Effect of Alumina Proportion on the Microstructure and Technical and Mechanical Characteristics of Zirconia-Based Porous Ceramics

Abstract

The current study investigates the process of preparing and analysing porous-structured ceramics made from zirconium, aluminium, and magnesium ceramic oxides. The starch consolidation casting (SCC) technique, with different types of starches (potato and tapioca), was used for this purpose. Our objective was to methodically examine the impact of different processing factors, such as the temperature at which pre-sintering and sintering occur, and the proportions of ceramic powders, on the microstructure, mechanical characteristics, and porosity of the resultant composites. Pre-sintering effectively reduced the rate of shrinkage during the final sintering stage; this resulted in more controlled and predictable shrinkage, leading to better dimensional stability and reduced risk of defects in the final product. A higher alumina content was associated with an increase in apparent porosity and a reduction in volume shrinkage and apparent densities. The mercury intrusion porosimetry (MIP) findings concluded that the prepared porous ceramics have a multi-modal pore structure. The highest calculated compressive strength was 76.89 MPa for a sample with a porous structure, which was manufactured using 20 wt.% tapioca starch and 30 wt.% alumina content. The main advantage of alumina is its ability to improve compressive strength by refining the grain structure and serving as a barrier against fracture development.

1. Introduction

The exceptional mechanical, thermal, and chemical properties of contemporary ceramic composites have garnered significant attention in material research due to their suitability for a wide range of industrial applications. The zirconia–alumina–magnesia (ZAM) composite is well regarded for its remarkable durability, thermal stability, and resistance to corrosion and wear, making it an extraordinary composite [1]. These features make ZAM composites a great choice for structural parts, applications involving high temperatures, and even biological applications such as dental and bone implant replacements. Researchers have recognised the importance of studying porous-structured ceramics as a key area for developing materials with better permeability, low density, and wide surface area. Porous ceramics provide significant benefits in several applications, such as bone substitution [2], catalysis [3], filtration [4], ceramic capacitors [5], fuel cell electrodes [6], and thermal insulation [7]. Therefore, it is essential to develop efficient and reliable methods for producing these materials with controlled porosity. Various manufacturing techniques have been examined to achieve the desired amount of porosity and strength in ceramic composites. The methodologies include freeze casting [8], replica techniques [9], foaming [10], and sacrificial template processes [11,12,13]. Out of these options, the method of starch consolidation casting (SCC) has shown great promise and versatility [14]. This method uses a naturally occurring and environmentally degradable starch as a material that forms pores. The starch has both ecological benefits and commercial advantages. The merits of SCC are its biocompatibility, affordability, eco-friendliness, and capacity to create precisely regulated porosity structures. The SCC method starts by blending ceramic particles with a slurry that is starch-based [15]. Subsequently, the blend is poured into moulds, left to desiccate, and ultimately subjected to sintering. During the sintering process of the ceramic matrix, the starch granules are completely incinerated, leading to the formation of a porous network [16,17,18,19,20]. Using SCC offers several advantages. Precise manipulation of the size and arrangement of pores is essential for tailoring the material’s properties to specific applications. The process is environmentally benign because of the biodegradable nature of starch, and enables the creation of complicated shapes with outstanding precision and uniformity [21,22]. Moreover, the approach may be customized to accommodate various ceramic systems, making it a versatile tool in ceramic processing. Because of their exceptional biocompatibility, mechanical strength, and controlled porosity, porous ceramics are widely prized for use in biomedical applications. These materials are often used in bone tissue engineering, dental implants, and medication delivery systems. Porous materials are particularly attractive due to their structure closely resembling that of bone tissue. Their interconnected pore structure emulates the natural architecture of bone, offering scaffolds that facilitate cell adhesion, proliferation, and vascularization [23]. The increase in surface area in porous materials, conversely, enhances adhesion with host tissues [24]. Research has shown that vascularized bone tissues may form on macro-porous alumina instead of dense alumina [25]. In porous ceramics, porosity and pore diameters significantly influence cellular activity, necessitating consideration of these parameters to achieve a balance between mechanical and biological activities. Bimodal-pore ceramics, including pores in two distinct size ranges, have shown good mechanical properties and have significant potential for biological applications [26].

The objective of this study is to fabricate and analyse ceramics with a porous structure, using zirconium dioxide (ZrO₂) as the main component and including other oxides like magnesium oxide (MgO) and aluminium oxide (Al₂O₃). One of the objectives is to fabricate a porous ceramic material with favourable mechanical characteristics by using cost-effective, readily accessible, and ecologically sustainable pore-forming agents (PFAs). This research will investigate the impact of progressive pre-sintering and sintering, as well as the influence of varying concentrations of ZrO₂ and Al₂O₃, on the composition of produced phases and the shape of pore surfaces. In addition, our research work will focus on examining the impact of varying the weight percentage of ceramic oxides on the technical and mechanical properties of the porous ceramics that are produced.

2. Materials and Experimental Procedures

2.1. Materials

The zirconia powder used in this study was monoclinic zirconia (Distributor: Interkeram Kft., Budapest, Hungary), which presents a purity of 99.99 wt.% and a density of 5.71g/cm³. The average particles size of zirconia was 0.1 μm. To enhance the properties of zirconia, aluminium oxide (Distributor: Interkeram Kft., Hungary) and magnesium oxide (Sigmaaldrich.com, Steinheim, Germany) were used as additive oxides. The purity of alumina and magnesia was about 99.99 wt.% and 99.5 wt.%, respectively. The densities of these additive oxides were 3.04 g/cm³ for alumina and 3.58 g/cm³ for magnesia. An X-ray diffraction analysis verified the purity of the zirconia, alumina, and magnesia ceramic oxides [27]. The average particle sizes of alumina and magnesia particles were 0.29 μm and 0.25 μm, respectively. Potato and tapioca starches (Distributor: Balance Food Kft., Hungary), known for their significant swelling capacity compared to other starches, have been used as agents for creating pores and shaping the structure. The experiments were conducted utilising the starch in its original form, without undergoing any further processing. The densities of potato and tapioca starches were 0.65 g/cm³ and 0.60 g/cm³, respectively. The purity was 98.2 wt.% for potato starch and 97.7 wt.% for tapioca starch. A dispersant comprising 1.5 wt.% of a commercially available ammonium salt of a polycarboxylic acid (Dolapix CE64, manufactured by Zschimmers and Schwartz in Germany) was used.

2.2. Sample Preparation

The present work included the preparation of ceramic composites with varying concentrations of zirconia and alumina. These composites were made utilising SCC processes, resulting in porous structures. The starch consolidation process begins by combining various concentrations of ceramic material with a suitable quantity of starch in a solution composed of distilled water and a small quantity of dispersant. The mixing operation was carried out at room temperature using a magnetic stirrer set at 600 rpm for a duration of two hours. Following a two-hour agitation period at ambient temperature, the mixtures were then transferred into a non-porous plastic mould. The moulds containing slurries were placed in the oven at a temperature of 80 °C for a duration of two hours. Subsequently, the porous-structured samples were removed from the moulds after being allowed to solidify for approximately 18 h at room temperature. Table 1 below displays all the various ceramic compositions that were created for this investigation. Furthermore, a porous structure sample without starch was produced for use in the mercury intrusion porosimetry (MIP) test, allowing for a comparison of the opening diameter in porous structures produced with and without starch. The sample comprises 67.5 wt.% ZrO₂, 2.5 wt.% MgO, and 30 wt.% Al₂O₃. This composition was chosen for its highest apparent porosity, as demonstrated by the results of the apparent porosity test. The apparent porosity test findings, indicate that the highest apparent porosity was observed in the sample with the highest alumina content (30 wt.%). Suspension concentration expresses the solid-to-liquid ratio. The expression of “Starch content P/T/50:50 P+T” in Table 1 denotes use of different types of starch: P for using potato starch, T for using tapioca starch, and 50:50 P+T when using potato and tapioca starch at 50:50 weight percentage ratio.

After the samples had been dried, they were gradually pre-sintered in a furnace at temperatures of 500 and 1100 °C for two hours with a heating rate of 2 °C/min. Following pre-sintering, the samples underwent gradually sintering at temperatures of 300 and 500 °C for one hour, and subsequently at 1600 °C for two hours with a heating rate of 2 °C/min. The gradual application of the pre-sintering and sintering process aims to minimize the rate of volume shrinkage and the incidence of surface fractures on the samples.

2.3. Characterization Techniques

The particle size distribution of potato and tapioca starches was assessed using laboratory sieve analysis equipment for one hour. Since the sieve method is dry and does not employ distilled water as a dispersion medium, unlike the laser method of measurement, the actual particle size will not be affected and no swelling or water absorption happens. The outcomes of sieve analysis are shown in a chart showing cumulative percentage passing on the Y-axis against sieve size on the X-axis. Finding the amount of powder retained in each sieve is the first step in calculating the percentage of powder that flows through them; hence, the following equation is employed:

% Retained = (W_sieve)/(W_total) × 100%

where W_sieve is the mass of powder in the sieve and W_total is the total mass of the powder. The next step is to compute the cumulative percentage of powder retained by each sieve, thus, sum up the entire quantity of powder that is retained in each sieve and the amount retained in the previous sieves. The cumulative percentage passing (D) of the powder is determined by subtracting the fraction retained (R) from 100%.

% Cumulative Passing = 100% − % Cumulative Retained

A moisture determination balance (RADWAG MA110R) was used to measure the moisture content of the potato and tapioca starches. The moisture determination balance composite of an electric heater to vaporise moisture and a highly accurate balance to measure the weight loss. The test started with placing 6 grams of starch in the empty balance pan, followed by evenly spreading the starch powder to make a thin, uniform layer. Upon initiation of the electrical heater, the drying process begins. The drying temperature was about 105 °C. The weight of the starch powder decreased throughout the drying process as moisture evaporated. The weight reduction was quantified by a very sensitive scale. Upon completion of the test, the moisture content displayed on the device’s screen was recorded. A hydrothermal experimental method was used to ascertain the swelling rate of starch granules in purified water. Since hydrothermal testing can evaluate several samples simultaneously and consumes less starch, it is straightforward and fast to determine starch swelling power. The hydrothermal technique is heating starch in water at controlled temperatures to examine its swelling properties. The first stage in this procedure involves preparing samples by measuring a certain amount of both potato and tapioca starches (about 1 g) and incorporating it into a defined volume of distilled water, often keeping a starch-to-water ratio of 1:100 (w/v). Subsequently, thoroughly mix the solution to ensure an even dispersion of starch granules. The mixing was performed using a magnetic stirrer at an estimated rotating speed of 200 rpm. The subsequent stage involves the hydrothermal treatment of starch particles. The starch–water solution is heated using a temperature-regulated magnetic stirrer, commonly maintained at 65 °C, for varying periods (10, 20, 30, 40, 50, and 60 min). Continuous stirring is needed to forestall sedimentation and assure continuous heating. The heat will cause a swelling of starch granules as they absorb water. After heating, the samples were left to return to ambient temperature. Filter paper was then employed to isolate the hydrated starch granules from the aqueous solution in the combination. Subsequently, the supernatant (the liquid above the deposited starch granules) was carefully poured. Following that, the mass of the sediment (expanded starch granules) was measured together with any remaining water. The swelling factor quantifies the extent to which starch granules enlarge in comparison to their initial volume. There are no units; hence, it is dimensionless. The below equation was used to quantify the swelling factor (SF) of potato and tapioca starch:

SF = Weight of swollen starch/Weight of dry starch

An X-ray diffractometer (XRD) equipped with Cu Kα radiation in the 2ϴ range of 3–90° was used to carry out the crystalline structure and phase identification of the prepared porous-structured ceramics. The step width was 0.01 and the exposure period per location was 2 s. Preparation for the XRD test involves finely grinding the samples to a powder to guarantee crystals are oriented randomly. After that, the sample is carefully placed in the sample holder inside the apparatus, making sure it is centred and aligned correctly within the XRD machine to prevent data errors. Scanning electron microscopy (SEM) was used to examine the pore size, pore surface morphology, and the surfaces of the pre-sintered and sintered porous-structured samples. It is important to thoroughly clean the samples’ surfaces before imaging them to remove any contamination. Image quality might be diminished due to dust obstructing the electron path. Since the ceramic is not a conductive material, a conductive carbon coating should be applied to the samples’ surfaces to eliminate the charging effects. After carefully coating the sample, secure it in the SEM chamber using a suitable stub and conductive adhesive (silver paste or tape). After entering the sample, allow the SEM chamber to reach the required vacuum level before proceeding with imaging. Subsequently, position the sample accurately inside the SEM chamber, ensuring that the area of interest is oriented towards the electron beam. Mercury intrusion porosimetry (MIP) was used to analyse the size distribution of the channels (opening) that connect the pores together. The benefit of mercury intrusion porosimetry is that it can measure openings of varying sizes, from the nanometre range in high-pressure equipment to the tens of micrometres in low-pressure equipment. Nevertheless, it should be noted that mercury porosimetry depends on the entrance of mercury into the pore space; hence, it can only detect open pores. The pore size distribution was conducted using mercury intrusion porosimetry (MIP) device with a maximum intrusion pressure of 400 MPa, enabling the detection of opening sizes as fine as 0.01 μm. Before analysis, the sample was evacuated to remove adsorbed material that may cover its porosity. After the pretreated solid sample is weighed, it is moved to the penetrometer. After that, the sample’s pore system is vacuum-degassed until the residual pressure is no more than 7 Pa. The penetrometer was filled using analytical-grade mercury. Under a vacuum, the sample is covered with mercury. A vacuum must be maintained in order for the penetrometer to receive mercury from the reservoir. The filling pressure of a filled penetrometer is equal to the applied pressure plus the pressure contribution from the mercury head that comes into contact with the sample. The usual filling pressure is around 4 kPa. Modestly introduce nitrogen or air to gradually raise the pressure, or introduce it continuously at a slow rate corresponding to the specific opening sizes of interest. The capillary tube records the change in mercury column length at the same time. The pressure decreases to ambient once the maximum required pressure is reached. Subsequent to taking readings under low-pressure circumstances, the mercury-filled penetrometer is moved to the instrument’s high-pressure port or unit and hydraulic fluid overlays on top. The hydraulic fluid introduces mercury into the porous system. The pressure of the system is raised to the highest point at which the low-pressure measurement was taken, and then the intrusion volume is recorded at this pressure; this is where subsequent intrusion volumes are determined. The decrease in the height of the mercury column is measured until it reaches the maximum pressure needed. The experiment used 480 nN/m and a contact angle of 141.3° (at 25 °C). Mercury intrusion tests were carried out on different types of samples. For the mercury intrusion test, it is crucial to meticulously regulate both the test parameters and ambient conditions to ensure the accuracy of the test. Mercury must be compressed in order to penetrate the pores of the material. The principle behind mercury intrusion porosimetry is that under pressure, only nonwetting liquids (those with a contact angle larger than 90°) may penetrate capillaries. The MIP test relies on measuring the amount of mercury that enters a porous structure as the applied pressure changes. The measurement encompasses just the pores that are capable of being penetrated by mercury at the imposed pressure. The equivalent diameters calculated by mercury porosimetry are size measurements based on the Washburn equation, i.e., derived assuming a porous medium with straight, constant-diameter, non-intersecting cylinder pores. This indicates that the mercury porosity results in a (equivalent) pore throat diameter distribution in the case of large pore cavities (cells) linked by tiny pore channels or throats (window cells).

p · r = −2 γ cosϴ

where p is the absolute applied pressure, r is the pore radius, γ is the mercury surface tension, and ϴ is the contact angle. The apparent densities of ceramic samples having a porous structure were determined by directly measuring both their weight and size sintering. The water absorption method was used in order to calculate the apparent porosity of the ceramic composite samples. In order to determine the apparent porosity (AP), the samples were weighed after sintering (W_sint.), then again after being immersed in water for 24 h at room temperature (W_imm.), and finally after being saturated with water (W_sat.). The apparent porosity (AP) was determined by solving for the following equation:

$\mathrm{}\mathrm{A}\mathrm{P}\mathrm{}=\mathrm{}{\displaystyle \frac{{\mathrm{W}}_{\mathrm{i}\mathrm{m}\mathrm{m}.\mathrm{}}-\mathrm{}{\mathrm{W}}_{\mathrm{s}\mathrm{i}\mathrm{n}\mathrm{t}.\mathrm{}}}{{\mathrm{W}}_{\mathrm{i}\mathrm{m}\mathrm{m}.\mathrm{}}-\mathrm{}{\mathrm{W}}_{\mathrm{s}\mathrm{a}\mathrm{t}.\mathrm{}}}}\times 100\%; (\%)$

The volume shrinkage (V_S) of the prepared porous-structured ceramics after sintering at different temperature was determined by measuring the volume of the samples before sintering (V₀) and the volume of the samples after sintering (V₁) using the equation below. The volume of the samples is calculated by using the average of three measurements taken with a digital Vernier calliper.

$V_{\mathrm{S}}={\displaystyle \frac{V_{0 }-V_{1 }}{V_{0 } }}\times 100\%; (\%)$

A universal testing machine was used to measure compressive stress of the sintered composite cylindrical samples with approximately 5 mm diameter and 8 mm height at 2 mm min⁻¹ loading rate. The findings of each test were repeated eight times for each composition in order to ensure that they were as accurate as possible. Compressive strength was measured by using the equation below, taking the average of eight samples for each composition.

σ = F/A

where σ is the compressive strength, F is the applied load, and A is the cross-sectional area of the samples A = π·r² mm².

3. Results and Discussions

3.1. Starch Characterization Test Results

Laboratory sieves were used to examine the particle size distribution of potato and tapioca starch powders. Sieve analysis presupposes that all particles are spherical or about round and will traverse the square opening when the particle diameter is less than the dimension of the square aperture in the screen. Figure 1 depicts the particle size distributions of potato starch and tapioca starch, using cumulative oversize data plotted vs. sieve size.

Cumulative oversize refers to the percentage of particles that remain over an appropriate sieve size. The data suggest that potato starch and tapioca starch have relatively equivalent particle size distributions, with modest changes in granule size proportions at particular sieve sizes. The potato starch curve demonstrates a somewhat steeper rise compared to the tapioca starch curve, suggesting that the particles in the former are dispersed across a little narrower range. Tapioca starch has a somewhat greater quantity of particles in the larger size range (>40 µm). The particle sizes of tapioca starch are approximately 78 μm, 48 μm, and 28 μm, whereas those for potato starch are around 75 μm, 45 μm and 28 μm for 90%, 50% and 10% of the total, respectively. The moisture content of potato and tapioca starches has been determined using a moisture determination balance. Upon completion of the test, the moisture content, expressed as a percentage, is displayed on the device’s screen. The measured moisture content of potato starch powder was about 15%, whereas that of tapioca starch powder was approximately 11%. Figure 2 illustrates the swelling properties of potato and tapioca starch. Tapioca starch exhibits a faster swelling rate and reaches its maximum swelling capacity faster than potato starch.

The rapid swelling is due to the increased amylopectin content in tapioca starch, which enhances water absorption and disrupts the granule structure more efficiently. In contrast, potato starch demonstrates a more gradual and slower increase in swelling capacity before reaching a plateau. Potato starch typically has larger granules, including around 80 wt.% amylopectin and roughly 20 wt.% amylose. It expands significantly, demonstrating more pronounced and gradual growth with increasing temperatures. Potato starch granules exhibit increased vulnerability to rupture after substantial swelling. Tapioca starch has a much higher amylopectin content, sometimes exceeding 85–90 wt.%, with a decreased amylose concentration. This composition facilitates rapid expansion, leading to expedited but less controlled swelling. Tapioca starch typically exhibits rapid swelling and gelatinization at lower temperatures compared to potato starch [28,29]. Research results indicate that tapioca starch expands more rapidly and stabilizes more quickly, while potato starch swells at a slower pace but may exhibit more swelling in the latter stages. The swelling properties of starch directly influence the pore structure of the resulting ceramic product. Potato starch, distinguished by its larger granules and superior swelling capacity, often produces ceramics with larger, more irregular voids. In contrast, tapioca starch, because of its small granules and controlled swelling properties, produces more uniform and smaller holes. The choice of starch is essential depending on the desired pore size distribution and application.

3.2. Apparent Density Test Results

The apparent density test results for the prepared porous-structured ceramic composites after pre-sintering and sintering are presented in Figure 3a,b, respectively.

The calculated apparent density of pre-sintered samples varies from roughly 1.84 g/cm³ for a sample prepared with tapioca starch and 15 wt.% alumina to about 1.28 g/cm³ for a sample prepared with potato starch and 30 wt.% alumina. The apparent density of sintered samples varies from roughly 3.19 g/cm³ for a sample prepared with tapioca starch and 15 wt.% alumina to about 2.4 g/cm³ for a sample prepared with potato starch and 30 wt.% alumina. According to the findings, alumina content negatively correlates with the apparent densities of the pre-sintered and sintered samples. A reduction in density was seen with an increase in alumina content due to the fact that the density of alumina, about 3.04 g/cm³, is lower than that of zirconia, which is approximately 5.71 g/cm³. The findings additionally indicate that samples produced with tapioca starch exhibit a higher apparent density compared to those produced with potato starch due to the larger swelling factor of potato starch, resulting in a higher degree of porosity. Heating with a slow rate of 2 °C/min during pre-sintering and sintering enabled good diffusion and binding of particles and subsequent pore (which are formed between the ceramic particles and not those pores formed by the swelling of starch granules) minimization.

3.3. Apparent Porosity Test Results

The apparent porosity test results for the prepared porous-structured ceramic composites after pre-sintering and sintering are presented in Figure 4a,b, respectively.

Figure 4 indicates that the porosity of ceramics is mostly influenced by the alumina content, since alumina has a significant impact on the density of the samples. The apparent porosities for all samples were increased with the addition of alumina. The calculated apparent porosity after pre-sintering varies from roughly 56% for a sample consisting of 15 wt.% alumina and made with tapioca starch to around 69% for a sample consisting of 30 wt.% alumina and made with potato starch. The apparent porosity of sintered samples ranges from approximately 50% for a sample prepared with using tapioca starch and 15 wt.% alumina to about 60% for a sample prepared with potato starch and 30 wt.% alumina. Due to its greater swelling factor compared to tapioca starch, potato starch has more pronounced form alterations during swelling and gelatinization. As a consequence, most samples exhibit a higher apparent porosity.

3.4. Volume Shrinkage Test Results

Figure 5a,b provide a comparison of the values of volume shrinkage of porous-structured samples after pre-sintering and sintering as a function of the alumina concentration as well as the type of starch present in the suspension. As can be seen from this figure, as the proportion of ceramic alumina powder in the suspension increases, the volume shrinkage reduces as the density of the samples is reduced. A matrix with a significant amount of porosity exhibited a propensity for shrinking during sintering as a result of the vacant space between the particles of the matrix. Nevertheless, the sintering process did not eliminate the presence of large pores in the matrix. In contrast, the higher shrinkage for the potato samples revealed the presence of a proportion of smaller and less stable pores inside the matrix, which were subsequently removed at 1600 °C. When starch is utilised as a pore-forming agent, the shrinkage that occurs during firing is solely attributed to the ceramic matrix. The residual pores resulting from the starch do not contribute to the shrinkage; rather, their size decreases proportionally to the whole shrinkage in the ceramic matrix. The sintering pressure, which is the driving force behind sintering, is not sufficient for pore shrinkage to happen since these pores have a high diameter and a low curvature. The volumetric sintering shrinkage of the prepared samples was reliant on the volume fraction of the ceramic particles, which decreased little among the various starch varieties. For each starch type, the increase in shrinkage after pre-sintering and sintering could be attributed to the sintering of the ceramic particles. As a result, the increased sintering shrinkage was caused by the less efficient ceramic packing in the matrix.

3.5. Microstructures of Porous-Structured Ceramics

The microstructures of the porous-structured ceramics that pre-sintered at 500 °C and 1100 °C progressive temperatures are shown in Figure 6a–c. Figure 7a–c illustrate the microstructures of porous-structured ceramics that were sintered at progressive temperatures of 300, 500, and 1600 °C. All of the samples that were utilized in the scanning electron microscopy test have a composition that is composed of 67.5 wt.% ZrO₂, 2.5 wt.% MgO, and 30 wt.% Al₂O₃, and they have the highest apparent porosity. Figure 6 and Figure 7 illustrate that in all the pre-sintered and sintered samples, large areas with a uniform distribution of alumina and zirconia grains were present. This confirms the effectiveness of the mixing procedure that was followed to produce a good dispersion of alumina particles in the zirconia matrix. There were two kinds of pores evenly dispersed throughout the ceramic grain matrix: larger, starch-granule-created pores that were either spherical or semi-spherical, and smaller, channel-like pores between the ceramic particles. Figure 8a,b compare the SEM micrographs of two samples: one that was sintered immediately without pre-sintering (a) and another that was pre-sintered before sintering (b). The scanning electron micrograph (a) clearly shows the ceramic shells contained inside the pores. The creation of these shells was greatly influenced by the interaction between ceramic powder and tapioca starch particles. When the starch particles in the matrix are heated, a layer of ceramic particles that had been attached to their surfaces separates from the rest of the matrix and shrinks with the starch particles.

The development of these shells is diminished when the samples undergo pre-sintering and subsequent gradual sintering. The diminished formation of these shells is attributed to the prolonged exposure to holding temperatures of 300 and 500 °C, which causes the ceramic particles around the swollen starch to undergo densification and compaction. SEM micrograph (b) illustrates the agglomeration of ceramic particles inside the pore. By holding temperatures of 300 and 500 °C for an extended duration throughout the progressive pre-sintering and sintering processes, the combustion of starch granules is prolonged. This allows the ceramic shells around the granules to coalesce inside the pores. In the present situation, the adhesion between the ceramic particles and the starch granules is stronger than the force of compaction or packing between the ceramic particles surrounding the swollen starch granules. The formation of this agglomeration of ceramic particles inside the pore occurs only when tapioca–potato (50:50 wt.% ratio) starches are used in the sample preparation, and is not seen when potato or tapioca starches are used. This suggests that combining these two starches will improve the connection between the gelatinized starch granules and the surrounding ceramic particles. Figure 9a–c present a comparison of the pore size and shape produced by three distinct types of starch after pre-sintering. Compared to microstructures produced using tapioca starch, porous ceramics manufactured with potato starch have bigger pores and a longitudinal pore morphology. The reason for this is because the swelling factor of potato starch is around 4.1, whereas that of tapioca starch is roughly 3.2. Following pre-sintering, the pores derived from the potato starch exhibited sizes greater than 150 μm and reached roughly 232 μm. The pores created by the tapioca starch have an average diameter of about 92 μm. Figure 9c depicts the existence of varying pore diameters as a result of the presence of two distinct starch kinds. Tapioca starch is responsible for the formation of small-sized pores, whereas potato starch granules contribute to the creation of medium-sized pores. The interaction between these two forms of starch leads to the formation of extremely large pores. The figure demonstrates that the largest pore has a significant size of around 444 μm, which results from the interaction between the two distinct types of starch granules.

A comparison of the pore size and shape after progressive sintering at 330, 500, and 1600 °C is shown in Figure 10a–c. The results are similar to those after pre-sintering: microstructures produced from using tapioca starch have smaller pores and approximately circular shape as compared with the pores produced using potato starch. As a consequence of the process of sintering, the ceramic particles that are situated in the vicinity of the pore surface experience volumetric shrinkage, which ultimately leads to a decrease in the measurement of the pores. When potato starch was used, the pores in the produced ceramics were larger than 100 μm, while the pores produced by tapioca starch were around 84 μm in diameter. In all cases, the pore morphology of the produced porous structure depends on the type and amount of pore-forming agent and the sintering temperature.

3.6. Mercury Intrusion Porosimetry Test Results

The pore size distribution, obtained using mercury intrusion porosimetry at a pressure of 400 MPa and analysed using the Washburn equation, is generally consistent with the diameters of the pore throats and necks. Figure 11 displays the distribution of pore opening sizes as a function of relative pore volume. Frequently, the analytical data reveal different groups of pore diameters, which may be related to the deliberate structuring of the materials. The mercury intrusion porosimetry findings indicate that the produced ceramics with porous structure have a multi-modal pore structure. The sample prepared without starch exhibits around 42% of the pore-opening volume within a 1–2 μm range. About 39% of the pore opening volume in the sample with tapioca starch is in the 3–7 μm range and about 10% is less than 0.01 μm. On the other hand, the potato starch sample has a bigger opening, which makes up around 45% of the volume of the pores in the 6–9 μm range and about 12% less than 0.01 μm. The sample produced with a (50:50) mixture of potato and tapioca starches has a distinct pattern: the curve representing the distribution of pore sizes shows four distinct peaks, each corresponding to a different diameter of pore opening. In this sample, almost 17% of the volume of pore openings had a diameter less than 0.01 μm, 4% had a diameter between 0.03 and 0.05 μm, 6% had a diameter between 2 and 6 μm, and approximately 17% had larger openings with diameters ranging from 10 to 100 μm. The predominant portion of the overall pore volume in all samples exists in pores with a diameter exceeding 1 μm.

The quantification can be determined by the median pore size (measured by volume) obtained from mercury intrusion measurements. The median pore diameter, also referred to as the most probable pore size, represents the point at which the concentration of pores corresponding to a specific size is highest. It is the pore size that corresponds to the peak of the pore distribution and has the highest probability of occurrence. Table 2 presents the most significant findings from the mercury intrusion porosimetry testing. The surface area is calculated by adding up the equivalent surface areas of all the pores in the test sample per unit mass.

To be precise, mercury intrusion porosimetry does not reveal the actual diameters of the macropores; rather, it reveals information on the characteristics that regulate the entry of mercury, namely, the openings connecting the neighbouring macropores. One notable aspect is that these samples with a porous structure can withstand being impregnated with mercury without crumbling. Figure 12 depicts the correlation between the applied pressure and the total volume that has been intruded. Under conditions of low pressure, the curve first exhibits a significant rise in the volume of intrusion. This suggests the infiltration of bigger pores that need just a little amount of pressure to allow mercury to enter. These sizable openings may include macropores, which play a vital role in determining material characteristics such as adsorption and catalysis. The rate of incursion volume accumulation may increase as the pressure increases, suggesting a shift from filling mostly bigger pores to smaller ones. It is common for the curve’s slope to vary during this transition zone. It shows that the material has pores of varying sizes, from very big to very tiny.

At certain points, the curve may flatten out, meaning that the incursion volume either grows very slowly or remains constant regardless of how much pressure is applied. This level indicates that most of the material’s pores that may be reached have been filled. Both macropores and linked pore networks may be present in these pores. Following the plateau, the curve may show a progressive rise in incursion volume at very high pressures. The existence of extremely tiny pores or empty spaces in this tail area can only be accessed at high pressures.

3.7. X-Ray Diffraction (XRD) Results

Figure 13 and Figure 14 show the diffraction patterns of the composite materials after the pre-sintering procedure. Figure 15 and Figure 16 show the diffraction patterns of the composite materials after the sintering process. The composition of the samples shown in Figure 13 and Figure 15 is 82.5 wt.% zirconia, 15 wt.% alumina, and 2.5 wt.% magnesia. The samples shown in Figure 14 and Figure 16 consist of 67.5 wt.% zirconia, 30 wt.% alumina, and 2.5 wt.% magnesia. The findings unambiguously demonstrate that the two largest peaks, namely, the (111) and (−111) planes at 2ϴ, have angles of around 28° and 31°, respectively. These peaks correspond to the monoclinic crystal structure of baddeleyite. Several peaks of corundum (α-Al₂O₃) were detected at 2ϴ = 26°, 35°, 38°, 43°, 52°, 57°, 66°, and 68° for planes (012), (104), (110), (113), (024), (116), (216), and (300), respectively. The findings indicate that the peak seen at a diffraction angle of 2ϴ = 37° is mostly caused by diffraction from the basal plane (111) of the periclase MgO phase.

The XRD patterns indicate that the intensity of the peaks related to zirconia increased after the sintering process, as shown in Figure 15 and Figure 16. Sintering improves the crystallinity of materials by facilitating the reorganization and bonding of atoms inside the crystal lattice. Consequently, the diffraction peaks related to zirconia may exhibit increased sharpness and intensity as a consequence of enhanced crystal arrangement. Furthermore, sintering often leads to the compaction of ceramic matrix, hence reducing porosity. Minimising the quantity of flaws and empty spaces in the material enables a larger density of atoms that may engage in X-ray scattering, resulting in increased peak intensities. Minimising stress may also amplify the magnitude of the diffraction peak patterns. The sintering process may have alleviated any pre-existing internal stresses or strains in the material. The decrease in strain might possibly lead to a more advantageous crystal structure for X-ray diffraction, resulting in heightened peak intensities. Sintering may sometimes increase the homogeneity and uniformity of the material, leading to enhanced X-ray diffraction outcomes. Possible flaws or irregularities in the pre-sintered material may have been minimised during the sintering process, resulting in clearer and stronger diffraction peaks. The results indicate that the peak intensity of the zirconia phase decreases as the concentration of alumina rises from 15 wt.% to 30 wt.%. As the concentration of alumina increases, the proportion of zirconia in the composite decreases. The diluting effect may result in a decrease in the strength of the XRD peaks associated with the zirconia phase. Increasing the alumina concentration causes the broadening of the zirconia peak, which indicates that the alumina acts as a grain-refining agent.

3.8. Compressive Strength Test Results

Figure 17 depicts the relationship between the compressive strength of the sintered ceramics and the alumina concentration (wt.%) for different kinds of starch. The highest calculated compressive strength was 76.89 MPa for a sample manufactured with 20 wt.% tapioca starch and 30 wt.% alumina content. The measured apparent density and apparent porosity were 2.78 g/cm³, and 52%, respectively. The lowest calculated compressive strength was 25.82 MPa for a sample, which was prepared using 20 wt.% potato starch and 17.5 wt.% alumina content. The measured apparent density of this sample was 2.76 g/cm³, whereas the observed apparent porosity was 56%.

Figure 17 demonstrate a positive correlation between the compressive strength of the produced porous-structured ceramics and the amount of alumina present. The main advantage of alumina is its ability to improve compressive strength by refining the grain structure and serving as a barrier against fracture development. Alumina, when added to zirconia, performs as an inhibitor of grain growth in the sintering process. Consequently, the zirconia grains undergo a reduction in size and a more even dispersion. A reduction in grain size results in a more uniform microstructure. A more homogeneous grain size distribution minimises the occurrence of weak areas in the material. The field of material science generally associates the existence of homogeneous grain sizes with improved mechanical properties, such as higher compressive strength. Alumina may improve the bonding between zirconia grains, hence improving the material’s cohesion and lowering its vulnerability to failure under stress. Alumina has a crucial role in the crack-deflection process. When a crack encounters a stiff alumina particle, it is either deflected or forced to change its path. A zigzag fracture pathway increases the energy required for crack propagation, hence greatly improving the toughness of the material. Fracture deflection is a process that absorbs energy that would otherwise encourage the spread of fractures. The material’s fracture resistance under compressive stresses is enhanced by the absorption of energy. The presence of alumina particles distributed throughout the zirconia matrix serves as an obstacle that impedes the propagation of cracks. The combined impact of grain refining and fracture deflection leads to a more robust material capable of withstanding greater compressive stresses. Smaller, homogeneous particles uniformly disperse stress over the material, decreasing the chances of localised failure. The mechanism of fracture occurred when the load application was parallel to the origin, as seen in Figure 18.

Compression failure occurs in brittle ceramics by progressive microfracture, when microcracks grow to create failure planes [30]. Under compression, microscopic cracks propagate steadily, widening in response to rising stress levels, and eventually coalesce to produce ultimate collapse [31,32]. Figure 16 shows the “kneading” phenomena, which is an intriguing finding since it shows the sample hardening at the start of the load application. The microstructural components in porous materials might fail due to the existence of many pores, which causes microcracking [33].

3.9. The Impacts of Pre-Sintering

Pre-sintering offers several benefits in ceramic technology. The current study employs pre-sintering to decrease the rate of shrinkage during the final sintering process and mitigate thermal stresses. Pre-sintering decreases the rate of volume shrinkage by allowing a substantial fraction of densification and particle rearrangement to take place during the pre-sintering stage. This minimises the amount of further shrinkage required during the final sintering process. By enabling partial densification of the ceramic material in the pre-sintering step, thermal stress is minimised during the final sintering process. Thermal stress and cracking might occur due to the rapid and uneven heating that takes place during the final sintering process. Pre-sintering mitigates these concerns by dividing the densification process into two parts. The SEM micrographs in Figure 19a,b show a comparison between those samples that was sintered with pre-sintering (a) and sample that was sintered without pre-sintering (b). These samples were prepared using 82.5 wt.% zirconia, 15 wt.% alumina, and 2.5 wt.% magnesia. Based on the findings, it seems that the pre-sintering method is effective in reducing the development of cracks on the surface of the samples.

4. Conclusions

This research work was devoted to the characterization of the porous-structured ceramics (from the system ZrO₂-Al₂O₃-MgO) prepared using 20 wt.% of different types of starch and 60 wt.% suspension concentration. Potentially derived conclusions are as follows:

1. As the alumina content in the pre-sintered and sintered samples increases, the apparent densities and volume shrinkage decrease while the apparent porosity increases.

2. SEM analysis of the microstructure illustrates that in all the pre-sintered and sintered samples, large areas with a uniform distribution of alumina and zirconia grains were present. It confirms the effectiveness of the mixing procedure that was followed.

3. SEM and MIP results reveals that the ceramic matrix contains two kinds of evenly distributed pores: big, spherical or semi-spherical pores produced by starch granules, and smaller, void-like channels that developed between ceramic particles as a consequence of sintering.

4. The presence of ceramic shells inside the pores is diminished when the samples undergo pre-sintering and subsequent gradual sintering. By holding temperatures of 300 °C and 500 °C for an extended duration throughout the progressive pre-sintering and sintering processes, the combustion of starch granules is prolonged. This allows the ceramic shells around the granules to coalesce inside the pores.

5. Compared to microstructures produced using tapioca starch, porous ceramics manufactured with potato starch have bigger pores and a longitudinal pore morphology.

6. The microstructures produced using 50:50 wt.% potato and tapioca starches have varying pore diameters. The interaction between these two forms of starch leads to the formation of extremely large pores.

7. Progressive pre-sintering and sintering of the samples results in a reduction in the rate of shrinkage and assists in reducing thermal stresses.

8. Increasing the alumina concentration causes broadening of the zirconia peak, which indicates that the alumina acts as a grain-refining agent.

9. The highest calculated compressive strength was 76.89 MPa (20 wt.% tapioca starch and 30 wt.% alumina), while the lowest calculated compressive strength was 25.82 MPa (20 wt.% potato starch and 17.5 wt.% alumina).

10. A positive correlation between the compressive strength of the produced porous-structured ceramics and the amount of alumina present is demonstrated.

11. The main advantage of alumina is its ability to improve compressive strength by refining the grain structure and serving as a barrier against fracture development.

In summary, the porosity, pore size, pore opening, and excellent mechanical compressive strength of the fabricated porous-structured ceramics render them appropriate for use in porous bone replacement applications.

Author Contributions

Conceptualization, R.A.S. and R.G.; methodology, R.A.S. and R.G.; software, R.A.S.; validation, R.A.S., R.G., M.T.M., T.T. and F.M.; formal analysis, R.A.S., R.G., M.T.M., T.T. and F.M.; investigation, R.A.S. and R.G.; resources, R.A.S. and R.G.; data curation, R.A.S., R.G., M.T.M., T.T. and F.M.; writing—original draft preparation, R.A.S.; writing—review and editing, R.A.S., R.G., M.T.M., T.T. and F.M.; visualization, R.G.; supervision, R.G.; project administration, R.G.; funding acquisition, R.A.S., R.G., M.T.M., T.T. and F.M. All authors have read and agreed to the published version of the manuscript.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data are contained within the article.

Acknowledgments

The authors would like to express their gratitude to the University of Miskolc, particularly Institute of Energy, Ceramic and Polymer Technology to use their laboratories.

Conflicts of Interest

The authors declare no conflict of interest.

Footnotes

Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Figures and Tables

Figure 1. Particle size distribution of potato and tapioca starches.

Figure 2. Swelling factor via mixing time for potato and tapioca starches.

Figure 3. Apparent density via Al2O3 content for porous ceramics: (a) after pre-sintering; (b) after sintering.

Figure 4. Apparent porosity via Al2O3 content for porous ceramics: (a) after pre-sintering; (b) after sintering.

Figure 5. Volume shrinkage via Al2O3 content for porous ceramics: (a) after pre-sintering; (b) after sintering.

Figure 6. SEM micrographs of porous ceramics after pre-sintering (a) using potato starch, (b) using tapioca starch, and (c) using tapioca–potato starch at 50:50 wt.% ratio.

Figure 7. SEM micrographs of porous ceramics after sintering (a) using potato starch, (b) using tapioca starch, and (c) using tapioca–potato starch at 50:50 wt.% ratio.

Figure 8. SEM micrographs of porous ceramics after sintering: (a) using tapioca starch without pre-sintering; (b) using tapioca–potato starch at 50:50 wt.% ratio with pre-sintering.

Figure 9. SEM micrographs of porous ceramics after pre-sintering (a) using potato starch, (b) using tapioca starch, and (c) using tapioca–potato starch at 50:50 wt.% ratio.

Figure 10. SEM micrographs of porous ceramics after sintering (a) using potato starch, (b) using tapioca starch, and (c) using tapioca–potato starch at 50:50 wt.% ratio.

Figure 11. The pore size distribution via relative pore volume.

Figure 12. The total volume intrusion via applied pressure.

Figure 13. XRD pattern for pre-sintered ceramic with 15 wt.% alumina.

Figure 14. XRD pattern for pre-sintered ceramic with 30 wt.% alumina.

Figure 15. XRD pattern for sintered ceramic with 15 wt.% alumina.

Figure 16. XRD pattern for sintered ceramic with 30 wt.% alumina.

Figure 17. Compressive strength vs. alumina content.

Figure 18. Fracture mode for sintered ceramic with 15 wt.% alumina.

Figure 19. SEM micrographs of sintered porous ceramics (a) with pre-sintering and (b) without pre-sintering.

Table 1

Mixture compositions (in wt.%).

m-ZrO2	MgO	Al2O3	Dolapix CE64	Starch Content P/T/50:50 P+T	Suspension Concentration
82.5	2.5	15	1.5	20	60
80	2.5	17.5	1.5	20	60
77.5	2.5	20	1.5	20	60
75	2.5	22.5	1.5	20	60
72.5	2.5	25	1.5	20	60
70	2.5	27.5	1.5	20	60
67.5	2.5	30	1.5	20	60

Table 2

Mercury intrusion porosimetry results.

Sample	Total Intruded Volume (cm3/g)	Total Pore Surface Area (m2/g)	Median Pore Size (μm)
Without starch	30.97	0.151	1.1947
With tapioca	77.21	5.685	3.2937
With potato	136.48	12.702	5.321
(50:50) P+T	212.21	24.454	24.0811

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