Figure 2 depicts the sinteringbehaviour of various glasses. Addition of CaF 2 up to 6.0 mol% improves sinterability, possibly due to the increase of viscous flow of the glass. It can also be noted that the increase of the sintering temperature from 950 °C to 1000 °C reduces the linear shrinkage of the investigated sam- ples, which is probably related to the volatility of this glass-ceramics at high temperatures. By comparing the crystallisation peak temperature of glass powders (Table 2) and their sinteringbehaviour (Figs. 2 and 3) it can be concluded that the sinterability depends on the temperature at which the crystallization starts, i.e CaF 2 addition. It means that with increasing the crys- tallization temperature, the glassy phase would have enough time for viscous flow which leads to complete densification . As discussed beforehand, the high- est density was observed in samples containing 6.0% mol CaF 2 (Fig. 3). It is also observed that the max- imum density of the samples was obtained at about 950 °C. The comparison of relative density of glasses shows that the sample G4 reaches an acceptable den- sification at the temperature ~140 °C higher than its crystallization peak temperature.
Diffractometer, using cuka radiation. The powder density was measured by using a pychnometer with xylene as the medium.. Particle size of the heat treated powders was measured using Malvern Laser based par ticle size analyzer, Malver n Instrument Ltd, U.K. The surface area of the powders was measured using a Qauntasorb BET surface area analyzing instrument. The sinteringbehaviour of the circular pellets was studied in a temperature controlled programmable furnace between the temperature ranges from 900 to 1050°C. Surface morphology of the sintered components was studied using a scanning electron microscope (Steroscan model-360).
The conventional sintering practice for CGO is to heat the powder compact at elevated temperature for a desired length of time to ensure that the maximum density of the ceramic compact is achieved. As shown in Figure 1 SSS, the relative density of CGO samples sintered in this way increases continuously with the heat treatment temperature, from 1250 °C to 1550 °C, with 2 hours of annealing time. The specimens begin to show an evident densification (~68%) at temperature higher than 1350 °C and exhibit a maximum density of about 92% after sintering at 1550 °C for 2 hours as seen in Figure 1. Furthermore, the increment in density between 1450 – 1550 °C is only about 3% signifying that a saturation point has been reached for the densification of CGO. According to the investigation by Chen and Wang , a success of two-step sintering strongly depends on the choice of temperatures T 1 and T 2 where T 1 is usually greater than T 2 .
a b s t r a c t
Cordierite–TiO 2 composite (5–20 wt.%) ceramics for various stoichiometric compositions were synthesised using conventional techniques with standard raw materials. The green composite ceramic samples obtained were dried at 100 ◦ C for 12 h, followed by pre-sintering at 800 ◦ C for 3 h. The powders were dry milled for 2 h and then compacted at 240 MPa to form pellets. The sample pellets were sintered at temperatures in the range of 600–1300 ◦ C for 3 h. The study of the different phases of the samples was performed using X-ray diffraction. The poor mechanical properties of cordierite have attracted the use of TiO 2 as a dopant. A study on the inﬂuence of TiO 2 addition to ﬂexural strength, fracture toughness, hard- ness, bulk density and thermal expansion was performed. Vickers hardness value of the cordierite–TiO 2 composites of 6.46 GPa is higher than other compositions of cordierite. The cordierite–TiO 2 exhibited a fracture toughness of 3.28 MPa m 1/2 . The ﬂexural strength of the cordierite–TiO 2 was 158.47 MPa. The effect of thermal expansion was very low in the cordierite–TiO 2 . Studies using X-ray diffraction and scanning electron microscope imag- ing on samples sintered at 1300 ◦ C conﬁrmed the presence of cordierite, along with small amounts of rutile and cristobalite.
was measured using Malvern Laser based particle size analyzer. Malvern instruments Ltd, UL. The surface area of the powders was measured using a Quantasorb BET surface area analyzing instruments. The sintering behavior of the circular pallets was studied in temperature controlled programmable furnace between the temperature ranges of 900-1050°C. Surface morphology of the sintered components was studied using a scanning electron microscope (SEM), Stereoscan model - 360.
This article discusses the corrosion of sintered Fe-Mn alloys, as influenced by spatial variability in microstructure and composition. Materials of interest were manufactured by mixing an iron powder with 25, 30 and 35wt.% of manganese powder, pressing the mixtures in a die and sintering. Particles that the sintered materials were comprised of possessed ferritic/martensitic core regions and austenitic pheripheries. While the thickness of austenite shell seemed to increase with increasing Mn content for Fe-25Mn and Fe-30Mn materials, it appeared to fall to its lowest value for the Fe-35Mn material. The corrosion potential of a material exposed to Hank's solution increased with increasing Mn content for Mn-poorer materials but fell to its lowest value for the Mn-richest material. The fractal dimension of an electrochemical noise generated in corroding material was much lower than 1.5 for Mn-poorer materials but increased to the "white noise" value of 1.5 for the Mn-richest material. Corrosion rates of all materials were higher than those reported for homogeneous Fe-Mn alloys. It was concluded that while the Mn-richest material was very likely undergoing general corrosion, Mn-poorer materials suffered from galvanic interaction between particle cores and pheripheries and their corrosion was dominated by localized events.
again, likely towards solid-state sinteringbehaviour (Eberstein et al., 2009; Amoros et al., 2019; Ryan et al., 2018).
Porosity loss can describe the degree of sintering, with crystal-free fully sintered samples reaching a final equilibrium porosity of ~0.03 (Wadsworth et al., 2016). Because rigid inclusions reduce the final attainable shrinkage, the final equilibrium porosity should be greater with more rigid inclusions. Over the course of an experiment, isolated pores grow in size and small pores are eliminated. This changes the porosity character of the sample from many but small pores at the start to few but larger pores at the end which becomes more pronounced with more crystals, probably due to differences in sintering rates caused by these heterogeneities (Amoros et al., 2019). The connectivity of pores also diminishes as samples sinter and has been estimated previously using the difference between total and connected porosity (using a combination of density measurements and pycnometry), from SEM images of the sintered end product (Eberstein et al., 2009; Winkel et al., 2012; Vasseur et al., 2013; Amoros et al., 2019), and using XCT (Colombier et al., 2017). With rigid inclusions, isolated porosity appears to develop earlier in the sintering process for rigid inclusion content >53 vol% (Amoros et al., 2019).
In Section 4.7 it was noted that the behaviour of composite material under flash sintering conditions does not seem to follow a straightforward relationship based on the relative volume fractions of the parent phases. The behaviour of composite materials undergoing flash sintering has not yet been explained by the theories presented in the literature, in particular the constrained sintering observed for titania with large alumina particles contained in the matrix  and in multilayer anode/electrode structures designed for solid oxide fuel cells . The interaction between different phases despite the lack of inter-diffusion evident in the layered structure using electron microscopy suggests that an interaction of defects between the two phases must occur in order to account for the flash sintering conditions required, strengthening the argument that defect populations contribute to flash sintering behavoiur. In addition the mechanism by which sintering additives change the flash sinteringbehaviour in materials such as alumina are not fully established [87,88]. The role of defects in mediating flash sintering remains an open question, and interesting insights may be obtained from experiments using different types of composite material structures.
The firing curves listed are recommendations. As measurement of the actual temperature in each furnace can lead to different results, adjustment of individual furnace parameters through trial firing may be necessary in individual cases. We will be more than happy to advise you in this respect.
Dental frames and crowns made of zirconiumdioxide should be fired in a furnace which is only used for these products. Firing restorations made of another ceramic material or baking of liquid ceramics in the same furnace can lead to impairment of sinteringbehaviour or local discolouring.
beginning of the heating step to the end of the dwell. For all the sintering experiments, the heating from room temperature to 600°C was controlled by a preset heating program and completed within 4 min, at a heating rate of 150°C/min. From 600°C to the desired temperature, heating rate was 100°C/min. When the required temperature was reached, the electric current was shut off, the applied stress released, and the specimens were immediately cooled down in the furnace. The sintering temperature was measured by an optical pyrometer which was implanted in the SPS apparatus at 3 mm from the top of the sample surface. Discs of 30 and 60 mm diameter of approximately 5 mm in height were produced. The 30 mm discs were produced for the purpose of optimisation and for further comparison with the 60 mm discs.
Considering the TMCs reinforced by in-situ TiC formation, a selection of carbon reactants have been added to Ti-alloy including carbon nanotubes, carbon fibre cloth, and graphene [20-23]. In this study, we offer an alternative route to prepare TiC reinforced TMC using graphite flakes, which are widely used as a cheap solid lubricant. The addition of graphite flakes to Ti-alloy powders during processing is also expected to aid the powder flow, and thereby minimize the deformation of Ti powder particles when using a low-energy ball milling process. The carbon- coated precursor composite powders are used as feedstock to form TiC phase by in-situ reacting with Ti matrix via spark plasma sintering (SPS) as an efficient method to produce TMC [24, 25]. Here the fabrication, microstructural characterization and dry sliding wear behaviour of the in-situ TiC reinforced TMCs have been evaluated, and wear mechanisms discussed.
In contrast to subtractive manufacturing technologies, additive metal manufacturing corresponds to the process of fabrication of metallic objects layer by layer using a 3D computer-aided design (CAD) model . To accomplish this, various additive manufacturing techniques have been developed so far to fabricate metallic components using metal powder, which can be classified into two process families: 1) Selective laser melting (SLM) also known as direct metal laser sintering (DMLS) or selective electron beam melting (SEBM); 2) Laser metal deposition (LMD), also known as direct laser fabrication (DLF) . In DMLS/SLM process, a high energy laser beam is used to selectively melt and join metal powder layer by layer until a fully dense and functional part is fabricated . Likewise, the SEBM process functions based on a layer by layer deposition of metal powder followed by subsequent melting of each layer; however, instead of a laser in an argon protected atmosphere, a high energy electron beam is used as the heat source in a vacuum environment . A DMLS/SLM fabricated part is known to have different properties than its SEBM fabricated counterpart, mainly ascribed to different heating and cooling rates that the material experiences during sintering process associated with each technique [83,84]. In addition, the former has been characterized by a superior surface finish (arithmetic average roughness (R a ) ~ 4-11 µm) in comparison with the latter
We demonstrated that for a typical cold sintering process of ZnO ion concentration of Zn 2 + in solution, instead of dissolution of grains, is crucial to densification and that desired ion species can be dissolved in solution prior to densification. Besides ion concentration, liquid phase fraction and temperature are also strong influencing parameters. Through pressure dependence studies, we demonstrated that critical pressures for transport correlate with hydrothermal pressures and share similar dependence with process temperature. The minimum overall pressure for full densification is also determined by this hydrothermal pressure when properly sealed. These results suggest that transient hydrothermal conditions are satisfied and important for cold sintering. Pressure and its load rate may be very important for a quick compaction stage through determining the property of seal, amount of liquid extruded, and initial density. However, pressure becomes less important for the latter stages and serves mainly to maintain hydrothermal conditions. For grain rearrangement and dissolution-reprecipitation events the effect of the solution becomes dominant as it offers lubrication, capillary driving forces and mass transport media for densification. This densification process is dynamic, as it requires controlled liquid loss and changing solution pH and concentration. Densification rate is thus competing against rate of liquid phase loss and chemical reactions. Therefore, the desired properties for materials undergoing cold sintering would be high solubility at around neutral pH, small tendency for hydroxide or hydrate formation, and fast mass exchange kinetics at the solid-liquid interface. This sintering method of ZnO requires low thermal energy, relatively low pressure, and can achieve near theoretical density with small particle size reduction. With the potential to be applied to other materials, the cold sintering process would be a valuable asset within the ceramic processing toolbox for both research and industrial production.
consumption for different class of ceramics is summarized. In all the cases, some MJ are needed for each kg of final product as shown in Table I. One should consider that the main part of the energy is consumed during the sintering process; i.e. considering the production of MgO refractory 3.0-6.3 MJ/kg are needed only for firing, while the overall energy consumption lays between 3.5-7.15 MJ/kg . Today the larger part of energy required for the firing process is provided by the combustion of “natural gas, liquefied petroleum gas (propane and butane) and fuel oil EL …; while heavy fuel oil, liquefied natural gas (LNG), biogas/biomass, electricity and solid fuels (e.g. coal, petroleum coke) can also play a role as energy sources for burners”. The energy cost of natural gas is around 0.1 €/kWh, that means that the firing process requires 0.05 - 0.2 € for kg of ceramic. These values seems to be quite low but one should consider that only in Italy the refractory production in 2011 was 486,336 tons and the production of tableware was 13,200 . Hence, the energy cost for refractory production in Italy can be estimated in the order of several millions of euros. This value is only referred to the energy cost, the actual cost of the firing process should account also for plant cost and maintenance. Typically the price of a tunnel kiln for porcelain/tableware is in the order of 500,000-1,000,000 €.
field assisted sintering techniques first in that it is a free sintering technique. The distinction from microwave sintering makes is easy to make, because of the obvious differences between the use microwaves and a dc E-fields and currents. There are no applied external pressures like it SPS , though the effect of E-fields during sinterforging is being investigated . Second, the sample is primarily heated by a furnace, rather than by the electric currents themselves. Third, the voltages and currents used, therefore the total energy used in flash sintering are much lower than the energies is SPS techniques . Fourth, the E-fields used so far in flash sintering are static and they remain constant throughout the sintering treatment as long as power supply current limits are not met. Reaching these limits results in reduction of the E-field to maintain the maximum current. Typical SPS schedules involve the pulsing of the applied current, another reason why it is known as Pulsed Electric Current Sintering (PECS) .
measuring the densification of silver nanoparticles films following photonic sintering. The absorption of light emitted by a flash lamp for varying thicknesses of silver nanoparticle layers was also measured. To determine the amount and depth of sintering, SEM images were taken of a cross section of a sintered film. To better understand the process through which the nanoparticles are sintered, we calculate the absorption of the light emitted by the flash lamp by the silver nanoparticle film using the Bruggeman effective medium theory. Using the heat transfer software package Fluent™ to model the temperature profile of the films during and following sintering, we propose a model for the photonic process.
Densification of mechanically alloyed powders involves thermally activated transition of powder particle system to thermodynamically more equilibrium state through a decrease of the free surface energy. In solid state densification process decrease in surface free energy is small when compared to other sintering processes. However the distance matter has to be transported is in the order of particle size. The entire sintering process can be divided into three distinct stages, the first stage starts as soon as some degree of atomic mobility is achieved. In this stage sharp concave necks will form between individual particles. About 5% linear shrinkage can be developed during this process. In the inter- mediate stage, high curvature formed in the first stage have been moderated and microstructure consists of three-dimensional inter penetrating network of solid particles and continuous, channel-like pores. In this stage 5-10% porosity will be persist which covers most of densification. Grain coarsening starts to become important at this stage. Grain coarsening intensity will be high during final stage.
V.V. Dabhade et al  reported that the mechanism for powder metallurgy could be volume diffusion or grain boundary diffusion or the simultaneous occurrence of both. For the densification process in a traditional powder sintering method, grain growth and neck growth are the critical mechanisms that achieve densification. Grain growth is caused by coarsening which is associated with either surface diffusion or evaporation/condensation, and generally not associated with densification . Hanna Borodianska et al.  have also studied densification by a grain-boundary diffusion mechanism as for conventional sintering and the contribution from the specific pressure-assisted mechanisms for hot pressing is insignificant. Actually, the diffusion rate is determined by temperature, powder characteristics and sintering time. J.-S. Lee et al.  established that volume diffusion would take a long time, about a few hours. However, with continuous high-pressures being applied and ultra-fast forming time, Micro-FAST occurs without the coarsening of grains during the densification process. Joule heating is the main heat source during the sintering process and thus it has different densification mechanisms compared to those of conventional processes (including FAST). For example, the mechanical plastic deformation and interface melting of particles make a great contribution to the densification of powders during the Micro-FAST .
Microstructures of starting aluminum powder and spark plasma sintered aluminum are presented in Fig. 3.1 (a-b). The starting aluminum consisted of spherical particles over a wide range of particle size (<40 μm). With the SPS processing parameters investigated in this work, near complete densification (~100% relative density) of aluminum was observed. It can be observed that the grains in the sintered aluminum samples were fairly spherical with some faceted boundaries due to deformation during SPS sintering. The grain size and distribution are not significantly different than that for starting powder suggesting full densification without significant grain growth during sintering.