# Sacrificial ceramics powders mix for titanium alloys complex shapes production by hybridization of SPS and 3D printing

Joseph Le Cloarec, Sylvain Marinel, Claude Estournès, Cendrine Folton, Moukrane Dehmas, Charles Manière

► To cite this version:

Joseph Le Cloarec, Sylvain Marinel, Claude Estournès, Cendrine Folton, Moukrane Dehmas, et al. Sacrificial ceramics powders mix for titanium alloys complex shapes production by hybridization of SPS and 3D printing. Powder Technology, 2024, 448, pp.120324. 10.1016/j.powtec.2024.120324. hal-04738249

HAL Id: hal-04738249

https://cnrs.hal.science/hal-04738249v1

Submitted on 15 Oct 2024

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# Sacrificial ceramics powders mix for titanium alloys complex shapes production by hybridization of SPS and 3D printing

Joseph Le Cloarec¹,²*, Sylvain Marinel¹, Claude Estournès², Cendrine Folton¹, Moukrane Dehmas², Charles Manière¹*

1. UNICAEN, CRISMAT, Normandie Univ, ENSICAEN, CNRS, 14000, Caen, France

2. CIRIMAT, Université de Toulouse, CNRS, INPT, UPS, 31062 Toulouse, France

E-mail address: joseph.le-cloarec@ensicaen.fr; charles.maniere@ensicaen.fr

## Keywords

Spark Plasma Sintering; Complex Shape; 3D printing; sintering simulation; Ti-6Al-4V;

## Abstract:

The production of complex shapes via Spark Plasma Sintering (SPS) is challenging due to intricate die configurations. The "DEFORMINT" approach combines 3D printing with SPS and sacrificial powder to fabricate fully dense shapes. Using the same powder for both the component and sacrificial part avoids distortions, but high costs of titanium alloys like Ti-6Al-4V are limiting. This study tested a low-cost ceramic powder that mimics Ti-6Al-4V's behavior, forming a pseudo-isostatic configuration above 1250 K. SPS experiments established a densification model and predicted distortions in a thick cone and a thin, complex turbine blade, showing good correlation with experimental results. The new sacrificial mix does not imply additional deformation and allows the fabrication of turbine blades. The only defect is a 600 µm skin effect at the contact interface.

## Graphical Abstract:

<figure><img src="image_2.png" /></figure>
$\theta$ Porosity

$\dot{\theta}$ Porosity elimination rate (s⁻¹)

$\underline{\sigma}$ Stress tensor (N.m⁻²)

$\sigma_{eq}$ Equivalent stress (N.m⁻²)

$\dot{\underline{\epsilon}}$ Strain rate tensor (s⁻¹)

$n$ Creep law stress exponent

$A_0$ Deformability pre-exponential factor (Ks⁻¹Pa⁻ⁿ)

$Q$ Deformability activation energy (J.mol⁻¹)

$T$ Temperature (K)

$\varphi$ Shear modulus

$\psi$ Bulk modulus

$\dot{\epsilon}_z$ Axial strain rate tensor component (s⁻¹)

$\sigma_z$ Axial stress tensor component (N.m⁻²)

$\theta_{cf}$ Critical final porosity

SPS Spark Plasma Sintering

Ti64 Contraction of Ti-6Al-4V

DEFORMINT refers to the DEFORMable INTERface method
# 1. Introduction

Spark Plasma Sintering (SPS) enables the fabrication of titanium alloys with advanced properties that cannot be achieved through traditional casting methods [1–5]. However, the manufacture of complex shape is very challenging [6–8] because it requires multiple punches with complex shapes [9], making the process difficult to manage and prone to brittleness. Additionally, the use of complex-shaped punches can lead to densification gradients [10] unless sacrificial powders are employed [11]. Like for powder metallurgy route [12], the use of differential punches motions could facilitate the full densification of complex shapes, but SPS devices are mostly simple double effect pressing configurations [13–16] without the possibility to control multiple punch motions. In this context, a new method, called “DEFORMINT”, emerges to allow the fabrication of complex shapes by SPS using a multiple powder assembly with a deformable interface [17]. This method implies sacrificial powders zones, which follow the sintering of the complex shape while ensuring a homogeneous stress distribution and a highly predictable complex shape evolution. The key point is the selection of the powder materials for the interface and sacrificial, with the right chemical, physical and mechanical properties. The interface, thin and chemically inert, should follow the densification of the complex shape and the sacrificial powder without disturbing mechanically their densification and to facilitate the post SPS ejection of the desired part from the sacrificial material. This interface should not sinter and react chemically with the powders; graphite or refractory ceramics like boron nitride are generally used. Generally, the interface is 3D printed [18] to facilitate the preparation and to benefit from the high shape resolution of 3D printing methods like the stereolithography [19–23]. In the case of the sacrificial powders, it should mimic as close as possible the sintering of the complex shape powder to prevent distortions of the part with regard to its desired geometry. The sacrificial powder should also have a sintering shrinkage equal or higher to the complex shape powder with similar temperature range to ensure the full densification of the complex part. The more effective solution to ensure minimal
distortion of the complex part is to use the same powder for the part and for sacrificial zones [24,25]. Consequently, this reduces the economic interest of the method, particularly when the complex part is made of high cost material. To remedy to that lock, the identification of low-cost composite powders represents a key aspect of the DEFORMINT method. Sacrificial materials are typically used in metallurgy as casting molds [26,27] or granules for pressure transmitting medium in quasi-isostatic configurations [28,29]. Here the challenge is to identify a powder mixture that mimics the sintering of the titanium alloy used for the complex shape. Alumina powder has already been used [30]. Despite very different SPS behavior, complex shapes were produced because the initial sintering shrinkage and the lower initial powder tap density allow enough shrinkage at high temperature. However, the high sintering temperatures difference can generate distortions and a powder with similar sintering temperature is desirable. In a previous study, a ceramic-metal powder mixture has been identified to reproduce a close sintering condition of Ti-6Al-4V powders[31]. However, to ease the preparation of the SPS assembly and improve the ejection of the complex shapes, a brittle sacrificial made of low-cost ceramic is more desirable.

A low-cost sacrificial ceramic powder mixture has been identified to replace the expensive use of Ti-6Al-4V as a sacrificial material. This paper investigates the sintering behavior of this mixture and its effect on the densification, microstructure and properties of the Ti-6Al-4V produced part. An important part consists of the identification of the sintering model parameters of this sacrificial mixture, which has a complex SPS behavior with swelling at high temperatures. A numerical tool capable of predicting the complex shape deformation during SPS is developed.
## 2. Experiment and method

### 2.1. Sintering behavior identification of sacrificial ceramic mixture with the swelling phenomenon.

A finite element model is used to simulate mechanically the powder assembly during the SPS. This model includes 3 materials sintered simultaneously, the Ti-6Al-4V powder of the complex shapes (called “shape powder”), the sacrificial powder and the thin interface powder between. This simulation is based on the continuum theory of sintering [32,33] described by the following equation, which includes the swelling pressure ($P_{swelling}$) active at the very end of sintering and which opposes the sintering generating negative densification (in the sacrificial zone) and pseudo-isostatic overpressure (on the complex shape).

$$ \underline{\sigma} = \frac{1}{A\sigma_{eq}^{n-1}} \left( \varphi \dot{\underline{\epsilon}} + \left( \psi - \frac{1}{3}\varphi \right) \dot{\underline{\epsilon}} \right) - P_{swelling} \underline{\dot{\epsilon}} \quad (1) $$

For SPS, the analytical equation can be formulated.

$$ \dot{\epsilon}_z = \frac{\dot{\theta}}{(1-\theta)} = A_0 \frac{\exp\left(\frac{-Q}{RT}\right)}{T} \left(\psi + \frac{2}{3}\varphi\right)^{-\frac{n-1}{2}} (1-\theta)^{\frac{1-n}{2}} (\sigma_z + P_{swelling})^n \quad (2) $$

In the latter, $\sigma_z$ is an applied compressive stress with a negative sign and Skorohod's [33] theoretical moduli can be used:

$$ \psi = \left(\frac{2}{3}\right) \frac{(1-\theta)^3}{\theta} \quad (3) $$

$$ \varphi = (1 - \theta)^2 \quad (4) $$

Sintering behavior is assessed in three main stages.

The first stage is to identify the stress exponent ($n$). This is achieved using the “stepwise” method [34,35] in the open porosity zone to avoid disturbances of the swelling phenomenon that take place at the very end of sintering. An isothermal SPS tests were performed at 1100 K to have enough sintering amplitude far from the swelling zone. During the isothermal holding, two pressure “jumps” were applied from 28 MPa to 56 MPa and from 56 MPa to 85 MPa. The porosity elimination rate before and after the pressure “jump” are used to determine the stress exponent with the following equations.
$$n = \frac{\ln(\frac{\dot{\theta}_1}{\dot{\theta}_2})}{\ln(\frac{\sigma_1}{\sigma_2})} \qquad (5).$$

The second stage consists of the identification of the temperature-dependent parameters ($A_0$, $Q$) still in the open porosity zone to avoid the swelling. The following linear regression is used with a heating ramp SPS test [36]. The SPS test was conducted with a heating rate of 50 K/min up to 1300 K. During the test, a pressure of 50 MPa was applied before the start of heating and remained constant throughout the process.

$$Y = \ln \left( \frac{T|\dot{\theta}|}{(\psi + \frac{2}{3}\varphi)^{\frac{-n-1}{2}} (1-\theta)^{\frac{3-n}{2}} |\sigma_z|^n} \right) = \ln(A_0) - \frac{Q}{RT} \qquad (6)$$

At this stage, the modeling parameters can simulate a typical densification curve that is well suited in the initial and intermediate sintering stage. To model the final stage swelling that reduces sintering and even causes negative densification, the swelling pressure ($P_{swelling}$) must be determined experimentally. This task (stage 3) is possible once all other parameters are determined by isolating the $P_{swelling}$ term in equation (2). This results in the following equation, in which all terms of the second hand are determined or calculated from the sintering curve.

$$P_{swelling} = \left( \frac{\dot{\theta}}{A(T) (\psi + \frac{2}{3}\varphi)^{\frac{-n-1}{2}} (1-\theta)^{\frac{3-n}{2}}} \right)^{\frac{1}{n}} \operatorname{sign}(\dot{\theta}) - \sigma_z \qquad (7)$$

$$\text{With } A(T) = A_0 \frac{\exp(-\frac{Q}{RT})}{T}.$$

This equation gives the experimental evolution of the swelling pressure that can be incorporated in equation (2) to model analytically the sintering curves and in equation (1) to simulate this behavior by the finite element method. A similar approach is used for conventional sintering to simulate the swelling of metal particles [37] and the liquid phase sintering of 3D printed ceramics [38].

GNU Octave 5.1.0 is used for the sintering analytic modeling, which solves ODE (Ordinary Differential Equations) and COMSOL multiphysics 5.0 for the sintering simulations.
## 2.2. Complex shape production with the DEFORMINT method combining 3D printing and SPS

The method used to produce complex-shaped parts is the DEFORMINT method combined with 3D printing and SPS. This method has been detailed in previous work [39]. Figure 1.a displays the main steps of the method. Initially, an interface is printed using stereolithography with a mixture of polymers and alumina. This interface has the hollow shape of the desired part. Subsequently, this interface is positioned in the SPS die. Its interior is filled with shape powder (Ti-6Al-4V), while its exterior is filled with sacrificial powder. The whole assembly forms a cylinder that can be easily sintered by SPS. Under the effect of temperature, the interface undergoes pyrolysis and degrades partially into graphite [18]. The interface then follows the shrinkage of the assembly during sintering, ensuring separation between the two powder beds. Finally, once sintering is complete, the partially degraded interface allows for easy ejection of the working part from the sacrificial part. For proper application of the method, both powder beds must sinter simultaneously. Otherwise, the stress field will be inhomogeneous during the process, which may result in deformation of the final part. In addition, the total shrinkage of the sacrificial powder must be higher or equal to the shrinkage of the shape powder to ensure full densification of the latter.

In this study, the shape powder is Ti-6Al-4V and the sacrificial powder is a mixture of different ceramic powders that mimics the sintering behavior of Ti-6Al-4V under SPS conditions. The chemical composition of this sacrificial powder is confidential because in patent process but its sintering behavior is studied and simulated.

Parts of two distinct geometries are produced using this method and their sintering is simulated. The first one is a simple cone in order to test the sacrificial powder on a thick shape. Figure 1.b shows the cone interface (on the left) and the geometric configuration of the simulation, which corresponds to the experimental setup (on the right), with the Ti-6Al-4V powder inside the
conical interface and the sacrificial powder outside. In order to save computing time, only a representative quarter of the whole assembly (green part) is simulated thanks to the shape symmetry. The second shape produced is more complex and challenging, a turbine blade is selected and is used to study a thin and highly complex shape. Figure 1.c shows the blade interface and the simulation configuration. The turbine blade shape has no symmetry; therefore, the whole assembly is simulated.

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

<figure><img src="image_4.png" /></figure>

<figure><img src="image_5.png" /><figcaption>Figure 1: a. DEFORMINT method combined with 3D printing and SPS to produce complex shape parts (cone example). b. Interface of cone part with the simulated corresponding part of the assembly in green and the experimental and simulation geometrical configuration. c. Interface of the blade assembly and the experimental and simulation geometrical configuration.</figcaption></figure>
The thermomechanical loading undergone by the cone assembly is shown in Figure 2.a. To facilitate progressive degradation of the interface, the temperature is increased at a rate of 25 K/min until it reaches 900 K. The temperature is then increased at a heating rate of 50 K/min, until it reaches 1350 K for the sintering part. The applied stress remains at a minimum of 7 MPa until 900 K, it is then increased to 50 MPa within one minute for reaching this temperature. This pressure is kept constant until the final temperature is reached. Once the temperature of 1350 K has been reached, the assembly is left to cool freely under a minimum pressure of 7 MPa.

<figure><img src="image_6.png" /><figcaption>Figure 2: Thermomechanical loading applied for cone (a) and blade (b) assembly.</figcaption></figure>

Figure 2.b shows the thermomechanical loading undergone by the blade assembly. It is very similar to that of the cone assembly, with one notable difference. The blade assembly does not reach 1350K. Instead, it is held at a temperature of 1150K for a period of one hour. The decision to reduce the maximum cycle temperature for the blade assembly is based on the fact that the blade is thinner and more fragile than the cone. At temperatures above 1250K, the sacrificial powder swells. This swelling causes significant stresses on the turbine blade, leading to breakage. By reducing the maximum cycle temperature, we prevent this problematic swelling from occurring. Furthermore, as illustrated in a previous work [31], by reaching too high temperature, significant thermal contraction differences between the shape material and the sacrificial material can occur during cooling. These differences in thermal contraction induce
significant residual stresses within the blade. As a result, these residual stresses increase the
risk of blade failure during cooling or ejection of the shape.

To allow the Ti-6Al-4V material to reach its maximum relative density, temperature and pressure are maintained for one hour. In addition, the assembly is cooled without the application of pressure to minimize the impact of thermal contraction deviations and residual stresses.

## 2.3. Instrumentation

The SPS apparatus used is an "FCT HPD25" equipped with graphite grade Mersen 2333 tooling (die, puncher, and spacer). The grade 5 powder of Ti-6Al-4V used is supplied by AP&C. The particle size distribution of this powder ranges from 10 µm to 100 µm, centered around 35 µm. During the SPS experiments, the temperature is measured using a K-type thermocouple inserted into the die, 3 mm from the outer surface. A graphite paper coated with boron nitride is inserted between the puncher and the sample to ensure electrical insulation of the sample during the process.

Interfaces were produced in "3DMix Al-E02" alumina resin with a ceramic stereolithography printer, ceramaker 100 from 3DCERAM SINTO®.

An EinScan- SP 3D scanner with an accuracy of 30 µm was used to scan the Ti-6Al-4V parts.

In order to reveal microstructure, complex shapes were etched with a solution of chloridric acid, fluoric acid and nitric acid: 30 % HCl; 10 % HF; 10 % HNO₃; 50 % H₂O.

A "CLEMEX CMT.HD" apparatus was used to measure hardness under a load of 0.5 KgF.

# 3. Results and discussions

## 3.1. Identification of the sacrificial powder mix stress exponent

In Figure 3 are reported the true strain and applied pressure of the ceramic sacrificial powder during the “pressure step” test. The porosity elimination rate has been calculated with the red parts reported on the true strain curve. Considering the values of stress, porosity rate elimination and equation (5), the stress exponent “n” has been identified. The “n” average value was found
to be 1.36. A “n” value between 1 and 2 is common in SPS [35,40,41]. A dominant diffusional mechanism is indicated by a “n” value of 1 while a “n” value of 2 indicates a dominant grain boundaries sliding [42]. It seems that the sacrificial powder dominant sintering mechanism is diffusional but a complex combination of the two mechanisms cannot be excluded.

<figure><img src="image_7.png" /><figcaption>Figure 3: True strain and applied stress of ceramic sacrificial powder during the “pressure step” SPS test.</figcaption></figure>

## 3.2. Identification of the sacrificial powder mix sintering parameters

The temperature and relative density of the sacrificial powder over time during SPS test are reported in Figure 4.a. Above 1250 K, the sacrificial part begins to swell reducing its relative density. This swelling is due to a chemical reaction between the compounds of the sacrificial powder. This reaction involves a reduction that releases gas forming porosity inside the sacrificial part when the porosity is closed. The linear regression based on equation (6) is reported in Figure 4.b. This linear regression yields a parameter $A_0$ of 61023 K.s⁻¹.Pa⁻ⁿ and an activation energy Q of 325 kJ.mol⁻¹. This regression is aimed at determining the temperature dependent parameters and is then conducted in the intermediate stage sintering. At the final stage, the inner pressure exerted by the swelling ($P_{swelling}$) as a function of temperature is shown
in Figure 4.c. The $P_{swelling}$ is calculated considering equation (7). A sigmoidal function was refined to fit the experimental data, its equation is reported in Figure 4.c.

<figure><img src="image_8.png" /><figcaption>Figure 4: a. Relative density of the sacrificial powder and temperature over time during SPS test. b. Linear regression of the sacrificial powder SPS experimental data. c. Pressure exerted by the swelling of the sacrificial powder over temperature and sigmoid fit of the $P_{swelling}$.</figcaption></figure>
The analytic model is reported in Figure 5. The basic model (without swelling) fits the experimental data until 1200 K. Above 1200 K, sintering kinetics is slowed by the pressure exerted by the swelling in sacrificial powder. The basic model does not predict this slowdown in sintering kinetics and the swelling that occurs at 1250 K. Considering this $P_{swelling}$ function as presented in Figure 4.c, the swelling model can predict those phenomena and fits the experimental data at final stage sintering. The comparison between the experimental sintering curve and the calculated sintering confirms the efficiency of the constructed model.

<figure><img src="image_9.png" /><figcaption>Figure 5: Analytic model of sintering of the ceramic sacrificial powder.</figcaption></figure>
### 3.3. Optimization of the sintering cycle for the cone and turbine blade.

The simulated average relative density curves of the sacrificial powder and the Ti-6Al-4V zones for the cone case and the turbine blade are reported in Figure 6.a and Figure 6.b, respectively. As neither of the Ti-6Al-4V nor sacrificial powder sinter below 1000 K, the simulation only displays the sintering above 1000 K. The cone assembly is heated from room temperature to 1350K with a heating rate of 50 K.min⁻¹. The match between the sintering of the sacrificial powder and the sintering of the Ti-6Al-4V powder is not perfect, but these curves confirm a similar sintering temperature and shrinkage magnitude between 1000 K and 1250 K. This small disparity in densification can induce stress on the Ti-6Al-4V shape. In addition, the swelling of the sacrificial powder that occurs above 1250 K can also cause stresses on the shape. Furthermore, residual stresses can also be caused by differences in thermal contraction coefficient between the two materials during cooling. However, the cone shape is sufficiently thick and mechanically resistant to undergo such stresses. The turbine blade is thinner and more brittle and breaks when the same cycle as the cone (Figure 6.a) is applied. For a more controlled production of a turbine blade, it is necessary to reduce the maximum temperature of the cycle undergone by the assembly. To determine this new dwell temperature, the main criterion is that it should be high enough to achieve an active sintering rate but lower than the swelling temperatures. Examining the experimental sacrificial material curve in Figure 4a, significant swelling starts at 1250 K, with the onset closer to 1200 K. A 50 K margin is selected to ensure no swelling occurs, and the dwell temperature is therefore set at 1150 K. Reducing the maximum temperature from 1350 K to 1150 K, the disparity in densification between the two powders is mitigated. This temperature reduction also avoids the swelling zone and reduces differences in thermal contraction on cooling. In order to maximize the densification of the turbine blade, the temperature of 1150 K is maintained for 1 hour. With the modified cycle (Figure 6.b), it is possible to produce turbine blades without breakage.
<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /><figcaption>Figure 6: Simulated average densification curves of the sacrificial powder and the Ti-6Al-4V zones (a) for the cone (b) for the turbine blade.</figcaption></figure>
### 3.4. Cone shape study

The simulation of the cone assembly sintering is reported in Figure 7.a. As expected, the sacrificial part sinters at slightly higher temperature than the Ti-6Al-4V part and reaches its maximum relative density around 1250 K. Above 1250 K, the sacrificial powder relative density decreases due to swelling. However, this difference in sintering between the two powders does not prevent the Ti-6Al-4V cone part to reach its maximum relative density at the end of the cycle where a pseudo-isostatic condition takes place. At the end of the cycle, a collapse seems to occur under the thin wall of the cone interface. Those walls are necessary to 3D printing of the interface. In Figure 7.b, the simulated cone and a 3D scan of the experimental cone are reported. On both shapes, the collapse is noticeable. This agreement indicates that the simulation is able to predict shape defects that occur during experimentation. In a previous article using the same Ti-6Al-4V powder in the sacrificial zone, it was concluded that this deformation was due to the more rigid nature of the interface c. Figure 7.c shows the microstructure at the core and edges of the cone. At the core, two populations of the α phase in size and morphology can be distinguished: i.e. a coarse α globular morphology and a finer α lamellar morphology. The presence of coarse α globular morphology implies that the β transus temperature is higher than the maximum applied temperature of 1350 K and that the applied stresses are sufficient to fragment a significant amount of the lamellae still present at 1350 K. The presence of finer α lamellar morphology assumes that the cooling rate is far enough to start nucleation but slow enough to avoid acicular morphology. The measured hardness is relatively constant over the cone section and is $439 \pm 5$ HV<sub>0.5</sub>, which is higher than that reported in the scientific community [43–45]. The relative density of the cone measured by Archimede method is 99.7%. The edge microstructure shows several differences from the core microstructure. Two skin effects are noticeable. Firstly, the microstructure of the first 20µm from the edge of the cone seems fully α phase at this scale of observation. The second skin effect occurs at the tip of the cone and is only a few microns thick. Finally, a third white phase surrounded by black
appears at the boundaries of the powders. Those disparities have already been studied in
previous work [39] where both powders (sacrificial and shape) were Ti-6Al-4V. In this previous
work, chemical analysis has shown that disparities are all due to pollution from the partially
degraded interface. The first skin effect is the development of α-case due to carbon
contamination. The third phase at grain boundaries is titanium carbide also due to carbon
contamination. Finally, the second skin effect is a titanium phase highly rich in aluminum in
contact with the alumina interface. Carbon contamination is the most likely explanation for
these skin effects and harder microstructures. However, when the interface is used with the
same Ti-6Al-4V powder in sacrificial, a traditional microstructure with lower hardness is
obtained [39]. Keeping in mind that the partial debinding of the interface is inherent, it seems
that microstructural defects and contamination can be either due to the sacrificial ceramic mix
or the consequence of the desorption of this mix, increasing the penetration of interface
contamination elements into the core of the part. This aspect is discussed in detail in the last
section.
<figure><img src="image_12.png" /><figcaption>(B) Simulated vs experimental shape</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>(C) Microstructures</figcaption></figure>

<figure><img src="image_14.png" /></figure>

Figure 7 a. Simulation of cone assembly sintering. b. The simulated and experimental cone shape. c. Core and edge microstructures of the experimental cone.

### 3.5. Turbine blade study

Figure 8.a presents the simulated sintering of the turbine blade assembly. As anticipated, the sacrificial powder exhibits sintering at a slightly higher temperature compared to Ti-6Al-4V. However, the blade reaches its maximum relative density at the end of the dwell. Towards the end of the dwell, distortions become apparent in the upper surface interface structure. For a
better examination of the turbine blade's internal interface, refer to the sectional view in Figure 8.b.

Prior to reaching the maximum temperature of 1137 K, the blade does not achieve complete densification, attaining a relative density of 75-80 % while the sacrificial powder reaches 65 % relative density. At the end of the holding period, the entire blade exceeds a relative density of 95 % as the sacrificial powder. Additionally, shape defects manifest in the twist of the blade. Both the upper and lower sections of the twist experience shrinkage. In a previous article [39], it was determined that these defects originate from a superior rigidity of the interface compared to powder beds rigidity (like for the depression of the cone).
<figure><img src="image_15.png" /><figcaption>(B) Central cutted view of above simulation</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>Figure 8: a. Simulated sintering of the blade assembly, b. Sectional view of the simulated blade assembly during sintering.</figcaption></figure>

Figure 9 presents different views of both simulated and experimental blades in order to compare the shapes. The sectional view reveals noticeable shrinkages in the twist of both the experimental and simulated blades. Plate distortions are also observable in the sectional views. Those two defects are slightly mitigated in the simulated blade and are more visible in the experimental blade. In the experimental top view, a visible roughness resembling wood-like lines can be seen, corresponding to the 3D printing layers of the interfaces. In addition, in the
top view, the twist is flattened near a vertical interface wall located within the sacrificial zone.

A corner part of the twist is broken but this breakage is due to handling and not to the sintering
process. It should be noted that there are no significant disparities in shape between the
simulated blade and the experimental blade. These experiments demonstrate that the use of the
ceramic sacrificial powder enables sintering processes to be monitored for complex shapes
without exacerbating shape distortions.

<figure><img src="image_17.png" /><figcaption>Figure 9 Several perspectives of both simulated and experimental blades</figcaption></figure>
Figure 10 reports the microstructural observations conducted on the blade. A porous skin of 600 µm thick is revealed by those micrographs that was not detected in the simulation. In the blade micrograph slice (Figure 10.a), the darker zone along the blade edges indicates the presence of this skin. Micrographs displayed in Figure 10.b and c demonstrate that the darker zone exhibits higher porosity compared to the brighter zone. In contrast, micrographs taken from the core of the blade (Figure 10.d and e) reveal a fully dense or near-fully dense material, characterized by brighter zones. The blade's overall relative density, as determined by Archimedes' method gives 93.45% of densification, with 2.37% open porosity and 4.17% closed porosity. Qualitatively, based on observation of the blade's microstructure, the 4.17% porosity seems to better represents the core blade porosity which densification is close to 96%.

The apparent low densification is due to the fact that the blade shape is very thin and highly sensitive to the skin effect. A thick or larger blade dimension should have a much higher density. In a previous study [39], it was established that the blade edges experience carbon and other element contamination during the sintering process. However, it has been already demonstrated that this skin phenomenon does not occur when Ti-6Al-4V is used as sacrificial powder. To facilitate the preparation of the sacrificial material, an aqueous slurry is employed.

One possible explanation for this skin phenomenon is that water desorption enhances the penetration of contaminants into the blade during the debinding stage. Consequently, carbon contamination could rigidify and impede the sintering mechanisms in the affected areas. The apparent core/shell microstructure of the particles in the skin seems to show this surface pollution of the particles. The presence of the thick skin at 1150 K dwell cycle limits the applicability of thin shape fabrication. However, this phenomenon is particularly seen because we use a small SPS device which ended with a blade thickness of about 2 mm. If bigger die is used, larger wall thickness shapes could be used. In this case, the skin can be eliminated by surface finishing. The simulation model used previously focuses on mechanical aspects and does not account for chemical interactions that can affect the sintering process. As a result, it could not predict this porosity skin defect.
Similarly to the cone shape, the blade's microstructure consists of a $\alpha + \beta$ mixture with a predominantly lamellar morphology of the $\alpha$ phase.

Hardness is highly variable across the turbine blade section and shows a very high average of 678 ± 46 HV<sub>0.5</sub> compared to those presented in the scientific community [43–45]. In some areas, due to high porosity, hardness is not properly measurable. As mentioned earlier, potential contamination can contribute to increased blade hardness by a mechanism of solid solution hardening. This higher hardness further supports the hypothesis of increased and deeper contamination during the debinding process, likely due to the use of an aqueous slurry in preparation. Notably, this skin effect appears to be reduced to a few microns when higher temperatures are applied, as observed in the case of the cone assembly.
<figure><img src="image_18.png" /><figcaption>Figure 10 Microstructures of the experimental blade (a) with view of the skin zones (b, c) and core zones (d, e).</figcaption></figure>
4. Conclusion

This study investigates the use of a cost-effective ceramic sacrificial powder mix for the Spark Plasma Sintering (SPS) fabrication of complex Ti-6Al-4V shapes. The powder mix is designed to replicate the sintering behavior of Ti-6Al-4V powder, exhibiting a similar shrinkage amplitude and a slightly higher sintering temperature. This ceramic powder mix also shows a swelling above 1250K. Experimental characterization of the sintering behavior of this powder mix was conducted, and finite element simulations using different materials were tested for a thick cone shape and a thin, highly complex turbine blade shape. The simulations predicted the main shape distortions originating from the interface, indicating that the sacrificial powder did not exacerbate the shape defects. However, due to the swelling of the sacrificial powder, the thin turbine blade shape experienced breakage requiring the modification of the cycle to a one-hour cycle at a lower temperature which avoids swelling. This cycle resulted in the formation of a 600 µm skin with higher porosity on the blade and increased hardness in the blade core. These findings suggest that the aqueous preparation of the sacrificial powder facilitates the penetration of carbon contamination into the blade during the debinding process. However, this skin effect is highly visible because the turbine blade shape is very small and thin. A larger at thicker blade would be less sensitive to the skin effect and a cycle at higher temperature like for the cone shape could be used to suppress the skin porosity.

5. Acknowledgements

The authors thank the Normandy RIN Region Project DEFORMINT n°2019111818193100000210000257 for the financial support of this Study and Christelle Bilot for the SEM images.
[1] A. Couret, T. Voisin, M. Thomas, J.-P. Monchoux, Development of a TiAl Alloy by Spark Plasma Sintering, JOM. 69 (2017) 2576-2582. https://doi.org/10.1007/s11837-017-2549-6.

[2] A. Couret, G. Molénat, J. Galy, M. Thomas, Microstructures and mechanical properties of TiAl alloys consolidated by spark plasma sintering, Intermetallics. 16 (2008) 1134-1141. https://doi.org/10.1016/j.intermet.2008.06.015.

[3] D. Annur, I. Kartika, S. Supriadi, B. Suharno, Titanium and titanium based alloy prepared by spark plasma sintering method for biomedical implant applications—a review, Mater. Res. Express. 8 (2021) 012001. https://doi.org/10.1088/2053-1591/abd969.

[4] Y. Long, H. Zhang, T. Wang, X. Huang, Y. Li, J. Wu, H. Chen, High-strength Ti-6Al-4V with ultrafine-grained structure fabricated by high energy ball milling and spark plasma sintering, Mater. Sci. Eng. A. 585 (2013) 408-414. https://doi.org/10.1016/j.msea.2013.07.078.

[5] T. Sadat, K. Hayashi, G. Haugou, H. Morvan, E. Markiewicz, L. Dubar, M. Bigerelle, K. Ameyama, G. Dirras, Mechanical Properties of Spark Plasma Sintering-Processed Pure Ti and Ti-6Al-4V Alloys: A Comparative Study between Harmonic and Non-Harmonic Microstructures, Compounds. 1 (2021) 41-57. https://doi.org/10.3390/compounds1010005.

[6] O. Guillon, J. Gonzalez-Julian, B. Dargatz, T. Kessel, G. Schierning, J. Räthel, M. Herrmann, Field-assisted sintering technology/spark plasma sintering: Mechanisms, materials, and technology developments, Adv. Eng. Mater. 16 (2014) 830-849. https://doi.org/10.1002/adem.201300409.

[7] X. Wei, O. Izhvanov, C. Back, C.D. Haines, D.G. Martin, K.S. Vecchio, E.A. Olevsky, Spark plasma sintering of structure-tailored ultrahigh-temperature components: First step to complex net shaping, J. Am. Ceram. Soc. 102 (2018) 1-12.
[8] Z.Y. Hu, Z.H. Zhang, X.W. Cheng, F.C. Wang, Y.F. Zhang, S.L. Li, A review of multi-physical fields induced phenomena and effects in spark plasma sintering: Fundamentals and applications, Mater. Des. 191 (2020) 108662.
https://doi.org/10.1016/j.matdes.2020.108662.

[9] T. Voisin, J.P. Monchoux, L. Durand, N. Karnatak, M. Thomas, A. Couret, An Innovative Way to Produce γ-TiAl Blades: Spark Plasma Sintering, Adv. Eng. Mater. 17 (2015) 1408–1413. https://doi.org/10.1002/adem.201500019.

[10] C. Manière, L. Durand, A. Weibel, C. Estournès, Spark-plasma-sintering and finite element method: From the identification of the sintering parameters of a submicronic α-alumina powder to the development of complex shapes, Acta Mater. 102 (2016) 169–175. https://doi.org/10.1016/j.actamat.2015.09.003.

[11] C. Manière, L. Durand, A. Weibel, G. Chevallier, C. Estournès, A sacrificial material approach for spark plasma sintering of complex shapes, Scr. Mater. 124 (2016) 126–128.
https://doi.org/10.1016/j.scriptamat.2016.07.006.

[12] K. Srinagesh, Powder metallurgy of iron and steel, Choice Rev. Online. 36 (1998) 36-0981-36-0981. https://doi.org/10.5860/CHOICE.36-0981.

[13] S. Grasso, Y. Sakka, G. Maizza, Electric current activated/assisted sintering (ECAS): A review of patents 1906-2008, Sci. Technol. Adv. Mater. 10 (2009).
https://doi.org/10.1088/1468-6996/10/5/053001.

[14] R. Orrù, R. Licheri, A.M. Locci, A. Cincotti, G. Cao, Consolidation/synthesis of materials by electric current activated/assisted sintering, Mater. Sci. Eng. R Reports. 63 (2009) 127–287. https://doi.org/10.1016/j.mser.2008.09.003.

[15] S.H. Risbud, Y.H. Han, Preface and historical perspective on spark plasma sintering, Scr. Mater. 69 (2013) 105–106. https://doi.org/10.1016/j.scriptamat.2013.02.024.

[16] Z.A. Munir, U. Anselmi-Tamburini, M. Ohyanagi, The effect of electric field and pressure on the synthesis and consolidation of materials: A review of the spark plasma
sintering method, J. Mater. Sci. 41 (2006) 763-777. https://doi.org/10.1007/s10853-006-6555-2.

[17] C. Manière, E. Nigito, L. Durand, A. Weibel, Y. Beynet, C. Estournès, Spark plasma sintering and complex shapes: The deformed interfaces approach, Powder Technol. 320 (2017) 340-345. https://doi.org/10.1016/j.powtec.2017.07.048.

[18] C. Manière, E.A. Olevsky, In situ partially degradable separation interface for fabrication of complex near net shape objects by pressure assisted sintering, Patent US 2021/0016499 A1, US2021/0016499A1 WO2019/191299A1, 2018.

[19] T. Chartier, A. Badev, Rapid Prototyping of Ceramics, Second Edi, Elsevier, 2013.
https://doi.org/10.1016/B978-0-12-385469-8.00028-9.

[20] J.W. Halloran, Ceramic Stereolithography: Additive Manufacturing for Ceramics by Photopolymerization, Annu. Rev. Mater. Res. 46 (2016) 19-40.
https://doi.org/10.1146/annurev-matsci-070115-031841.

[21] J. Deckers, J. Vleugels, J.P. Kruth, Additive manufacturing of ceramics: A review, J. Ceram. Sci. Technol. 5 (2014) 245-260. https://doi.org/10.4416/JCST2014-00032.

[22] Y. Lakhdar, C. Tuck, J. Binner, A. Terry, R. Goodridge, Additive manufacturing of advanced ceramic materials, Prog. Mater. Sci. 116 (2021) 100736.
https://doi.org/10.1016/j.pmatsci.2020.100736.

[23] Z. Chen, Z. Li, J. Li, C. Liu, C. Lao, Y. Fu, C. Liu, Y. Li, P. Wang, Y. He, 3D printing of ceramics: A review, J. Eur. Ceram. Soc. 39 (2019) 661-687.
https://doi.org/10.1016/j.jeurceramsoc.2018.11.013.

[24] C. Manière, E. Torresani, E. Olevsky, Simultaneous Spark Plasma Sintering of Multiple Complex Shapes, Materials (Basel). 12 (2019) 557.
https://doi.org/10.3390/ma12040557.

[25] C. Manière, G. Lee, E.A. Olevsky, Flash sintering of complex shapes, Appl. Mater. Today. 26 (2022) 101293. https://doi.org/10.1016/j.apmt.2021.101293.

[26] J.E. Kanyo, S. Schafföner, R.S. Uwanyuze, K.S. Leary, An overview of ceramic molds
for investment casting of nickel superalloys, J. Eur. Ceram. Soc. 40 (2020) 4955-4973.
https://doi.org/10.1016/j.jeurceramsoc.2020.07.013.

[27] G. Vidyarthee, N. Gupta, H. Khandelwal, A Study of Ceramic Core for Investment Casting, in: 2021: pp. 379-389. https://doi.org/10.1007/978-981-16-0673-1_31.

[28] S. Hocquet, V. Dupont, F. Cambier, F. Ludewig, N. Vandewalle, Densification of complex shape ceramics parts by SPS, J. Eur. Ceram. Soc. 40 (2020) 2586-2596.
https://doi.org/10.1016/j.jeurceramsoc.2019.10.038.

[29] Q. Zhao, H. Wang, B. Tu, P. Xu, Z. Jing, W. Wang, Z. Fu, KNbTeO6 transparent ceramics prepared by the combination of pressure-less sintering and pseudo hot isostatic pressing, J. Eur. Ceram. Soc. 43 (2023) 4226-4231.
https://doi.org/10.1016/j.jeurceramsoc.2023.03.025.

[30] A. Van der Laan, R. Epherre, G. Chevallier, Y. Beynet, A. Weibel, C. Estournès, Fully coupled electrothermal and mechanical simulation of the production of complex shapes by spark plasma sintering, J. Eur. Ceram. Soc. 41 (2021) 4252-4263.
https://doi.org/10.1016/j.jeurceramsoc.2021.02.010.

[31] J. Le Cloarec, PhD, Pressure assisted sintering of multi-powders assembly to produce complex shapes: coupling between SPS and 3D printing, Université de Caen Normandie, 2023.

[32] E.A. Olevsky, Theory of sintering: From discrete to continuum, Mater. Sci. Eng. R Reports. 23 (1998) 41-100. https://doi.org/10.1016/S0927-796X(98)00009-6.

[33] V.V. Skorohod, Rheological basis of the theory of sintering, Nauk. Dumka, Kiev. (1972).

[34] W. Li, E.A. Olevsky, J. McKittrick, A.L. Maximenko, R.M. German, Densification mechanisms of spark plasma sintering: Multi-step pressure dilatometry, J. Mater. Sci. 47 (2012) 7036-7046. https://doi.org/10.1007/s10853-012-6515-y.

[35] C. Manière, J.S. Diatta, C. Harnois, C. Couder, C. Bilot, S. Marinel, Pressure assisted sintering stress exponent assessment methods: Accuracy analysis and effect of sintering stress, Mech. Mater. 181 (2023) 104664.
[36] C. Manière, L. Durand, G. Chevallier, C. Estournès, A spark plasma sintering densification modeling approach: from polymer, metals to ceramics, J. Mater. Sci. 53 (2018) 7869–7876. https://doi.org/10.1007/s10853-018-2096-8.

[37] C. Manière, T. Grippi, S. Marinel, Estimate microstructure development from sintering shrinkage: A kinetic field approach, Mater. Today Commun. 31 (2022) 1–29. https://doi.org/10.1016/j.mtcomm.2022.103269.

[38] C. Manière, C. Harnois, S. Marinel, 3D printing of porcelain and finite element simulation of sintering affected by final stage pore gas pressure, Mater. Today Commun. 26 (2021) 102063. https://doi.org/10.1016/j.mtcomm.2021.102063.

[39] J. Le Cloarec, S. Marinel, C. Estournès, M. Dehmas, C. Folton, C. Manière, Stereolithography coupled with spark plasma sintering to produce Ti-6Al-4V complex shapes, J. Manuf. Process. 114 (2024) 122–135. https://doi.org/10.1016/j.jmapro.2024.01.084.

[40] J. Langer, M.J. Hoffmann, O. Guillon, Direct comparison between hot pressing and electric field-assisted sintering of submicron alumina, Acta Mater. 57 (2009) 5454–5465. https://doi.org/10.1016/j.actamat.2009.07.043.

[41] C. Manière, J.S. Diatta, C. Couder, C. Harnois, S. Marinel, Spark plasma sintering grain growth assessment by densification kinetics analysis, Scr. Mater. 228 (2023). https://doi.org/10.1016/j.scriptamat.2023.115346.

[42] M.N. Rahaman, Sintering of Ceramics, 2007.

[43] Q. Yan, B. Chen, N. Kang, X. Lin, S. Lv, K. Kondoh, S. Li, J.S. Li, Comparison study on microstructure and mechanical properties of Ti-6Al-4V alloys fabricated by powder-based selective-laser-melting and sintering methods, Mater. Charact. 164 (2020). https://doi.org/10.1016/j.matchar.2020.110358.

[44] O.O. Ayodele, M.A. Awotunde, M.B. Shongwe, B.A. Obadele, B.J. Babalola, P.A. Olubambi, Densification and microstructures of hybrid sintering of titanium alloy,
Mater. Today Proc. 28 (2019) 781-784. https://doi.org/10.1016/j.matpr.2019.12.297.

[45] N. Singh, R. Ummethala, P.S. Karamched, R. Sokkalingam, V. Gopal, G. Manivasagam, K.G. Prashanth, Spark plasma sintering of Ti6Al4V metal matrix composites: Microstructure, mechanical and corrosion properties, J. Alloys Compd. 865 (2021) 158875. https://doi.org/10.1016/j.jallcom.2021.158875.