by natural cooling.
## 2.2. Sintering process

The sintering process followed the recommended Nanoe temperature profile, but was adapted to our equipment which is a high-temperature dilatometer (Linseis Dil L75 2400°C). The process started with heating from room temperature (RT) to 2200°C at 10°C/min under argon with a 1-hour dwell time. The final step involved controlled cooling from 2200°C back to RT at a maximum rate of 10°C/min. However, in this study, additional sintering conditions were explored (as with the debinding) to evaluate their influence on densification and microstructure. A master sintering curve (MSC) study was conducted to assess the sintering behavior of the powder. This study included three tests at constant heating rates of 2, 5 and 10°C/min up to 2200°C. This was followed by a 2h-holding and rapid cooling of 20°C/min to room temperature. Dilatometry measurements continuously monitored length variations, enabling the identification of critical sintering stages and the calculation of thermal expansion coefficients and sintering activation energy. The density of sintered parts was measured using Archimedes' method.

## 2.3. Characterization

The characterization of printed samples included microstructural analysis, thermogravimetric analysis, relative density measurements and hardness tests. Microstructural analysis was conducted using scanning electron microscopy (SEM Jeol 7200 LV) to investigate grain size, morphology, and porosity distribution. SEM images provided insight into grain growth mechanisms and pore evolution throughout the sintered samples. The mean grain size was determined using the linear intercept method with Mendelson's 1.56 scaling factor [19]. Additionally, thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) were performed using a TA Instruments Q500 system to study the decomposition process during thermal debinding. TGA provided insights into mass loss as a function of temperature, allowing for the evaluation of carbon conversion efficiency. The choice of the debinding atmosphere was a critical parameter, influencing binder removal and carbon conversion. The sample volumetric mass was measured by the Archimedes' method, and the Vickers hardness was measured on the device "Falcon 500G2" from "Innovatest". The phase proportions and the ratio of the eliminated phases were determined for vacuum debinding and nitrogen debinding by successive measurements of the weights. The acetone elimination ratio ($x_{acetone}$) is obtained by measuring the weight before and after chemical debinding. The same procedure is applied after thermal debinding ($x_{thermally eliminated}$) and after last debinding in air at 575°C to estimate the
proportion of remaining carbon (x<sub>carbon</sub>). The weight ratio of SiC powder (X<sub>SiC</sub>) and all other weight loss phases (X<sub>acetone</sub>, X<sub>thermally eliminated</sub>, X<sub>carbon</sub>) can be determined by the following equations:

$$
\left\{
\begin{array}{l}
X_{\text{acetone}} = x_{\text{acetone}} \\
X_{\text{thermally eliminated}} = x_{\text{thermally eliminated}}(1 - X_{\text{acetone}}) \\
X_{\text{carbon}} = x_{\text{carbon}}(1 - X_{\text{acetone}} - X_{\text{thermally eliminated}}) \\
X_{\text{SiC}} = 1 - X_{\text{acetone}} - X_{\text{thermally eliminated}} - X_{\text{carbon}}
\end{array}
\right.
\qquad (1)
$$

3. Results and discussions

3.1. SiC powder analysis

The analysis of silicon carbide (SiC) powder using SEM. As filaments have about 50vol% of binder, it was removed by vacuum debinding and the powder was obtained after grinding with an agate mortar to break soft agglomerates. The SEM granulometry was conducted by measuring the grain with the software “imageJ”. The histogram in Figure 1 indicates that most grains are between 0.1 and 0.3 µm, with a peak around 0.15 µm, while a few larger grains above 0.5 µm suggest a slight clumping. Overall the grain distribution is asymmetric but largely submicronic, which is required to favor pressureless sintering. The SEM image confirms that the grains have sharp edges, which is common for SiC powders made by grinding. The calculated average grain size is 0.25 µm.

<figure><img src="image_2.png" /><figcaption>Figure 1. SEM granulometry analysis on the debinded Zetamix® SiC filament.</figcaption></figure>

3.2. Thermal Analysis of the Debinding Process

Figure 2 presents the TGA-DSC curves for the debinding process of SiC samples sintered under N₂ atmosphere (A) with a 10°C/min heating ramp up to 700°C and a 1-hour holding time and
under vacuum in P-SPS (B). The TGA curves (black) illustrate the weight loss associated with the decomposition of the polymer binder, while the DSC curves (red) highlight the thermal events (endothermal/exothermic) occurring during each stage of this process. Both samples exhibit significant mass loss in the early stages, corresponding to the removal of the organic binder. The DSC of both tests indicates very similar profile with mostly small and negative values. This indicates that both debinding methods are slightly endothermal, which helps protect the sample from uncontrolled hot spot formation that can arise in oxygen atmosphere with strong exothermic behavior. The TGA curves are also very similar with about 20% mass loss in both cases and between 220°C and 450°C. Normally, the carbon yield is more effective in nitrogen because vacuum is generally associated with increased removal of volatile species [20]. Here the difference in carbon yield is not clear but as the P-SPS heating rate is higher than that for the nitrogen, which is in a furnace, higher carbon yield is expected in nitrogen.

<figure><img src="image_3.png" /></figure>

<figure><img src="image_4.png" /><figcaption>Figure 2. TGA-DSC analysis of the debinding process of SiC samples. (A) Under N₂ and (B) under vacuum, with a 10°C/h ramp up to 700°C and a 1-hour holding time.</figcaption></figure>
### 3.3. Master Sintering Curve Behavior

Figure 3 illustrates the shrinkage behavior and resulting microstructures of SiC sintered at different heating rates (2 °C/min, 5 °C/min, and 10 °C/min). All these samples were dedined in P-SPS configurations (vacuum). The shrinkage curves plotted in time (figure 3a) show that a higher heating rate leads to a faster densification process, which is expected. Typically, samples sintered at lower heating rates (which remain longer at intermediate temperatures) exhibit shrinkage at lower temperatures. What is very unconventional is the distribution in temperature (figure 3b). Here, the expected MSC behavior is fully inverted: the fast heating rate not only leads to quicker densification but also shifts the onset of shrinkage to a lower temperature. The other noticeable effect is that the total shrinkage is grater for fast heating and lower for lower heating rates. The microstructure confirms this tendency with less porosity for the higher heating rate. The other element arising by the microstructure aspect at low heating rates (2 and 5 °C/min) is that their aspect suggests a strong coarsening with rounded porosity interconnected network. Since coarsening reduces densification mechanisms, this may explain the disturbance observed on the shrinkage curves. Considering the fact that the MSC behavior is fully inverted, coarsening is dominant and absolutely need to be countered by conditions favorable to densification mechanisms. This confirms that increasing the heating rate can enhance densification due to the presence of strong coarsening, but it may also introduce internal stresses, which should be carefully controlled to avoid defects. The heating rate is fixed to 10°C/min for all subsequent sintering experiments. Higher temperatures than 2200 °C are avoided to limit grain growth, and the carbon content must also be increased to mitigate coarsening effects, as previously demonstrated in the literature [14].
<figure><img src="image_5.png" /><figcaption>Figure 3. Sintering shrinkage behavior of SiC at different heating rates. (A) the linear shrinkage as a function of time (B) and temperature.</figcaption></figure>

### 3.4. Effect of Debinding Atmosphere on SiC Densification

Figure 4 compares the sintering shrinkage behavior of 3D-printed SiC filaments sintered at 10 °C/min, following prior debinding under vacuum and nitrogen ($N_2$) atmospheres. The shrinkage in the build direction (Z) and the in-plane direction (R, i.e., the plane of the deposited layers) exhibits strong anisotropy, with a shrinkage difference of approximately 13% for the vacuum-debonded cycle and 9% for the nitrogen-debonded cycle. This anisotropy may arise from the porous architecture induced by 3D printing or from grain orientation effects. It is difficult to assess particle orientation from the sintered microstructure, as grain growth tends to reduce the anisotropy, and EBSD (Electron Backscatter Diffraction) analysis is challenging on nonconductive ceramics. However, some degree of particle orientation appears to be present in the SEM image of the printed and debonded specimen surface. The orientation distribution analysis is presented in Figure 5, using "OrientationJ".

The relative density curves (Figure 4) show that the debinding atmosphere has a significant influence on the sintering process. Nitrogen debinding appears to promote greater
rearrangement in the green body, resulting in a higher initial relative density. In contrast, under
vacuum debinding, the onset of shrinkage occurs at a higher temperature (~1700 °C) and
follows a more gradual trend. For the nitrogen-debound sample, an active shrinkage mechanism
is observed between 500 and 1500 °C, resulting in about 3% densification. This behavior
suggests the presence of residual carbon, which may react with SiC and assist in removing the
oxide phase, an effect known to enhance densification while limiting coarsening. Above
1500 °C, both samples show similar sintering temperatures, but the nitrogen-debound sample
exhibits greater radial shrinkage, indicating enhanced densification. In summary, nitrogen
debinding leads to improved green body rearrangement, additional low-temperature shrinkage,
and preservation of sintering mechanisms, all contributing to a denser final microstructure.

<figure><img src="image_6.png" /><figcaption>Figure 4. Sintering shrinkage and relative density curves of 3D printed SiC filaments sintered at 10 °C/min with prior debinding under vacuum and nitrogen (N₂) atmospheres, with Z referring to the building direction and R to the in-plane directions.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Figure 5. Particle orientation was analyzed using the SEM image of the printed, debound green specimen using OrientationJ with a 3-pixel boundary detection. The red circle highlights the dominant grain boundary, which is also represented by the red curve in the orientation distribution. The black arrow indicates the particle orientation.</figcaption></figure>

### 3.5. Microstructural Comparison

The SEM images compare the microstructures of SiC samples debound under nitrogen and vacuum (Figure 6). The sample debound under N₂ exhibits a more uniform microstructure with fewer large pores, indicating enhanced densification. In contrast, the sample sintered under vacuum shows a higher number of irregular and larger pores, suggesting coarsening partially countered by higher heating rate. The difference in porosity could be attributed to a higher carbon content in nitrogen which removes the oxide trace in favor of densification mechanisms. Another reason could be due to the crystallinity of graphite. As the TGA suggests similar carbon yield after debinding, it is also known that vacuum favor higher graphite phase crystallinity while nitrogen can favor amorphous or turbostratic carbon which has high defect but higher reactivity [20]. It is possible that lower carbon crystallinity enhances the cleaning effect of carbon on the oxide phase which favors densification over coarsening. In order to verify this aspect, we completed this study by an EDS analysis comparing the nitrogen-debound microstructure with the highly porous microstructures obtained under vacuum dedinding.
<figure><img src="image_8.png" /><figcaption>(B)</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>Figure 6. SEM micrographs of sintered SiC samples (at 10 °C/min) under different atmospheres: (A) under N₂ and (B) under vacuum.</figcaption></figure>

Figure 7 presents the EDS mapping and elemental quantification of SiC samples sintered with debinding under N₂ and vacuum. The atomic composition tables indicate that both samples contain silicon (Si) and carbon (C) as the main elements, confirming the SiC structure. However, significant differences in oxygen (O) content are observed. The vacuum-debound sample shows a higher oxygen percentage, suggesting inefficient oxide removal by carbon. In contrast, the sample debound under N₂ contains significantly less oxygen (0.7 at.% nitrogen vs. 6.7 at.% for vacuum), indicating that the nitrogen atmosphere is more efficient at removing the oxide phases. The EDS mapping images further highlight these trends, with oxygen highly concentrated in the large pores where coarsening seems to be highly active. For the nitrogen debound microstructure, the oxygen is limited to a few residual phases near some pores.

To more accurately determine the residual carbon content compared to EDS analysis, successive chemical and thermal debinding steps were performed, as described in the characterization section. The results are presented in Figure 8, showing the weight ratio of each phase under different debinding conditions. In the nitrogen case, 78.7 wt% corresponds to the SiC powder (ceramic phase). Acetone debinding removes 4.88 wt% of soluble organic components, while thermal debinding in nitrogen eliminates 15.6 wt% of volatile organics elements. The remaining residual carbon accounts for 0.76 wt% of the initial composition. In contrast, vacuum thermal debinding removes a greater amount of volatile matter, with a total mass loss of 16.3 wt%. As a result, the residual carbon (needed to reduce the oxide phase) is lowered to 0.54 wt%. The filament exhibits a grey appearance compared to the black color observed after nitrogen debinding. The uniform black color indicates a good distribution of residual carbon, whereas the grey, vacuum-debound filament suggests heterogeneous carbon retention, consistent with the EDS map.
This observation aligns with the mechanism described by Clegg [21], who showed that free carbon in boron-doped SiC reacts with surface SiO₂ at a lower temperature than the intrinsic reaction between SiC and SiO₂. Without this carbon, the SiC–SiO₂ reaction occurs at a higher temperature, which can fracture interparticle contacts and promote coarsening. In our nitrogen-debound samples, the presence of residual carbon appears to prevent this reaction by removing surface oxides earlier in the process, leading to lower oxygen content and more effective densification. In contrast, insufficient carbon retention in vacuum-debound samples likely delays oxide removal, resulting in higher oxygen levels and a coarser, more porous microstructure. These results confirm that debinding under nitrogen leads to an efficient oxides removal and a denser microstructure.

<figure><img src="image_10.png" /><figcaption>Figure 7. EDS analysis of sintered SiC samples debound under different atmospheres. (A) N₂ and (B) vacuum. The top tables show atomic composition, the middle images present SEM micrographs, and the bottom images display EDS elemental mapping of Si, C, and O.</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>Figure 8. Mass loss during the different debinding stages: proportions for nitrogen (A) and vacuum (B) conditions, and (C) a detailed comparison of thermal debinding losses and residual carbon at this stage.</figcaption></figure>

### 3.6. Hardness behavior

The hardness of 3D-printed SiC in this study shows a clear dependence on the debinding atmosphere. Using the Vickers HV1 method (10 s dwell time), samples debound in nitrogen reached an average hardness of 2,646.88 HV1, while those debound in vacuum averaged only 1,414.43 HV1.

This difference highlights the role of residual carbon: nitrogen debinding preserves some carbon, which promotes densification. In contrast, vacuum debinding removes more carbon, limiting densification and resulting in lower hardness.

In the literature, pressureless sintered SiC typically shows hardness between 2,000–2,400 HV1, and hot-pressed or HIPed SiC can reach up to 2,800–3,000 HV1 [22] [23] [24]. The nitrogen-debound SiC in this work is in the upper range of these values, demonstrating that additive manufacturing (FFF) can produce competitive hardness when optimized. The vacuum-debound parts, however, fall well below this range, due to poor densification.

These results confirm the importance of controlling the debinding atmosphere in order to achieve high-performance 3D-printed SiC components.
### 3.7. Sintering mechanisms analysis

In order to capture the difference of the debinding conditions on the sintering kinetics, the sintering behavior has been modeled. As the master sintering curve is inverted and cannot be exploited, the model where identified by direct regression approach (2) using modified Skorohod's moduli [25] and the continuum theory of sintering [26].

$$Y = \ln\left(\frac{-3(1 - \theta)^3}{T\theta\psi}\right) = \ln\left(\frac{r\eta_0}{\alpha}\right) + \frac{Q}{RT} \quad (2)$$

Where $\Theta$ is the porosity, $r$ the initial particle radius, $T$ the absolute temperature, $\alpha$ the surface energy, $\eta_0$ the viscosity pre-exponential constant, $\psi$ the bulk modulus, and $Q$ the activation energy of sintering.

The regression Y(1/RT) and modeled porosity curves are reported in figure 9 for the prior vacuum (A) and nitrogen (B) debinded samples. The activation energy is slightly higher for the vacuum debinding behavior with 495 kJ/mol and 489 kJ/mol for the nitrogen debinding that has the best properties. This value strongly depends on the additives but 643 kJ/mol were obtained for boron/carbon additive [27] and about 500 kJ/mol between 0.75 and 0.85 relative density in ref [28] with similar additives. For the other parameters, the bulk modulus uses modified Skorohod equation with an initial critical porosity fixed to 0.6 and a final critical porosity corresponding to the final relative density as we do not model grain growth in this model. Then, if we take 3 J/m² for the surface energy for SiC [29] and the measured initial radius (see Figure 1), the viscosity pre-exponent factor is then $\eta_0=8.18E-3$ Pas/K for the nitrogen case. This model reflects the slightly higher densification rate for nitrogen debinded, but for the thermally activated part of the model at the intermediate stage sintering is very similar.

<figure><img src="image_12.png" /><figcaption>Figure 9. Sintering model assessment and verification for the vacuum debound case (A) and nitrogen debinding (B).</figcaption></figure>
4. Conclusion

This study demonstrates the critical impact of the debinding atmosphere on the densification,
microstructure, and hardness of 3D-printed silicon carbide (SiC) parts sintered at 2200 °C.
Using fused filament fabrication (FFF) and Zetamix® SiC filaments, we show that nitrogen
debinding significantly enhances densification compared to vacuum debinding. The presence
of residual carbon in the nitrogen atmosphere favors oxide removal and suppresses grain
coarsening, leading to a denser microstructure and higher hardness.

Hardness measurements confirm this trend, with nitrogen-debound samples achieving an
average of 2,646 HV1, positioning them within the upper range of pressureless sintered and
hot-pressed SiC found in the literature. In contrast, vacuum-debound parts show a much lower
hardness (~1,414 HV1) and a porous microstructure, highlighting incomplete densification.

Moreover, modeling and dilatometry reveal that coarsening dominates the sintering behavior,
reversing the expected Master Sintering Curve (MSC) trend. A high heating rate (10 °C/min)
and sufficient carbon retention are shown to be essential to promote densification while
countering coarsening.

Overall, this work provides new insights into optimizing additive manufacturing of SiC,
particularly through careful control of the debinding process, paving the way toward for high-
performance, pressurelessly sintered ceramic components.

5. References

[1] M. Belmonte, "Advanced Ceramic Materials for High Temperature Applications," *Adv. Eng. Mater.*, vol. 8, no. 8, pp. 693-703, 2006, doi: 10.1002/adem.200500269.

[2] "Review of additive manufacturing methods for high-performance ceramic materials | The International Journal of Advanced Manufacturing Technology." Accessed: Feb. 18, 2025. [Online]. Available: https://link.springer.com/article/10.1007/s00170-019-03669-3

[3] T. Ayode Otitoju, P. Ugochukwu Okoye, G. Chen, Y. Li, M. Onyeka Okoye, and S. Li, "Advanced ceramic components: Materials, fabrication, and applications," *J. Ind. Eng. Chem.*, vol. 85, pp. 34-65, May 2020, doi: 10.1016/j.jiec.2020.02.002.

[4] T.-T. Xu, S. Cheng, L.-Z. Jin, K. Zhang, and T. Zeng, "High-temperature flexural strength of SiC ceramics prepared by additive manufacturing," *Int. J. Appl. Ceram. Technol.*, vol. 17, no. 2, pp. 438-448, 2020, doi: 10.1111/ijac.13454.

[5] H. Masuda, Y. Ohta, and M. Kitayama, "Additive Manufacturing of SiC Ceramics with Complicated Shapes Using the FDM Type 3D-Printer," *J. Mater. Sci. Chem. Eng.*, vol. 07, no. 02, Art. no. 02, Feb. 2019, doi: 10.4236/msce.2019.72001.

[6] A. Bratten, R. Chen, J. Rittenhouse, M. Leu, and H. Wen, "Improved additive manufacturing of silicon carbide parts via pressureless electric field-assisted sintering," *Int. J. Appl. Ceram. Technol.*, vol. 19, no. 5, pp. 2480-2488, 2022, doi: 10.1111/ijac.14105.

[7] Y. Katoh and L. L. Snead, "Silicon carbide and its composites for nuclear applications – Historical overview," *J. Nucl. Mater.*, vol. 526, p. 151849, Dec. 2019, doi: 10.1016/j.jnucmat.2019.151849.

[8] T. Hinoki et al., "Silicon Carbide and Silicon Carbide Composites for Fusion Reactor Application," *Mater. Trans.*, vol. 54, no. 4, pp. 472-476, 2013, doi: 10.2320/matertrans.MG201206.
[9] H. Tanaka, "Silicon carbide powder and sintered materials," *J. Ceram. Soc. Jpn.*, vol. 119, no. 1387, pp. 218-233, 2011, doi: 10.2109/jcersj2.119.218.

[10] K. Raju and D.-H. Yoon, "Sintering additives for SiC based on the reactivity: A review," *Ceram. Int.*, vol. 42, no. 16, pp. 17947-17962, Dec. 2016, doi: 10.1016/j.ceramint.2016.09.022.

[11] Z. H. Huang, D. C. Jia, Y. Zhou, and Y. G. Liu, "A new sintering additive for silicon carbide ceramic," *Ceram. Int.*, vol. 29, no. 1, pp. 13-17, Jan. 2003, doi: 10.1016/S0272-8842(02)00082-2.

[12] Z. Aygüzer Yaşar, V. A. DeLucca, and R. A. Haber, "Effect of boron carbide additive and sintering temperature - Dwelling time on silicon carbide properties," *Ceram. Int.*, vol. 47, no. 5, pp. 7177-7182, Mar. 2021, doi: 10.1016/j.ceramint.2020.11.071.

[13] M. Omori and H. Takei, "Pressureless Sintering of SiC," *J. Am. Ceram. Soc.*, vol. 65, no. 6, pp. c92-c92, 1982, doi: 10.1111/j.1151-2916.1982.tb10460.x.

[14] L. Stobierski and A. Gubernat, "Sintering of silicon carbide. Effect of carbon," *Ceram. Int.*, vol. 29, no. 3, pp. 287-292, Jan. 2003, doi: 10.1016/S0272-8842(02)00117-7.

[15] Z. A. Yaşar and R. A. Haber, "Effect of Carbon Addition and Mixture Method on the Microstructure and Mechanical Properties of Silicon Carbide," *Materials*, vol. 13, no. 17, Art. no. 17, Jan. 2020, doi: 10.3390/ma13173768.

[16] L. Stobierski and A. Gubernat, "Sintering of silicon carbide II. Effect of boron," *Ceram. Int.*, vol. 29, no. 4, pp. 355-361, Jan. 2003, doi: 10.1016/S0272-8842(02)00144-X.

[17] "Silicon carbide filaments: Microstructure | Journal of Materials Science." Accessed: Feb. 18, 2025. [Online]. Available: https://link.springer.com/article/10.1007/BF01112277

[18] L. Novakova-Marcincinova, J. Novak-Marcincin, J. Barna, and J. Torok, "Special materials used in FDM rapid prototyping technology application," in *2012 IEEE 16th International Conference on Intelligent Engineering Systems (INES)*, Jun. 2012, pp. 73-76. doi: 10.1109/INES.2012.6249805.

[19] M. I. Mendelson, "Average Grain Size in Polycrystalline Ceramics," *J. Am. Ceram. Soc.*, vol. 52, no. 8, pp. 443-446, 1969, doi: 10.1111/j.1151-2916.1969.tb11975.x.

[20] B. P. Chang, A. Rodriguez-Uribe, A. K. Mohanty, and M. Misra, "A comprehensive review of renewable and sustainable biosourced carbon through pyrolysis in biocomposites uses: Current development and future opportunity," *Renew. Sustain. Energy Rev.*, vol. 152, p. 111666, Dec. 2021, doi: 10.1016/j.rser.2021.111666.

[21] W. J. Clegg, "Role of Carbon in the Sintering of Boron-Doped Silicon Carbide," *J. Am. Ceram. Soc.*, vol. 83, no. 5, pp. 1039-1043, 2000, doi: 10.1111/j.1151-2916.2000.tb01327.x.

[22] W. E. Lee and M. Rainforth, *Ceramic Microstructures: Property control by processing*. Springer Science & Business Media, 1994.

[23] M. Caccia and J. Narciso, "SiC Manufacture Via Reactive Infiltration," in *Processing and Properties of Advanced Ceramics and Composites VI*, John Wiley & Sons, Ltd, 2014, pp. 15-25. doi: 10.1002/9781118995433.ch2.

[24] A. Can, M. Herrmann, D. S. McLachlan, I. Sigalas, and J. Adler, "Densification of liquid phase sintered silicon carbide," *J. Eur. Ceram. Soc.*, vol. 26, no. 9, pp. 1707-1713, Jan. 2006, doi: 10.1016/j.jeurceramsoc.2005.03.253.

[25] C. Manière, J. S. Diatta, T. Grippi, C. Bilot, G. Riquet, and S. Marinel, "A robust thermomechanical sintering simulation for 3D printed parts with internal lattices," *Acta Mater.*, vol. 274, p. 120024, Aug. 2024, doi: 10.1016/j.actamat.2024.120024.

[26] E. A. Olevsky, "Theory of sintering: from discrete to continuum," *Mater. Sci. Eng. R Rep.*, vol. 23, no. 2, pp. 41-100, Jun. 1998, doi: 10.1016/S0927-796X(98)00009-6.

[27] D. A. Ray, S. Kaur, R. A. Cutler, and D. K. Shetty, "Effect of Additives on the Activation Energy for Sintering of Silicon Carbide," *J. Am. Ceram. Soc.*, vol. 91, no. 4, pp. 1135-1140, 2008, doi: 10.1111/j.1551-2916.2008.02271.x.

[28] D. C. Jana, G. Sundararajan, and K. Chattopadhyay, "Effective Activation Energy for the Solid-State Sintering of Silicon Carbide Ceramics," *Metall. Mater. Trans. A*, vol. 49, no. 11, pp. 5599-5606, Nov. 2018, doi: 10.1007/s11661-018-4884-9.

[29] "Sintering Theory and Practice | Wiley," Wiley.com. Accessed: Apr. 25, 2025. [Online]. Available: https://www.wiley.com/en-us/Sintering+Theory+and+Practice-p-9780471057864