Microstructure refinement-homogenization and flexural strength improvement of Al2O3 ceramics fabricated by DLP-stereolithography integrated with chemical precipitation coating process

In this study, the chemical precipitation coating (CP) process was creatively integrated with DLP-stereolithography based 3D printing for refining and homogenizing the microstructure of 3D printed Al2O3 ceramic. Based on this novel approach, Al2O3 powder was coated with a homogeneous layer of amorphous Y2O3, with the coated Al2O3 powder found to make the microstructure of 3D printed Al2O3 ceramic more uniform and refined, as compared with the conventional mechanical mixing (MM) of Al2O3 and Y2O3 powders. The grain size of Al2O3 in Sample CP is 64.44% and 51.43% lower than those in the monolithic Al2O3 ceramic and Sample MM, respectively. Sample CP has the highest flexural strength of 455.37±32.17 MPa, which is 14.85% and 25.45% higher than those of Samples MM and AL, respectively; also Sample CP has the highest Weibull modulus of 16.88 among the three kinds of samples. Moreover, the fine grained Sample CP has a close thermal conductivity to the coarse grained Sample MM because of the changes in morphology of Y3Al5O12 phase from semi-connected (Sample MM) to isolated (Sample CP). Finally, specially designed fin-type Al2O3 ceramic heat sinks were successfully fabricated via the novel integrated process, which has been proven to be an effective method for fabricating complex-shaped Al2O3 ceramic components with enhanced flexural strength and reliability.


Introduction 
Alumina (Al 2 O 3 ) ceramic has been widely used in thermal management applications, such as substrates, heat sinks, and packaging materials for integrated and low toughness of structural ceramics make it difficult to shape and machine complex-shaped ceramic parts [4]. The conventional ceramic-shaping methods, such as dry pressing, isostatic pressing, gel casting, injection molding, etc. [5][6][7], require the use of molds, leading to increased manufacturing cost [8], and cannot be used to fabricate Al 2 O 3 ceramic parts with highly complex geometries and interconnected channels [9], thereby limiting the wide application of Al 2 O 3 components. So there has been strong demand to develop a more effective approach to sidestep the aforementioned limitations and fabricate complexshaped Al 2 O 3 ceramic parts.
Additive manufacturing, also referred to as 3D printing and rapid prototyping, is a series of advanced manufacturing technologies to construct complicated prototypes from 3D CAD models [10]. The introduction of 3D printing technologies into the fabrication of advanced ceramics can address the issues such as mold dependence and difficulties in shaping complex-shaped ceramic parts [8]. Due to the excellent technical advantages of stereolithography such as high dimensional accuracy and good surface finish, efforts on applying stereolithography to ceramic fabrication have been made [11]. The stereolithography based additive manufacturing, including stereolithography apparatus (SLA) and digital light processing (DLP), is based on the controlled light-induced layer-by-layer photopolymerization of a liquid photosensitive resin mixed with ceramic powder [8,12]. For the DLP-stereolithography, a digital micromirror device (DMD) is used to dynamically project the mask image of sliced layers onto the surface of photosensitive resin, and the shape of a whole thin layer can be solidified simultaneously; hence it has a higher building speed compared with the SLA-stereolithography and is advantageous for the fabrication of ceramic components with a very good feature resolution [13,14].
To date, numerous studies have confirmed the viability and effectiveness of the stereolithography to fabricate Al 2 O 3 ceramic parts. For instance, the Al 2 O 3 ceramic windowpanes with fully dense microstructure were fabricated by Griffith and Halloran [15], who were the first to adopt stereolithography for ceramic freeform fabrication in 1996. The drying and debinding processes of SLA-stereolithography were optimized by Zhou et al. [16], and a defect-free Al 2 O 3 cutting tool with a relative density of 99.3% was obtained. The DLP-stereolithography was used to prepare Al 2 O 3 ceramic parts with good surface quality and a relative density of 93.2%, which demonstrated the effectiveness of DLP-stereolithography [11]. Works by Santoliquido et al. [17] and Shuai et al. [18] have detailed the use of DLP-stereolithography for the fabrication of Al 2 O 3 architectures with fine and complex lattice structures. However, it is worth noting that the presence of photosensitive resin in printed green parts and its removal during the debinding process can result in a large porosity in the debound parts [19], which can lengthen the pathways for substance migration at high temperatures and be an obstacle for the sintering densification [20,21]. For this reason, an increased sintering temperature is required for the debound samples to form dense ceramic parts. The exorbitant high sintering temperature (1550-1750 ) of ℃ stereolithography-based 3D printed Al 2 O 3 ceramics [15,16,[22][23][24][25][26][27] is expected to cause the abnormal grain growth and heterogeneous microstructure of Al 2 O 3 ceramics. For instance, as presented in Refs. [20,24], the grain sizes of the Al 2 O 3 ceramics fabricated by stereolithography were 24 and 12 times larger than the average particle sizes of the feedstock Al 2 O 3 powders, respectively; in addition, a satisfactory microstructure uniformity of the Al 2 O 3 ceramics prepared by SLA-stereolithography can hardly be obtained even if the Al 2 O 3 samples were sintered at a relative low temperature (1550 ) [27]. The abnormal grain ℃ growth and non-uniform microstructure can result in lower flexural strength and reliability of ceramics [28,29], which will further shorten service lifetime of ceramic substrates and heat sinks [30] and thus restrict their application in electronics. Therefore, it is urgent to refine and homogenize the microstructure of Al 2 O 3 ceramics prepared by stereolithography-based 3D printing, in order to improve the flexural strength and reliability of Al 2 O 3 ceramics.
Yttria (Y 2 O 3 ) is commonly used as the sintering additive in preparation of Al 2 O 3 ceramics [31][32][33][34][35][36]. The addition of Y 2 O 3 , which strongly segregates or precipitates in the form of yttrium aluminates at Al 2 O 3 grain boundaries due to its limited solubility (< 10 ppm) in Al 2 O 3 crystal lattice [31,32], is known to inhibit grain growth of Al 2 O 3 and represent an effective way for microstructure refinement, through changing the activation energy of grain boundary diffusion and motion [33,34]. As a result, the mechanical strength of Al 2 O 3 or Al 2 O 3 -based ceramic can be noticeably enhanced [35,36]. This enhancement can be significantly improved by the uniformly dispersed Y 2 O 3 or yttrium aluminate second phases in Al 2 O 3 matrix [37]. However, Y 2 O 3 is conventionally doped by the mechanical mixing (MM) process [35], and it is difficult to obtain a homogeneous microstructure, especially a uniform distribution of yttrium aluminates [38]. In order to obtain the evenly distributed Y 2 O 3 in Al 2 O 3 matrix, the co-precipitation method was used to prepare Al 2 O 3 -yttrium aluminum garnet (YAG) composite ceramics [39,40]; but great amounts of alumina and yttria precursors (nitrates) were used, and thus the yield of the composite powder is low and its preparation cost is high. Compared with the conventional MM process, the chemical precipitation coating (CP) process, which has been employed in the preparation of YAG, SiC, and translucent Al 2 O 3 ceramics [41][42][43], shows more homogeneous mixing performance and can improve the microscopic uniformity of ceramics. Besides, only a small amount of Y-precursor was used in the CP process, which can make the yield of the composite powder to be higher than that prepared via the co-precipitation method. It is a remarkable fact that the dispersive performance of additives (Y 2 O 3 ) in Al 2 O 3 matrix is expected to lead to the morphology change of second phases (yttrium aluminates), and the thermal conductivity of ceramics can be highly affected by the grain growth and the morphology of second phases, as commonly seen in AlN ceramic [44,45]. Nevertheless, this CP process has never been used for the preparation of stereolithographybased 3D printed Al 2 O 3 ceramics. To the best of our knowledge, few researches have evaluated the thermal conductivity of stereolithography-based 3D printed Al 2 O 3 ceramics, which is significantly essential for thermal management applications of Al 2 O 3 ceramics. So far, little attention has been devoted to the effect of the way of introducing additives on the thermal conductivity of Al 2 O 3 ceramics.
In this study, a CP process was used to fabricate an amorphous Y 2 O 3 coating on the surface of Al 2 O 3 powder, in order to improve the dispersive homogeneity of Y 2 O 3 in Al 2 O 3 matrix. The coating effectiveness and integrity were examined in detail by transmission electron microscopy (TEM) and X-ray photoelectron spectroscopy (XPS) measurements. In addition, to demonstrate the advantage of the CP process, the phase transformation, microstructure, and thermal and mechanical property comparisons were made between the CP and conventional MM processed Al 2 O 3 samples. The current study shows that, as compared with the conventional MM process, the CP process can effectively refine and homogenize the microstructure of the DLP-stereolithography based 3D printed Y 2 O 3 -Al 2 O 3 ceramics, thus enhancing their flexural strength and reliability without decreasing thermal conductivity.

1 Preparation of Y 2 O 3 -Al 2 O 3 composite powders
Commercially available α-Al 2 O 3 powder with an average particle size (D 50 ) of 200 nm (TM-DAR, Taimei Chemicals Co., Ltd., Japan) was used as starting material. Y 2 O 3 additives were used to refine microstructure and enhance mechanical strength of Al 2 O 3 ceramics [33,35,38]. Two batches of Y 2 O 3 -Al 2 O 3 composite powders were prepared by the CP and MM process, and both of them had the same content of Y 2 O 3 . For the conventional MM route, 95 wt% of Al 2 O 3 and 5 wt% of Y 2 O 3 (D 50 = 500 nm, Shanghai Macklin Biochemical Co., Ltd., China) powders were ball milled in ethanol for 6 h employing a planetary ball mill (QM-QX4, Nanjing NanDa Instrument Plant, China), with the milling speed set to 250 r/min. The MM system consisted of powders/ethanol/zirconia ball at a weight ratio of 1:3:2 in a Teflon container. After ball-milling, the powders were dried in a rotary evaporator and granulated through a 150 μm sieve. For the CP route, the non-aqueous precursor and precipitant solutions were used for avoiding hard agglomeration of Y 2 O 3 -coated Al 2 O 3 composite powder during drying, and the whole CP process is shown in Fig. 1. The particle agglomeration is controlled by the network of self-equilibrated forces resulted by the tensile action of capillary bonds bridging the gaps between the constituting particles, and the capillary force (F cap ) calculated by integrating the LaplaceYoung equation is proportional to the surface tension of liquid medium [46]; the surface tension of ethanol (22.3 mN/m) is much less than that of water (72.8 mN/m) [47], and thus the ethanol was used as the solvent for the precursor and precipitant in this work; moreover, this similar treatment approach can also be found in Refs. [48,49] for decreasing agglomeration degree of synthesized powders. Firstly, Y(NO 3 ) 3 ·6H 2 O (99.99% purity, Shanghai Aladdin Bio-Chem Technology Co., Ltd., China) was dissolved in ethanol to prepare the precursor solution, where the concentration of Y 3+ was accurately controlled to www.springer.com/journal/40145 0.05 mol/L. Secondly, the Al 2 O 3 powder was slowly added in the prepared precursor solution, and mixed with 1.5 wt% of dispersants of polyethylene glycol (PEG 2000). The weight ratio of the Y 2 O 3 (calculated based on the transformation of Y(NO 3 ) 3 to Y 2 O 3 ) to Al 2 O 3 was set to 5:95. Thirdly, the above suspension was dispersed in a bath under ultrasound for 60 min, and then vigorously stirred for 120 min to prevent the sedimentation of ceramic particles. Fourthly, the mixed solution (pH = 12.66) of ethylenediamine (EDA) and ethanol with the weight ratio of 1:2 was prepared as the precipitant solution, and was slowly dropped into the ceramic suspension under strong mechanical stirring to tailor the pH value of the suspension in the range of 9.3-9.5. During the dripping process of the precipitant solution, the precipitations of yttriumamine complex were formed and simultaneously deposited on the Al 2 O 3 powder surface, which can act as the preferential heterogeneous nucleation sites [50]. After the above precipitation coating process, the resultant suspension was continuously stirred for 120 min. Then, the as-prepared composite ceramic powders (Al 2 O 3 coated with the yttrium-amine complexes) were washed with ethanol and air-dried at 60 ℃ for 10 h. Finally, the dried products were calcined at 450 ℃ for 2 h to decompose yttrium-amine complexes into Y 2 O 3 , and the calcined powder was granulated through a 150 μm sieve.

2 Preparation of UV-curable ceramic suspension
Before the preparation of UV-curable ceramic suspensions, the pure Al 2 O 3 , MM, and CP processed Y 2 O 3 -Al 2 O 3 powders were separately surface modified with oleic acid (OA, Shanghai Aladdin Bio-Chem Technology Co., Ltd., China). Firstly, the ceramic powder was dispersed in ethanol, and 1.0 wt% of OA with respect to the mass of ceramic powder was used as a surface modifier. The suspension was ball-milled for 2 h employing a planetary ball mill to facilitate physical adsorption of OA on powder surface. Then the suspension was dried at 50 ℃ for 12 h to remove ethanol and the remaining powder was thermally treated at 80 ℃ for 6 h to promote chemical adsorption [27,51]. Finally, the treated ceramic powder was de-agglomerated by passing them through a 150 μm sieve.
The UV-curable ceramic suspension was prepared by adding 79 wt% of above modified ceramic powder into the photosensitive resin, which was fabricated by mixing ethoxylated pentaerythritol tetraacrylate (PPTTA, Royal DSM, the Netherlands), 1,6-hexanediol diacrylate (HDDA, Royal DSM, the Netherlands), di-functional aliphatic polyurethane acrylate (U600, Royal DSM, the Netherlands), octanol (Shanghai Aladdin Bio-Chem Technology Co., Ltd., China), and polyethylene glycol (Shanghai Aladdin Bio-Chem Technology Co., Ltd., China) with commercial dispersant (BYK9077, BYK Additives & Instruments, Germany). The above ceramic suspension was ball-milled for 6 h at 350 r/min using a planetary ball mill. After the ball-milling process, the photoinitiator (Irgacure819, BASF, Germany), with an effective absorption peak range well matched with the wavelength of the UV light used for DLP processing in this work, was mixed into the homogeneous suspension to produce a UV-curable ceramic suspension.

3 DLP-stereolithography based 3D printing
3D printing was performed at room temperature by using a DLP-stereolithograpgy based apparatus. The UV light source of the DLP printer is below the vat and has a wavelength of 405 nm with a light intensity of 10.5 mW/cm 2 . A 3D model was first created using the UG software and output to a STL file, and then the STL file was imported into the stereolithography machine and sliced into 2D images. The Al 2 O 3 green bodies were obtained by DLP-stereolithography using the above-mentioned ceramic suspensions. During the DLP-stereolithography process, the UV light selectively cured the photosensitive resin in ceramic suspension based on the 2D images and created cross-linked polymer networks to bond ceramic particles together. The photopolymerisation process was generally proceeded in a layer-by-layer pattern. The layer thickness was set to 20 μm, and the cure depth was 83.00±2.51 μm by adjusting the exposure time to 3.0 s, giving a high vertical resolution and an adequate integration between layers. Once a single layer was cured, the vat was tilted down to detach the cured layer and the building platform was lifted to allow recoating the suspension layer at the bottom of the vat. Then the new layer was cured subsequently in exactly the same fashion. These steps were repeated until the whole green body was eventually fabricated.

4 Debinding and sintering
For printed Al 2 O 3 green bodies, the post-processing steps including debinding and sintering were carried out to obtain Al 2 O 3 ceramic parts. A two-step debinding process was adopted in this work, in which green bodies were firstly debound under vacuum to decelerate pyrolysis rate of resins and then debound in air to completely remove the residual carbon [16]. The debound samples were subsequently sintered in a muffle furnace (HTK 16/18, Thermconcept, Germany) at 10 ℃/min from room temperature to 800 ℃, then at 5 ℃/min up to 1650 ℃, with a final plateau of 2 h. Finally the furnace was cooled at 5 ℃/min to 800 ℃, and then naturally cooled to room temperature. The sintered specimens were machined and polished to evaluate their thermal and mechanical properties. The 3D printed Al 2 O 3 ceramics fabricated using the MM and CP processed (i.e., ball-milled and coated) Y 2 O 3 -Al 2 O 3 composite powders are referred to as Sample MM and Sample CP, respectively. In addition, the monolithic Al 2 O 3 reference samples prepared by pure Al 2 O 3 powder are referred to as Sample AL.

5 Characterization
The microstructure and element analysis of the coated Al 2 O 3 powder were investigated by a transmission electron microscope (TEM, Talos F200S, Thermo Fisher Scientific Inc., USA) coupled with an energy dispersive spectroscope (EDS). The TEM sample (coated Al 2 O 3 powder) was dispersed in ethanol with ultrasonic treatment and then dropped onto a holeycarbon-coated copper grid. An X-ray photoelectron spectroscope (XPS, Escalab 250Xi, Thermo Fisher Scientific Inc., USA) was used to study the surface chemistry of the MM and CP processed Y 2 O 3 -Al 2 O 3 powders. The contents of Y and Al in the MM and CP processed Y 2 O 3 -Al 2 O 3 powders were determined by an X-ray fluorescence spectrometer (XRF-1800, Shimadzu Co., Ltd., Japan), in order to identify whether both composite powders have the same chemical compositions.
The viscosities of the ceramic suspensions were tested using a rotational rheometer (MCR 301, Anton Paar, Austria). To study the decomposition profile of the 3D printed green body, TG-DSC analyses were conducted by a simultaneous thermal analyzer (STA449F3, Netzsch, Germany) at a heating rate of 10 /min over the temperature range from room ℃ temperature to 600 . The relative density of the ℃ sintered Al 2 O 3 ceramic samples was measured by Archimedes method using an analytical balance with an accuracy of 0.0001 g. The bending strength of samples with a size of 1.5 mm × 2.0 mm × 25 mm was evaluated by three-point bending tests [52]. Finish grinding of four long faces was performed using a 600 grit diamond wheel, and the four long edges were rounded with a radius of 0.15 mm. The two faces with the size of 2.0 mm × 25 mm were gently polished to 1 μm by applying diamond paste. The loading experiments were performed using a universal mechanical testing machine (Inspekt Table Blue 05, Hegewald & Peschke, Germany), with the crosshead speed set to 0.5 mm/min and the supporting span of 20 mm. A laser thermal conductivity instrument (LFA 447, Netzsch Instruments Co., Ltd., Germany) was used to determine the thermal diffusivity (α) of the sintered Al 2 O 3 samples with a size of 10 mm × 10 mm × 2 mm at room temperature. The blocks were spray-deposited with a thin layer of colloidal carbon to enhance the absorption of the laser pulse. The value of the thermal conductivity (λ) for samples was calculated by where ρ is the density and C is the specific heat capacity of the prepared samples [53]. In the present work, the thermal conductivity of Sample AL was where ω i is the mass fraction of each phase which can be determined by XRD analysis, and C i is the corresponding specific heat capacity for the constituents (Al 2 O 3 and YAG). A documented value of 0.60 J·g -1 ·K -1 is used for the specific heat capacity of YAG [56,57]. The phase composition of the sintered samples was determined by X-ray diffraction (XRD, D8 Advance, Bruker Corporation, Germany). A scanning range of 2θ from 10° to 80° was applied. The mass percentage of the phases in the sintered sample was also semiquantitatively estimated by analyzing the reference intensity ratio (RIR) value taken from the X-ray pattern [58]. The weight fraction of phase 1 in the sintered ceramics can be calculated by where X 1 is the weight fraction of phase 1, I i is the integrated intensity of the highest peak of the i-th phase in the analyzed ceramic, RIR i is the reference intensity ratio of the i-th phase (taken from the powder diffraction database), and n phase is the number of phases in the prepared ceramics.
To identify the fracture surface features and global distribution profile of Y element throughout the 3D printed samples, the images of fracture surface and the corresponding X-ray mapping analysis of Y element were studied by a scanning electron microscope (SEM, LYRA 3 XMU, Tescan, Czech) coupled with an energydispersive spectroscope (Inca X-Max50, Oxford Instruments, UK). Moreover, the prepared samples were polished with diamond paste and then thermally etched at 1550 ℃ for 30 min. The microstructures of the polished and fracture surface of the sintered ceramics were characterized by the SEM. The average grain size was determined by the Nanomeasure software, and at least 600 grain sizes were statistically analyzed for each sample. The Christiansen uniformity coefficient (CU) was used to quantitatively determine the distribution uniformity of the grain size in the sintered ceramics, which can be calculated using the following equations [59]: where x is the average grain size, which can be calculated by grain grain 1 grain, n grain is the total number of grains, and the larger CU value indicates the more uniform microstructure of sintered ceramic.

1 Characterization of CP processed Al 2 O 3 powder
The weight percent of Y/(Y+Al) of the MM and CP processed Y 2 O 3 -Al 2 O 3 composite powders were measured to be 7.12 and 7.40 wt%, respectively, by an XRF spectrometer, indicating that the composite powders prepared by MM and CP processes have nearly the same Y 2 O 3 content. The microstructure of the CP processed composite powder was analyzed. TEM results of the CP processed Y 2 O 3 -Al 2 O 3 powder are presented in Fig. 2. The morphology of the CP processed Y 2 O 3 -Al 2 O 3 powder is shown in Fig. 2(a), which displays that an evident shell layer was closely, uniformly attached to the Al 2 O 3 particle surface. From a high resolution image as shown in Fig. 2(b), the CP processed Y 2 O 3 -Al 2 O 3 powder has a typical core-shell structure with a relatively smooth surface layer, which is about ~2.31 nm thick. According to a further fast Fourier transform (FFT) pattern analysis, the dispersive diffraction halo in Fig. 2(c) shows that the shell is amorphous, and Fig. 2(d) presents the lattice fringes of Al 2 O 3 (2-13) and (-114) planes with interplanar spacing of about 0.201 and 0.259 nm, respectively, indicating that the core is Al 2 O 3 crystalline. This result verifies that a uniform amorphous deposition was formed on the surface of the Al 2 O 3 particle. To check the element compositions of the amorphous layer, the EDS analyses were conducted on both amorphous layer (shell) and the bulk of the Al 2 O 3 particle (core), shown in Figs. 2(e) and 2(f), respectively. The major elements in the shell are Al, Y, and O with minor C and Cu. Obvious Al, O, C, and Cu peaks can be detected in the core, meanwhile detecting weak Y peak at about 1.93 keV. The C and Cu elements come from the holey-carbon-coated copper grid used in the TEM sample preparation. The spectral peak for Y in the shell is much higher than that in the core, suggesting the shell contains more Y than in the core. This EDS result indicates that the chemical composition of the shell is Y-O compound. The TEM and EDS results suggest that the Al 2 O 3 particle is encapsulated by the amorphous layer of Y-O compound.
The XRD patterns for the MM and CP processed Y 2 O 3 -Al 2 O 3 composite powders are shown in Fig. 3. The diffraction peaks of the MM processed composite powder can be identified as phases of α-Al 2 O 3 and cubic Y 2 O 3 , indicating that the MM processed powder is a mixture of α-Al 2 O 3 and Y 2 O 3 powders. However, the XRD pattern of the CP processed composite powder only shows sharp diffraction peaks assigned to α-Al 2 O 3 , and a very broad peak of Y 2 O 3 is present at 29.15° (which is the diffraction angle of the strongest   Table 1. The Y/Al atomic ratio of the CP processed Y 2 O 3 -Al 2 O 3 powder is calculated to be 0.2081, which is much greater than the Y/Al atomic ratio of the MM processed Y 2 O 3 -Al 2 O 3 powder (0.0023) and stoichio-metric mixture (95 wt% Al 2 O 3 + 5 wt% Y 2 O 3 ) (0.0238). The much higher value of the Y/Al atomic ratio in the CP processed Y 2 O 3 -Al 2 O 3 powder can evidence that the surface of the CP processed Al 2 O 3 powder is enriched with Y, i.e., Y is localized in the surface layer of the CP processed Al 2 O 3 powder [62]. It can be concluded from the TEM, XRD, and XPS results that the surface of the Al 2 O 3 ceramic powder is covered with an amorphous Y 2 O 3 layer prepared by the CP process.

2 Characterization of the ceramic suspensions and printed green bodies
It is important to maintain a good flowability and a low viscosity for the ceramic suspension, because the photocurable suspension should be spread and recoated by the blade [63].   To get a dense ceramic part, the organic components in the printed green body must be removed by a debinding step. A two-step debinding process, consisting of a vacuum debinding step followed by an air debinding step, was adopted in this work [16]. The lower solid loading of ceramic suspension implies the higher organic components in the printed green body, which can amplify the signal intensity of TG-DSC test and improve its accuracy to analyze the thermal decomposition of the green ceramic body. So the green compact manufactured by the Al 2 O 3 suspension with low solid loading (69 wt%) was used for the TG-DSC test, and the test results are shown in Fig. 6. Due to the fact that the TG-DSC analysis cannot be conducted in vacuum, the N 2 was used to simulate the oxygen-free environment for determining the thermal degradation behavior of printed green body in vacuum debinding step, and the test result is shown in Fig. 6(a). It can be seen that the thermal degradation of the organic components in the green body takes place over a broad temperature range between 150 and 500 , ℃ accompanied by weight loss and endothermic reactions. This ensures a continuous formation of thermal degradation products without temporary accumulation which can affect the structural integrity of the parts [23,64]. However, there exists a domain where the degradation rate is increased at the temperature range of 300-425 , as can be seen from the peak of DTG ℃ curve, manifesting the majority of photopolymer is degraded at this temperature regime. The DSC curve depicted in Fig. 6(a) features three endothermic peaks at 232, 373, and 415 , with the peak observed at ℃ 373 being the most prominent. Therefore, a slow ℃ heating rate (1 /min) was used in the whole vacuum ℃ debinding process and hold points were introduced at the onsets of the endothermic peaks.
Afterwards, the air debinding step was required to remove the residual carbon in the black vacuumdebound body, thus preventing the generation of cracks in the sintered body [16]. The TG-DSC result for the vacuum-debound body tested in air is shown in Fig. 6(b). Unlike the thermal degradation behavior of green body in the absence of oxygen, an exothermic peak at 396 and a domain where the decomposition ℃ rate is increased at the temperature range of 346-446 are present in Fig. 6(b), due to the oxidative ℃ decomposition of the residual carbon, and its weight loss and decomposition rate are much lower than those of the green body degraded in oxygen-free environment. In addition, the green Al 2 O 3 bulk sample was manufactured by the ceramic suspension with the 79 wt% solid loading, and its weight losses in the vacuum and air debinding processes were measured to www.springer.com/journal/40145 be 21.32%±0.15% and 2.12%±0.04%, respectively, indicating the weight loss in the second air debinding process is much lower than that in the first vacuum debinding process. This is conductive to avoiding formation of defects (such as cracking, delamination, blistering, and bloating) in the debound body. Thus, the heating rate of the air debinding was reduced in the domain, and a hold point was introduced at the onset of the decomposition peak at 346 . ℃ An assessment of the distribution of Y 2 O 3 in the printed green bodies was conducted via the elemental mapping analyses using SEM and EDS, as shown in Fig. 7. A significant difference in Y 2 O 3 dispersion is observed in the MM and CP processed green compacts. The MM processed compact has large Y-patches distributed heterogeneously. The size of these Y-patches varies from 0.72 to 3.84 μm, and the average value is about 1.92±0.86 μm, which is much higher than the average particle size (500 nm) of Y 2 O 3 powder, manifesting the apparent agglomeration of Y 2 O 3 particles during the MM process. On the other hand, the CP processed compact shows no apparent aggregate patches, and Y 2 O 3 is uniformly distributed in the Al 2 O 3 matrix. This finding indicates that the CP process can lead to an obvious enhanced homogeneous distribution of the sintering additives (Y 2 O 3 ) in Al 2 O 3 matrix as compared with the conventional MM process. This change in Y 2 O 3 distribution is expected to tailor the microstructure of Al 2 O 3 ceramic and morphology of second phase, which are essential to the mechanical and thermal properties of Al 2 O 3 ceramic.

3 Phase analysis and microstructure of sintered Al 2 O 3 ceramics
The XRD patterns of the 3D printed Al 2 O 3 ceramic samples are shown in Fig. 8, and the peak intensities are normalized relative to the Al 2 O 3 (211). Diffraction peaks of α-Al 2 O 3 , as the only phase for Sample AL, are exhibited in Fig. 8(a) Figure 9 demonstrates the SEM observation on the microstructures of different sintered Al 2 O 3 samples and their grain size distributions. In Fig. 9, the white zone at grain boundary areas corresponds to the Y 3 Al 5 O 12 second phase (Fig. 8), while the gray and black sections are Al 2 O 3 grains and pores, respectively. Compared with Sample AL (Fig. 9(a)), Samples MM and CP (Figs. 9(b) and 9(c)) possess higher porosities, corresponding to the relatively higher density of Samples AL as shown in Fig. 10. It suggests that the densification of Al 2 O 3 matrix was slightly inhibited by the addition of 5 wt% of Y 2 O 3 . On one hand, the yttrium segregation at Al 2 O 3 /Al 2 O 3 interfaces can block the diffusion of ions along grain boundaries, leading to a reduced grain boundary diffusivity and hence a decreased densification rate [33,34,67]. On the other hand, once above the Y 2 O 3 solubility-limit in Al 2 O 3 , the extra Y 2 O 3 would react with Al 2 O 3 matrix and then form Y 3 Al 5 O 12 precipitations via the solid state reactions (Eqs. (5)- (7)). The Y 3 Al 5 O 12 precipitations, present at grain boundaries and multigrain junctions in Al 2 O 3 matrix, can result in Zener pinning action on grain boundary mobility and finally retarding the densification of Al 2 O 3 matrix [37,68].
As shown in Fig. 9(a), the 3D printed monolithic Al 2 O 3 ceramic possesses an obvious coarsening and inhomogeneous microstructure (grain size up to 4.78 μm, with a low CU of 0.50), and Sample AL contains a few of elongated alumina grains with the evidence of abnormal grain growth. The introduction of Y 2 O 3 enables the Al 2 O 3 matrix grain size to be refined, due to the pinning effect of Y 3 Al 5 O 12 precipitations [33,37]. The mean grain sizes of Samples MM and CP are smaller than that in Sample AL by 26.78% and 64.44%, respectively. Furthermore, the grain size of Sample CP (~1.70 μm) is 51.43% smaller than that of Sample MM (~3.50 μm). It suggests that the different introducing ways of Y 2 O 3 would result in different growth trends for the Al 2 O 3 grains. Compared to Sample MM, Sample CP shows much refined Y 3 Al 5 O 12 precipitations and an increased degree of Y 3 Al 5 O 12 distribution range at grain boundary areas in the matrix. This precipitation feature displayed in Sample CP can lead to more areas of Zener pinning on grain boundary migration, and hence enable a much refined and homogenous Al 2 O 3 matrix microstructure.In addition, in both Ref. [24] and the present work, the same Al 2 O 3 powder was used; although the sintering temperature (1600 ℃) in Ref. [24] is lower than that (1650 ℃) in the present work, the grain size of Sample CP is much less than the value reported in Ref. [24] (2.6 μm). The microstructure refinement is often expected to fabricate Al 2 O 3 ceramic with higher strength according to the Hall-Petch relationship [69]. Furthermore, Sample CP has the narrowest grain size distribution, with a CU of 0.70, which is 12.90% and 40.00% higher than those of Samples MM and AL, respectively. This finding demonstrates that the CP process can make the microstructure of Al 2 O 3 ceramic more uniform, which can play a crucial role in improving the reliability of Al 2 O 3 ceramics.

4 Properties of 3D printed Al 2 O 3 ceramics
The physical properties of the 3D printed Al 2 O 3 ceramics are shown in Fig. 10. The relative density of pure Al 2 O 3 ceramic (Sample AL) is higher than those of Y 2 O 3 -Al 2 O 3 system (Samples MM and CP), demonstrating that the addition of 5 wt% of Y 2 O 3 is detrimental for the densification of Al 2 O 3 ceramic. In addition, Sample CP has a higher relative density than Sample MM, which is consistent with the denser microstructure of Sample CP (Fig. 9(c)) compared to that of Sample MM ( Fig. 9(b)), illustrating that the CP process can improve the relative density of the Y 2 O 3 -Al 2 O 3 system compared with the conventional MM process.
As displayed in Fig. 10, Sample AL has the lowest flexural strength, which is 8.45% and 20.29% lower than those of Samples MM and CP, respectively, i.e., the addition of Y 2 O 3 is favorable to the improvement of flexural strength of the 3D printed Al 2 O 3 ceramic. The improved flexural strength of the Y 2 O 3 -Al 2 O 3 system is resulted from the fine grain strengthening and dispersion strengthening of the Y 3 Al 5 O 12 particles [70,71]. The flexural strength of Sample CP is 14.85% higher than that of Sample MM, owing to the finer Al 2 O 3 grain and more homogeneous microstructure of Sample CP (Fig. 9(c)); also the relative density improvement can lead to an increase in flexural strength of Sample CP, according to the empirical suggestion for the strength of ceramics [1]. Fracture surface SEM micrographs of the 3D printed ceramic samples are shown in Fig. 11. Fractures predominantly occurred in an intergranular mode for Samples AL and MM (Figs. 11(a) and 11(b)). On the other hand, a large proportion of Al 2 O 3 grains fractured transgranularly in Sample CP (Fig. 11(c)). This indicates that the Al 2 O 3 grain boundaries in Sample CP should be much stronger than those in Samples AL and MM, leading to the increased flexural strength for Sample CP.
Furthermore, reliability analysis was conducted by estimating the Weibull modulus of the flexural strength distribution using at least 16 samples [72]. The two-parameter Weibull distribution was calculated by the equation: where P f = (n -0.5)/N is an estimator of the fracture probability of the n-th ranked sample, n is the rank of the bending strength data, N is the total number of samples tested (N = 16 in the present work), m is the Weibull modulus, σ n is the measured bending strength, and σ 0 is the Weibull material scale parameter [72]. A higher m value means a better uniformity of flexural strength and a higher reliability of ceramic. The flexural strength distribution of Samples AL, MM, and CP is shown in Fig. 12 (Fig. 9(b)) can make the homogeneity of Sample MM deterioration, thus leading the reduction of Weibull modulus. Compared to the Weibull modulus for Sample MM, a very drastic increase from 9.59 to 16.88 in the Weibull modulus for Sample CP was obtained, indicating a much higher reliability and repeatability. These above findings illustrate that the CP process is beneficial to simultaneously enhance the flexural strength and reliability of the 3D printed Al 2 O 3 ceramic. It can be seen from Fig. 10 that the thermal conductivity of pure Al 2 O 3 ceramic (Sample AL) is 14.54% and 15.37% higher than those of Samples MM and CP (Y 2 O 3 -Al 2 O 3 system), respectively. The addition of 5 wt% of Y 2 O 3 causes a fall in thermal conductivity compared with that of pure Al 2 O 3 and this is accompanied by the decreased relative density and formation of the second phase (Y 3 Al 5 O 12 ). The pores and second phases can enhance phonon scattering and decrease the effective conductive mean free path, therefore reducing the thermal conductivity according to the kinetic theory of phonons in solids [73,74]. Based on the Maxwell model, the thermal conductivity of ceramic sample with isolated pores dispersed in a ceramic matrix is written as [75] where λ is the apparent thermal conductivity of ceramic with pores, φ is the porosity of ceramic and φ = 1 -RD% (relative density), and λ 0 is the thermal conductivity of completely dense ceramic without pore (φ = 0 where λ c , λ s , and λ p are the thermal conductivities of the composite ceramic, Al 2 O 3 , and Y 3 Al 5 O 12 phase, respectively, and V p is the volume fraction of Y 3 Al 5 O 12 phase. In the present work, a value of λ p = 10.7 W·m -1 ·K -1 was used for the thermal conductivity of pure Y 3 Al 5 O 12 phase [80], which is much lower than the thermal conductivity of dense Al 2 O 3 ceramic (34.32 W·m -1 ·K -1 ). If the morphology of second phase is not completely isolated, the effect of the connectivity of second phase on the thermal conductivity of composite system cannot be ignored. An appropriate model is given by the effective medium theory (EMT), and the thermal conductivity of composite system is given by [77,81] The calculated thermal conductivities derived from Eqs. (10)-(12) are shown in Table 3. The formation of Y 3 Al 5 O 12 precipitations with a low thermal conductivity can result in a drop in thermal conductivity of Y 2 O 3 -Al 2 O 3 system based on the ME and EMT model analyses. However, the thermal conductivities of Y 2 O 3 -Al 2 O 3 system estimated by the ME and EMT models are higher than the measured values of thermal conductivities of Samples MM and CP since the phonon scattering caused by the pores, and the second phase precipitations and grain boundaries can decrease the phonon mean free path and further lead to the reduction of thermal conductivity [74,77]. Thus, the addition of Y 2 O 3 into Al 2 O 3 is expected to reduce the thermal conductivity of 3D printed Al 2 O 3 ceramic.
The grain boundaries can act as the scattering sites for phonons and reduce the thermal conductivity of ceramic [77]. This reduction can be enhanced by decreasing the grain size, which is attributed to the increased number of grain boundaries per unit length of heat path. However, compared to the coarse grained Sample MM, the fine grained Sample CP has a similar value of thermal conductivity, rather than a lower value, as shown in Fig. 10, because the thermal conductivity of ceramic is also highly affected by second phase morphology [45]. The SEM image and corresponding EDS mapping for Y on the fracture surfaces of Sample MM sintered at 1650 ℃ are shown in Figs. 13(a1) and 13(a2). Sample MM has large Y patches distributed heterogeneously (Fig. 13(a2)), and these patches are on the average size of 2.85±1.24 μm, which is comparable to the grain size of Sample MM (3.50±1.68 μm), manifesting an apparent agglomeration of Y 3 Al 5 O 12 phase during sintering. A presence of semi-connected Y 3 Al 5 O 12 aggregates can be observed in Fig. 13(a3), and the Y 3 Al 5 O 12 aggregates distributes continuously along Al 2 O 3 grain boundaries in Sample MM. The second phased distribution of Sample CP sintered at 1650 ℃ is shown in Figs. 13(b1)-13(b3). Sample CP has small Y patches distributed homogeneously (Fig. 13(b2)), and these patches are on the average size of 0.71±0.26 μm, which is much lower than the grain size of Sample CP (1.70±0.63 μm). Combined with the results obtained from SEM micrograph of Sample CP  ( Fig. 9(c)), the small-sized Y 3 Al 5 O 12 particles tend to be concentrated at multigrain junctions without continuous distribution along Al 2 O 3 grain boundaries. The morphologic change of the second phase from semiconnected (Sample MM) to isolated (Sample CP) can contribute to an improvement of the thermal conductivity of Al 2 O 3 ceramics, according to the comparison between EMT model (Eq. (12)) for the interconnected second phase and ME model (Eq. (11)) for the isolated second phase [45,76]. In addition, compared to Sample MM, the higher relative density of Sample CP can lead to an increase in thermal conductivity. From all the above analysis, Sample CP has a close thermal conductivity to Sample MM due to the combined contribution of grain size, the second phase morphology, and relative density.
According to the above research results, the CP process is an effective method to enhance the flexural strength and reliability of 3D printed Al 2 O 3 ceramic. Typical Al 2 O 3 ceramic components with complex shapes can be fabricated by a novel approach integrating DLP-stereolithography and CP process, and the images of the 3D-printed green and sintered bodies with different fin configurations are shown in Fig. 14. The surface smoothness of the green and sintered parts appears to be fine. In addition, the thermal debinding shrinkages of the components in the length, width, and height were measured to be 2.23%, 2.49%, and 1.89%, respectively. It should be noted that these shrinkage ratios are very small (lower than 3%), indicating that the thermal debinding has no apparent effect on the dimensions of the 3D-printed parts [20]. This is important for constructing parts with precise sizes and geometries. The average sintering shrinkages in the length, width, and height reached 20.90%, 21.33%, and 23.34% from their initial dimensions, respectively. The height shrinkage is larger than those in other directions, and this is caused by the layer-by-layer forming characteristic derived from stereolithography-based 3D printing technology [82]. On the whole, the sintering shrinkages of the parts are uniform without obvious geometry deformation. The successful manufacturing of the pin-type heat sinks as shown in Fig. 14 can show potential application in thermal management in electronics and automotive industries [83]. Therefore, the fundamental study of this work can offer an alternative approach to fabricate ceramic heat sinks with complex shapes and excellent mechanical performance.

Conclusions
In this study, a special Al 2 O 3 ceramic with complex shape, high strength, and fine grained and homogeneous microstructure was successfully fabricated by a novel approach integrating DLP-stereolithography and CP process. The CP process was used to synthetize Y 2 O 3coated Al 2 O 3 composite powder, which was then used to print complex-shaped Al 2 O 3 bodies via DLPstereolithography. It was found that: 1) The microstructure, phase, and surface element present analyses for the CP processed Al 2 O 3 powder clearly demonstrate that an amorphous Y 2 O 3 layer can be fully wrapped on the surface of Al 2 O 3 powder, which will be favorable for improving the dispersive homogeneity of Y 2 O 3 in Al 2 O 3 .
2) The phase and microstructure comparisons among the 3D printed Al 2 O 3 samples prepared via the pure Al 2 O 3 powder, CP, and MM processed Y 2 O 3 -Al 2 O 3 powders show: i) The introduction of Y 2 O 3 into Al 2 O 3 can result in the generation of Y 3 Al 5 O 12 phase precipitations, which can reduce the grain size of Al 2 O 3 due to the precipitation pinning; ii) the CP process can make the microstructure of the 3D printed Y 2 O 3 -Al 2 O 3 system denser, more uniform and refined, compared to the conventional MM process.
3) Sample CP has a flexural strength of 455.37± 32.17 MPa and a Weibull modulus of 16.88, which are 15.10% and 76.02% higher than those obtained for Sample MM, and are 25.45% and 13.82% higher than those of Sample AL, respectively, due to the grain refinement and microstructure uniformity enhancement in Sample CP. This result indicates that the CP process is conductive to simultaneously improving the flexural strength and reliability of the 3D printed Al 2 O 3 ceramic.
4) The CP process can refine the microstructure of Y 2 O 3 -Al 2 O 3 system at no expense of the thermal conductivity. The fine grained Sample CP has a close thermal conductivity (28.95±0.28 W·m -1 ·K -1 ) to the coarse grained Sample MM (29.16±0.55 W·m -1 ·K -1 ), because the CP process can facilitate the formation of Y 3 Al 5 O 12 phase as isolated pockets at corners of Al 2 O 3 grains, which can benefit the decrease of phonon scattering caused by the second phase. 5) Some fin-type Al 2 O 3 ceramic heat sinks were successfully fabricated without obvious geometry deformation via the CP process followed by DLPstereolithography, which may offer a new opportunity for thermal management applications of Al 2 O 3 ceramic.