Evolution of the microstructure and mechanical properties of stereolithography formed alumina cores sintered in vacuum

Journal Pre-proof Evolution of the microstructure and mechanical properties of stereolithography formed alumina cores sintered in vacuum He Li, Yongsheng Liu, Yansong Liu, Qingfeng Zeng, Jing Wang, Kehui Hu, Zhigang Lu, Jingjing Liang

PII:

S0955-2219(19)30792-7

DOI:

https://doi.org/10.1016/j.jeurceramsoc.2019.11.047

Reference:

JECS 12872

To appear in:

Journal of the European Ceramic Society

Received Date:

24 October 2019

Revised Date:

13 November 2019

Accepted Date:

14 November 2019

Please cite this article as: Li H, Liu Y, Liu Y, Zeng Q, Wang J, Hu K, Lu Z, Liang J, Evolution of the microstructure and mechanical properties of stereolithography formed alumina cores sintered in vacuum, Journal of the European Ceramic Society (2019), doi: https://doi.org/10.1016/j.jeurceramsoc.2019.11.047

This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier.

Evolution of the microstructure and mechanical properties of stereolithography formed alumina cores sintered in vacuum

He Lia,b, Yongsheng Liua,b* [email protected], Yansong Liua,b, Qingfeng Zenga* [email protected], Jing Wanga,b, Kehui Huc, d, Zhigang Luc, d, Jingjing Liange

a

Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University,

Xi’an, Shaanxi 710072, China NPU-SAS Joint Research Center of Advanced Ceramics, Northwestern Polytechnical University, Xi’an, Shaanxi 710072,

ro of

b

China

Department of Mechanical Engineering, Tsinghua University, Beijing 100084, China

Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China

*Corresponding

author. Tel.: +86 029 88495179; Fax: +86 029 88494620

re

e

State Key Laboratory of Tribology, Tsinghua University, Beijing 100084, China

-p

d

lP

c

Abstract

na

Stereolithography (SL) was used to form alumina ceramic cores. The effect of sintering temperature on the microstructure and mechanical properties of the alumina ceramics are investigated, which were sintered in vacuum. The results indicate

ur

that, as the sintering temperature increased the particle size of alumina slightly increased, and the interlayer spacing first decreased and then increased. The open porosity of alumina ceramics significantly decreased as the sintering temperature in

Jo

vacuum increased. The flexural strength, hardness and modulus increased as the sintering temperature increased. When sintered at 1150 °C, the flexural strength was found to be 33.7 MPa, the shrinkage was 2.3%, 2.4%, and 5.3% in the X, Y, and Z directions, respectively, and the open porosity was 37.9%. These results are similar to those found from sintering at 1280 °C in air.

Keywords: stereolithography; aluminum oxide; sintering temperature; vacuum; ceramic cores

1. Introduction Ceramic cores with appropriate flexural strength, open porosity, and high temperature stability have been used for superalloy hollow blades with complex internal structures [1]. Alumina ceramics have good metallurgical chemical stability

and creep resistance, which can guarantee the dimensional accuracy of blades with complex internal cavity structures. Therefore, this material is suitable for utilization in advanced single crystal and eutectic casting conditions [2-4]. Generally, an alumina ceramic core must satisfy the following requirements [5]: (1) moderate strength (20 to 30 MPa) in order to resist mechanical and thermal damage in high temperature alloy casting and solidification process; (2) chemical stability such as that the core does not contaminate the alloy or react with molten metal during the contact process; (3) the thermal expansion rate of the core should be as small as possible to ensure the inner cavity has the right size during casting, and the ceramic core will not be damaged or destroyed due to a phase change; (4) refractoriness of the core should be higher than the casting temperature of the alloy to ensure that there is no softening or deformation during casting and solidification; (5) the alumina should be easy to remove. Ceramic cores are usually prepared though investment casting method using silica or alumina based materials for their

ro of

desired thermal properties, lack of reaction with the superalloy, and low thermal expansion [6]. It is known that the preparation of ceramic cores though investment casting method is time-consuming, expensive, and low precision. Therefore, the development of a fast and inexpensive method to prepare ceramic cores is of interest.

-p

Stereolithography (SL) is one of the popular 3D printing methods. Bae et al. [7] fabricated a ceramic investment casting mold using SL, and the sintering shrinkage was 10.7%. Zanchetta et al. [8] fabricated complex 3D polymer-derived ceramic

re

structures using SL and fabricated dense, crack-free silicon oxycarbide-based microparts. This technology offers an efficient way to fabricate complex ceramic parts and with low time and cost. Moreover, it provides a new method for forming

lP

ceramics. However, printed green bodies contain a large number of photosensitive resins that must be removed during the debinding and sintering processes. Removing the photosensitive resins will lead to shrinkage of more than 10% and deformation of ceramic parts, which limits its application in SL during fabrication of ceramic cores due to the high precision

na

and mechanical properties of ceramic core [9]. Several efforts have focused on increasing the inorganic powder content in the ceramic slurry in order to reduce shrinkage and deformation via cumbersome preparation procedures. Li et al. [10] tried

ur

to use surfactants to prepare alumina suspensions, and it was shown that oleic acid provides the best flexural strength. 40 vol% alumina suspensions were obtained and a sintering density of 95% can be reached. Zhang et al. [11] managed to prepare

Jo

alumina slurries (60 vol%) with 5 wt% KOS110 and high solid loading. SL is a relatively new method for production of alumina ceramics [12]. Currently, the preparation of ceramic slurries, printing of green bodies, debinding, and sintering procedures of green bodies all require further development. According to Xing et al. [13], cracks would form between two adjacent layers in alumina ceramics formed via SL. Prevention of interlayer delamination in the green and sintered bodies is one of the critical issues that deserves attention. Zhou et al. [14] found that residual carbon in the body would lead to crack formation in the body during sintering, but sintering in vacuum would suppress defect formation in the alumina body. In addition to crack formation during sintering, large shrinkage is also harmful to the ceramic cores. According to Gonzalez et al. [15], when sintered at 1600 °C, the expected shrinkage of the ceramic materials fabricated with SL was 20% ~ 25% in the X or Y directions, and 25% ~ 30% in the Z direction. Hu et al. [16] fabricated ceramics from ceramic suspensions with high

solid loading with the shrinkage of 12.5%. Li et al.[17] found that 1280 °C was the best sintering temperature sintered in air atmosphere. Sun et al. [18] found that alumina ceramic sintered in vacuum could have fewer residual pores. To prevent the crack formation, shrinkage should be controlled, and a vacuum sintering process was proposed to improve the quality of the alumina ceramic fabricated with SL. In this study, we attempted sintering in vacuum in an effort to control shrinkage and deformation of fabricated alumina ceramics. In addition, the evolution of the microstructure and mechanical properties of alumina ceramics obtained by SL were investigated. Apart from ceramic cores, the sintered method could also be applied in other ceramic parts. 2. Experimental 2.1. 3D printer, raw materials, and green body fabrication Green bodies were provided by Tsinghua University, and a detailed preparation process can be found in another paper [16]

ro of

. Fig. S1 shows the 3D printer (AutoceraM, Beijing Ten Dimensions Technology Co., Ltd.) used in this study. The

irradiation source was a 405 nm LED. The particle size distribution of the alumina powders (AW-SF, Henan Hecheng Inorganic New Material Co., Ltd.) is shown in Fig. S2. The alumina powder was dried at 200 °C for 5 h in a blast air oven

-p

and was subsequently used for ceramic slurry preparation. 495 g of Al2O3 powder was slowly added to 100 g of a

photosensitive resin (Al100-1, Beijing Ten Dimensions Technology Co., Ltd.); the mixture was stirred vigorously to obtain

re

the ceramic slurry. After the slurry was stirred evenly, it was ball-milled for 2 h with a planetary grinding mill at 400 rpm. Then the slurry was vacuum defoamed for 10 min to obtain the final alumina ceramic slurry. A cuboid model (50 mm×4

lP

mm×3 mm) was drawn using UG software and imported into STL format. The model file was transferred to the 3D printer. The ceramic slurry was then transport to the 3D printer. The exposure energy was 10 mW/cm2, and the single layer exposure

na

time was 10 s. The green bodies were printed with 0.1 mm layer thickness. The excess slurry was cleaned with ethanol, yielding the green bodies shown in Fig. S3.

ur

2.2. Debinding and sintering process

The green bodies were debound in a muffle furnace (Hefei Ke Jing Materials Technology Co., Ltd. China). First, the

Jo

green bodies were heated to 200 °C at a heating rate of 2 °C/min. Second, the samples were heated to 550 °C at a heating rate of 1 °C/min and sustained for 2 h. Third, the samples were heated to 1000 °C at a heating rate of 5 °C/min and sustained for 2 h. The temperature was allowed to cool to 600 °C at a heating rate of 5 °C/min, and the temperature of the samples subsequently dropped to room temperature as the furnace cooled. The samples were then transferred to a vacuum hot-pressing furnace (ZT-40-207, Shanghai Guanbo Electric Furnace Co., Ltd. China), which is shown in Fig. S4. After the vacuum pressure dropped to 10-3 Pa, the samples were heated to 800 °C at a heating rate of 10 °C/min. Then the samples were heated to the target temperature (T = 1100, 1150, 1200, 1250, 1300, or 1350 °C) at a heating rate of 5 °C/min and sustained for 2 h. The debinding and sintering profiles are illustrated in Fig. 1. The samples were denoted as S(T).

Fig. 1. Debinding and sintering of green bodies: (a) debinding; (b) sintering.

ro of

2.3. Sample characterization Powder X-ray diffraction (XRD) patterns were recorded with a Bruker D8 FOCUS (Bruker Corporation, Germany) Xray diffractometer using a Cu Kα radiation source. The 2θ diffraction angle was scanned in 0.02°s-1 increments from 10° to 90°. The voltage was set to 40 kV and the generator current was 30 mA. Raman spectra were gathered from the samples

-p

with a confocal Raman microscope (Alpha300R, WITec) with a 532 nm laser at 50 mW and 50 μm slit width. X-ray photoelectron spectroscopy (XPS) was performed using a photoelectron spectrometer (Axis Supra, Shimadzu). Scanning

re

electron microscope (SEM, Helios G4 CX, FEI Corporation) and transmission electron microscope (TEM, Tecnai G2 F20, FEI Corporation) images were gathered to study the microstructure of the ceramics. Energy dispersive spectroscopy (EDS)

lP

was also performed. A double Cs corrector TEM (Themis Z, FEI, USA) was also used to examine the micromorphology of the samples.

na

The open porosity and bulk density of the samples after sintering process was measured with the Archimedes method. The accuracy of the balance was 0.0001 g (Mettler Toledo, Switzerland).

ur

The flexural strength of the sintered samples was tested with an electronic universal testing machine (CMT4304, SUNS, China) using the three-point flexure method. The loading speed was 0.5 mm/min and the span was 30 mm.

Jo

The surface roughness of the sample was measured with an atomic force microscope (Dimension Icon, Bruker, USA), and the spring constant was 0.40 N/m. Nanoindentation tests were performed with a nano-test apparatus (TI980, Hysitron, USA). Static indentation was applied at room temperature. The cross-sectional area of an indentation is shown in Fig. 2, which illustrates the relationship between F, A, hc, and h. First, the head approached the surface of the sample at 2 mN/s, and the sample was loaded to a maximum load of 10 mN over 5 s. The sample was unloaded over a period of 2 s. A load-displacement curve was recorded on the connected computer during the test, and 5 points on each sample were tested. The hardness and modulus of the materials were calculated with the following equations:

FMax A

H

Er

(1)

S A

2

(2)

1 v 2 1 vi 2 + E Ei

1 Er S

dF dh

(3)

(4) h hMax

F B(h h f ) m hc

h

(5)

FMax S

(6)

ro of

A 24.5hc2

(7)

A 24.5hc2 102.25hc is the hardness,

contact stiffness,

FMax

is the maximum load,

is the conversion factor,

E

vi

hc

is the

is Poisson's ratio of the material,

is Poisson's ratio of the indenter (0.07),

is the displacement after complete unloading,

is the contact depth, and

B

and

m

Er

is

are

is a constant

Jo

ur

na

related to the shape of the indenter (0.75).

v

S

lP

hf

is the projected contact area at maximum load,

is the elasticity modulus,

the elastic modulus of the diamond indenter (1140 GPa), fitting parameters,

A

-p

H

re

where

(8)

Fig. 2. Nanoindentation schematic. A Vickers hardness test was performed with a micro/macro automatic hardness testing unit (LM248AT, LECO, USA). The load was 1000 g, the dwell was 15 s, and 5 points on each sample were tested. The pore diameter distribution was measured with a mercury porosimeter (AutoPore IV 9500, Micromeritics, USA). 3. Results and discussion

3.1. Microstructure and composition After sintering at different temperatures in vacuum, delamination in all the samples was clearly seen in the SEM images, which are shown in Fig. 3. The interlayer spacing (shown in Fig. 4) first decreases and then increases as the sintering temperature increases. The interlayer spacing reached a minimum value of 5.94 μm after sintering at 1200 °C. Compared with the microstructure of the sintered samples observed by Chen et al. [19], no cracks were found here, which may be due to the slow heating rate and vacuum atmosphere promote junction formation between Al2O3 particles. According to the description from Chen et al. [19], the sintering process could heal some small cracks. Although delamination between layers occurred less often due to shrinkage during sintering, delamination phenomenon could not be eliminated. It is difficult to prevent delamination between layers by regulating the sintering temperature. Results from ceramics prepared with SL by Zhou et al. [20] show that the thickness of a single layer should be larger than that of sliced

ro of

layers to ensure continuous bonding between layers and obtain a three-dimensional solid. Furthermore, layers in the X- Y plane are exposed to a homogeneous planar light source, and joining occurs through uniform polymerization in the X-Y plane. Interfacial joining between layers is caused by exposure along the Z direction, and the energy distribution may be

-p

different from that in the X-Y plane. Interfacial polymerization can be different from that in the X-Y plane, which may attract more photopolymers [21]. Therefore, layer delamination of the printed bodies is reflected in the sintered samples.

re

Meanwhile, sintering was conducted in vacuum, and relative pressure is generated in the samples at high temperature [22], which promotes bonding between particles. The combination of heat and pressure increase the plasticity and fluidity of the

lP

alumina particles, which determine the interlayer spacing at various sintering temperature. When sintered at 1100 °C, a relatively large interlayer spacing between layers was created due to resin volatilization. When the temperature was increased to 1150°C, the sintering driving force increases and promotes bonding between layers, so the spacing between

na

layers decreases. However, when the sintering temperature is further increased to 1300°C, the sample shrinks seriously

Jo

ur

between layers, so the interlayer spacing increases.

Fig. 3. SEM images showing delamination during sintering at (a) 1100, (b) 1150, (c) 1200, (d) 1250, (e) 1300, and (f)

ro of

1350 °C.

Fig. 4. Variation in interlayer spacing in samples sintered at different temperatures.

Fig. S5 shows a magnified view of the samples in Fig. 3, illustrating the dispersion and combination of particles in the samples sintered at different temperatures. The SEM images indicate that the combination of particles in the sample after

-p

sintering is rather weak, and small particles are randomly scattered and attach to each other around large particles. Meanwhile, the intergranular bond is more compact as the sintering temperature increases, and the intergranular pore

re

decreases gradually. According to the two-particle model put forward by Kuczynski [23], increasing the sintering temperature will accelerate the neck growth rate of spherical particles, which promotes intergranular binding. The neck growth rate is

lP

defined as follows:

rM P0 3 2

3 2

2R T d

)

1 3

2 1 3 3

r t

(9)

2

x is the neck growth rate of the spherical particle contact area, T is the sintering temperature, and P0 is the r

vapor pressure.

ur

where

(

3

na

x r

3 2

When the sintering temperature increases, the vapor pressure increases exponentially and the growth rate accelerates,

Jo

which densifies the material during sintering. This observation is consistent with the results shown in Fig. 5, where particles are more likely to combine at higher sintering temperature, while the particles remain independently dispersed at lower sintering temperature.

The average particle size in each sample was measured, and the results are shown in Fig. S6. This figure shows that the particle size increases as the sintering temperature increases. When the sintering temperature increases, the sintering driving force increases, which causes grain size to increase. According to the crystal growth model put forward by Hillert [24], the relationship between sintering temperature and crystallite size is defined as follows:

D2 D02 tC exp[ Q / ( RT )]

(10)

where D is the crystallite size, C is a constant, and Q is the activation energy for grain growth. One can see from this equation that the crystallite size is increases as the sintering temperature increases, which is consistent with the conclusion in Fig. S6. According to the densification and grain growth of Al2O3 nanoceramics during pressureless sintering [25], the average grain size is known to increase as the sintering temperature increases. This phenomenon is similar to what is seen in 3D printed green bodies in this case. Results from sintered nanosized α-Al2O3 powder by Karagedov et al. [26] show that densification during the initial sintering stages was attributed to dense packing of small crystallites, especially within the formed aggregates. These aggregates were considered an obstacle to form high-density ceramics. Stepped growth of alumina was evident when the sample was magnified further, which is shown in Fig. 5. There are two aspects should be noticed. First, as the sintering temperature increases, the bond between particles becomes more

ro of

compact, the size of voids between particles decreases. When sintered at relatively lower temperatures (e.g., 1100 and 1150 °C), there is a large number of voids between particles. However, when sintered at relatively higher temperatures, such as 1300 and 1350 °C, most of the particles become interconnected. Second, large particles were found to form ledges.

-p

According to the growth of α-Al2O3 within a transition alumina matrix researched by Morrissey et al. [27], larger particles have curved edges that can extend into the fine-grain matrix. These α-Al2O3 grains frequently overlap with smaller grains in

re

the transition alumina matrix, and this growth is accompanied by a thickening of the α phase. The surface steps are significant, not only for their part in transformation, but also because their thickness is uniform. Gao et al. [28] found that

lP

stepped growth of alumina is a common phenomenon in sintered alumina. According to the Kossel-Stranski crystal-growth model [29], steps form after two-dimensional nucleation and subsequently move. The steps move in parallel and eventually disappear, so only a few particles can be observed in Fig. 5. Chen et al. [30] found that the vacuum atmosphere could result

na

the more uniform distributed particles compared to the normal pressures. This is because that the driving force in vacuum during sintering process was relatively higher than normal pressures, especially for the period after the pores are formed.

Jo

particles.

ur

The small particles could grow fast due to the high driving force during sintering, then lead to the uniform distributed

ro of

Fig. 5. SEM images (10000× magnification) of samples sintered at (a) 1100, (b) 1150, (c) 1200, (d) 1250, (e) 1300, and (f) 1350 °C.

TEM images of the micromorphology are shown in Fig. 6. Large irregularly-shaped particles in the figure may be composed of a number of small particles that have been sintered together at high temperatures. This is because clear edges

-p

and lines can be seen in the samples sintered at 1150 or 1200 °C. According to Montero et al. [31], the boundaries between particles will be difficult to distinguish in a sample with an excessively thick α-Al2O3 support. If some particles are tightly

re

bound together during sintering, the interface between different particles may not be observable. Gupta et al. [32] found that alumina particles area mostly irregular and spherical, and that few particles are agglomerated during synthesis of Al2O3

lP

nanofluids. HRTEM images are shown in Fig. S7 and selected area electron diffraction (SAED) images of the sintered samples are shown in Fig. S8. The clear and regular lattice in Fig. S7 shows that oxygen and aluminum atoms are regularly

na

arranged in the sintered samples. The lattice spacing was measured to be 0.25, 0.34, 0.24, 0.21, 0.20, and 0.24 nm from the HRTEM images which were sintered at 1150, 1200, 1250, 1300, and 1350 °C, respectively. Combined with the XRD results in Fig. 8, one can conclude that the lattice planes are (1 0 4), (0 1 2), (1 1 0), (1 1 3), (2 0 2), and (1 1 0), respectively. The

ur

distribution of Al, O, and C are shown in Figs. S9-S11, respectively, which show the distribution of Al, O, and C are uniform. There seems no obvious distinction in the elemental distribution in samples sintered at different temperatures. The

Jo

EDS curves shown in Fig. S12 confirm the particles are composed of O and Al. The atomic fraction obtained from TEM is shown in Table 1. Only very low levels of C (<2%) are present detected. The accuracy of EDS elemental analysis is usually less than 3%, so the presence of C is negligible. Sample of S(1150) was prepared using a focused ion beam system, and the sample was observed with the double Cs corrector TEM. The results are shown in Fig. 7. The macroscopic morphology of these samples in Fig. 7a illustrate the porosity of the 3D printed samples. From the HRTEM and atomic-resolution STEM images, no dislocation could be found in the samples. According to the nonstoichiometric dislocation cores in α-alumina researched by Shibata et al. [33], the implication for mechanical properties is that the mobile high-temperature dislocation core structures consist of two closely-

spaced partial dislocations. Table 1. Atomic fraction determined from TEM C (%)

O (%)

Al (%)

S(1100)

0.62

52.08

47.30

S(1150)

0.46

46.58

52.96

S(1200)

0.95

55.22

43.83

S(1250)

2.23

60.14

37.63

S(1300)

0.94

48.48

50.58

S(1350)

0.87

46.40

52.73

lP

re

-p

ro of

Sample ID

Jo

ur

na

Fig. 6. TEM images of samples sintered at (a) 1100, (b) 1150, (c) 1200, (d) 1250, (e) 1300, and (f) 1350 °C.

Fig. 7. TEM images of S(1150): (a) macroscopic morphology in TEM samples, (b) HRTEM images, (c) atomicresolution STEM images, and (d) SAED patterns.

The phase composition was determined using XRD patterns from the samples sintered at different temperatures in vacuum; the XRD patterns are shown in Fig. 8. The XRD patterns indicate that all samples are α-alumina. The peaks of the XRD curves are located at 25.6°, 35.1°, 37.8°, 41.7°, 43.4°, 46.2°, 52.6°, 57.5°, 59.8°, 61.2°, 66.5°, 68.2°, 70.4°, 74.3°,

-p

ro of

77.2°, 80.7°, 84.4°, and 86.4°.

Fig. 8. XRD patterns from samples sintered at different temperatures in vacuum.

Raman spectra from the samples sintered at different temperatures in vacuum are shown in Fig. 9. These results show

re

that all samples have the same peak, meaning they have the same chemical bond structure. The peaks in Raman spectra are located at 378, 418, 578, 645, 751 cm-1. The Raman spectra of synthetic sapphire (α-Al2O3) reported by Kadleikova et al. [34]

lP

is expected to have 7 bands with peaks at 378, 418, 432, 451, 578, 645, and 751 cm-1, respectively. Compared with the Raman spectrum obtained from the alumina ceramics sintered at different temperatures in vacuum, most of the peaks are at

na

the same locations. XPS data from the samples sintered at different temperatures in vacuum are shown in Fig. 10. The peaks at 529, 282, 117, and 72 eV correspond to the O 1s, C 1s, Al 2s, and Al 2p levels, respectively. The peaks in the alumina

ur

ceramics sintered at different temperatures are identical, indicating that the chemical bonds in different samples are identical. The sintering temperature in vacuum does not affect the chemical bond structure. Atomic ratios acquired from

Jo

XPS in samples sintered at different temperatures are shown in Table 2. These data indicate that if the samples were etched

for 1 min at 5 keV before the measurement, the content of C would be reduced to some extent. This proves that samples placed in air will absorb CO2 and other substances, producing errors in measurements from the sample surface, thus etching is required before the measurement.

Jo

ur

na

lP

re

-p

ro of

Fig. 9. Raman spectra from samples sintered at different temperatures in vacuum.

Fig. 10. XPS curves from samples sintered at different temperatures in vacuum: (a) full spectrum; (b) O 1s; (c) Al 2p; (3) C 1s. Table 2. Atomic ratio (XPS) in different sintered samples (etch: samples etched for 1 min at 5 keV before measurement) Sample ID

O

C

Al

S(1100)

47.92

14.58

37.50

S(1100) etch

52.81

2.43

44.76

S(1150)

48.47

14.23

37.30

S(1150) etch

52.51

3.22

44.27

S(1200)

52.46

28.15

19.38

S(1200) etch

57.82

7.93

34.25

S(1250)

48.81

21.63

29.56

S(1250) etch

53.35

5.48

41.18

S(1300)

48.87

13.62

37.52

S(1300) etch

49.98

11.22

38.80

S(1350)

48.85

12.90

38.25

S(1350) etch

52.77

2.54

44.69

3.2. Physical properties Shrinkage in samples sintered at different temperatures in vacuum is shown in Fig. 11. Shrinkage along the X, Y, and Z

ro of

directions increased from 1.6% to 6.5%, from 1.8% to 6.6%, and from 3.5% to 11.8%, respectively, as the sintering temperature increasing from 1100 to 1350 °C. These results show that shrinkage increases as the sintering temperature increases. Shrinkage is also different along three directions: shrinkage along the X direction is nearly the same as that along the Y direction, while the shrinkage in the Z direction is much larger. From the sintering model put forward by Kingery [35],

where

(

6

2

Dc0 LV V0 13 43 13 ) r t 1 RT

(11)

re

L L

-p

particle rearrangement occurs during the middle and later stages of sintering. Shrinkage is defined as follows:

L is the shrinkage in the sintered samples and T is the sintering temperature. L

which reflects the trend shown in Fig. 11.

lP

Kingery’s sintering model illustrates that shrinkage in sintered samples increases as the sinter temperature increases,

na

Obviously, shrinkage along the Z direction is much larger than that along the X and Y directions, with the ratio in shrinkage close to 2. This indicates that shrinkage of 3D printed alumina green bodies is not uniform after sintering. Layerby-layer deposition during 3D printing and the resulting shrinkage are quite different along the Z direction and the X/Y

ur

directions. Delamination also occurs in dense 3D printed ceramics [36]. The microstructure shown in Fig. 3 reveals gaps

Jo

between layers in the sintered samples. Weak binding between layers along the Z direction may lead to larger shrinkage compared with that along the X and Y directions. Dehurtevent et al. [37] also found that an alumina ceramic fabricated using SL exhibit anisotropic shrinkage. Although shrinkage can be compensated by using a magnified design model, serious deformation will occur when the structure is complex, such as an alumina ceramic core. Braga et al. [38] also found that higher shrinkage could cause more cracks and deformation in ceramics. According to Li et al. [39], deformation occurs due to uneven shrinkage during sintering. Shrinkage could be reduced by increasing the powder content in the slurry and decreasing the heating rate during sintering. However, increasing the powder content in the slurry may be rather difficult due to its rather high viscosity and low fluidity. While the anisotropic shrinkage was dangerous for product integrity after sintering, and also will result the deformation of the ceramic parts, so decrease the shrinkage was rather important for the

stereolithography formed ceramics parts. As the shrinkage was increasing as the increasing sintering temperature in vacuum atmosphere, the relatively lower sintering temperature would lead to the lower shrinkage and lower deformation for the final

Fig. 11. Shrinkage in samples sintered at different temperatures in vacuum.

ro of

ceramic parts.

-p

The bulk density and open porosity in the samples sintered at different temperatures in vacuum are shown in Fig. 12. The bulk density increased from 2.4 to 2.9 g/cm3 and open porosity decreased from 38.7% to 23.8% as the sintered

re

temperature was increased from 1100 to 1350 °C. A schematic diagram showing possible structural changes in the samples during sintering is shown in Fig. 13. Fig. S6 shows that the alumina particle size gradually increases as the sintering

lP

temperature increases. Although the number of pores in the sample decreases, the contact area between larger particles increases, which reduces the size of pores in the samples. Increasing the sintering temperature leads to decreased open

na

porosity. Fig. 11 indicates that shrinkage increases as the sintering temperature increases, thus the volume decreases significantly as the sintering temperature increases. Therefore, the bulk density increases as the sintering temperature increases due to volume contraction during sintering at high temperature. Increased powder loading ensures the green body

ur

obtain higher density [40], thus sintering produced low porosity sintered parts. However, during preparation of ceramic cores, the bulk density should not be too high and there should be adequate porosity to ease removal. Ceramic cores require open

Jo

porosity of >20% to ensure air permeability during investment casting and liquid permeability during removal [41]. Compared with the data in Fig. 12, the open porosity is greater than 20% when the sintering temperature ranges from 1100 and 1350 °C, which means the sintering temperatures used in this study provide the desired open porosity. The pore diameter distribution obtained from mercury porosimeter is shown in Fig. 14 and Table 3. Fig. 14 shows that the sintered samples are composed of alumina particles with different particle sizes, which is consistent with the SEM results in Fig. 5. Table 3 shows that the median pore diameter decreases as the sintering temperature increases. This observation is consistent with the theoretical assumption in Fig. 13.

lP

re

-p

ro of

Fig. 12. Bulk density and open porosity in samples sintered at different temperatures in vacuum.

Jo

ur

na

Fig. 13. Schematic diagram showing structural changes during sintering process.

Fig. 14. Pore diameter distribution obtained from mercury porosimeter.

Table 3. Pore diameter distribution obtained from mercury porosimeter.

Sample ID

Total intrusion volume (mL/g)

Total pore area (m2/g)

Median pore diameter (volume)(nm)

Median pore diameter (area)(nm)

Average pore diameter (4V/A)(nm)

S(1100)

0.14

1.76

324.90

307.31

317.90

S(1150)

0.14

1.62

354.29

325.20

353.72

S(1200)

0.12

1.24

399.33

371.58

397.06

S(1250)

0.12

31.76

395.25

4.22

15.32

S(1300)

0.11

24.29

393.97

6.48

18.47

S(1350)

0.12

24.37

415.40

4.03

18.91

ro of

3.3. Mechanical properties The flexural strength, Vickers hardness, load-displacement curves, and nanoidentation hardness are shown in Fig. 16. The flexural strength increased from 20.3 to 138.9 MPa as the sintering temperature increased from 1100 to 1350 °C. According to Ryskewitsch empirical equation below [42], flexural strength of the ceramic decreases as the porosity increases:

where

exp(

p)

is the strength,

(12) 0

is the strength at zero porosity,

p

-p

0

is the porosity, and

is a constant.

re

The experimental data obtained in this study were fit to Eq. (14); this curve is shown in Fig. 15, where the porosity

lP

should be larger than 20% and lower than 40%. The fitting curve is consistent with the data, and the difference between the experimental data and the theoretical prediction is small. When sintered in vacuum, the open porosity decreased as the sintering temperature increased, as shown in Fig. 12. Eq. (14) indicates the flexural strength should increase as the sintering

Jo

ur

na

temperature increases. The values from the fitting curve are consistent with the experimental results shown in Fig. 16(a).

Fig. 15. Fit to the Ryskewitsch empirical equation. From Jiang et al. [43], the ceramic core must be strong enough to be resistant to thermal shock and sufficiently resistant to dimensional changes. The flexural strength of the ceramic material they prepared ranged from 10.0 and 24.8 MPa.

Generally, the flexural strength of the ceramic core should be higher than 20 MPa and lower than 50 MPa to prevent striking damage during high temperature alloy casting and solidification. From the data shown in Fig. 16, one can see that the flexural strength was 20.3 MPa when sintered at 1100 °C and 33.7 MPa when sintered at 1150 °C. In order to ensure the ceramic cores had sufficient strength, the sintering temperature should be greater than 1150 °C. The delamination results shown in Figs. 3 and 4, the sintering temperature should be between 1150 and 1250 °C in order to ensure low spacing between adjacent layers. Because the structures of the ceramic cores are very complex, smaller shrinkage will reduce deformation in the structure. Fig. 11 shows that shrinkage increased as the sintering temperature increased. Therefore, 1150 °C was considered to be the best sintering temperature. In summary, when sintered at 1150 °C in vacuum, the obtained ceramic is suitable for preparation of ceramic cores. The interlayer spacing distance was 6.44 μm, the shrinkage was 2.3%, 2.4%, and 5.3% along the X, Y, and Z directions, respectively, the bulk density was 2.43 g/cm3, the open porosity was

ro of

37.9%, and the flexural strength was 33.7 MPa.

The Vickers hardness and nano-indentation tests were used to explain the change in material hardness as the sintering temperature changed. All test results show that the hardness increased as the sintering temperature increased. Manshor et al. shows that the hardness of a zirconia-toughened alumina ceramic composite increased due as the bulk density increased.

-p

[44]

This means that denser materials have higher hardness, which is in line with the results obtained from Fig. 12. Nano-

re

indentation tests were conducted to understand the nanomechanical properties of the sintered alumina ceramics. The variation in hardness follows the same trend as the Vickers hardness. According to Xing [45], large differences in instability

lP

values were observed during nano-indentation tests due to the porosity and non-compactness of the alumina ceramics. In order to ensure the data were reliable, we used an AFM to measure the surface roughness of the sample; the results are

Jo

ur

na

shown in Figs. 17 and S13, and in Table 4.

ro of

-p

Fig. 16. Mechanical properties of the samples sintered at different temperatures in vacuum: (a) flexural properties, (b) Vickers hardness, (c) load-displacement curves from nano-indentation tests, and (d) hardness obtained from nano-

Jo

ur

na

lP

re

indentation tests.

Fig. 17. A 3D diagram of the sample surface obtained from AFM for samples sintered at (a) 1100, (b) 1150, (c) 1200, (d) 1250, (e) 1300, and (f) 1350 °C.

Table 4. Surface roughness measured with an atomic force microscope (AFM). Sample ID S(1100) S(1150) S(1200) S(1250) S(1300)

94.6±2.6 97.2±3.6 91.3±1.8 95.0±5.1 90.7±4.0 96.5±5.5

ro of

S(1350)

Surface roughness Rq (nm)

Green bodies undergo contraction during sintering; pores disappear and the material density increases to form the ceramic. The driving force in sintering is the reduction in the surface energy of the powders in the green body. The sintering process is driven by reduction of the total boundary area caused by the substitution of a low energy grain boundary for a

-p

high energy grain surface and contraction of the green body. Gustavo et al. [46] found that vacuum sintering transparent alumina at low pressure slows down the sintering rate, which accelerated grain growth and intragranular porosity. Thus,

re

alumina compacts with very low porosity and grain size can be obtained by sintering at the temperature corresponding to

lP

maximum shrinkage. Srdic et al. [47] observed fast densification in nanocrystalline zirconia doped with alumina at very low temperature, which can be explained due to vacuum sintering, and the dispersed powder has very small crystallite size and narrow size distribution. Vacuum sintering can remove photosensitive resin and eliminate gas in the green bodies, and

na

densification occurs at a faster rate. High bulk density and flexural strength could be reached at relative low temperature. As shown in Fig. 18, gas produced by photosensitive resin volatilization is quickly removed in vacuum, and densification

ur

occurs at a faster rate. In the meantime, the high vacuum pressure between particles reduces bonding between particles. Therefore, sintering in vacuum can be used to produce ceramics with higher flexural strength at lower temperature.

Jo

Combing Kuczynski’s double ball model [48] shown in Fig. 19, the particle migration velocity is equal to the neck volume growth when the sintering temperature increases gradually. The particles gradually grew up and rearranged, while particles in vacuum would promote particle rearrangement over a shorter period of time. Compared with sintering in air, vacuum sintering could proceed at lower temperature while producing a ceramic with equal strength. Sintering in air at 1280 °C and sintering in vacuum at 1150 °C would produce ceramics with the same flexural strength.

ro of

lP

re

-p

Fig. 18. Schematic diagram showing the vacuum sintering process.

Fig. 19. Kuczynski’s double ball model.

na

4. Conclusions

The influence of sintering temperature on the microstructure, physical properties, and mechanical properties of alumina

ur

ceramics that were sintered in vacuum. The interlayer spacing first decreased and then increased, and the particle size increased, as the sintering temperature increased. Sintering temperature would not influence the phase composition and

Jo

chemical bond structure. Shrinkage and bulk density were found to increase as the sintering temperature increased, while open porosity was found to decrease as the sintering temperature increased. The sintering temperature influences the amount and size of the pores. Flexural strength, Vickers hardness, and nano-indentation test results show that the flexural strength and hardness increased as the sintering temperature increased because the bulk density increased. Overall, 1150 °C was determined to be the best sintering temperature in vacuum, yielding a ceramic with shrinkage of 2.3%, 2.4%, and 5.3% in the X, Y, and Z directions, respectively, bulk density of 2.43 g/cm3, open porosity of 37.9%, and flexural strength of 33.7 MPa.

Declaration of interests

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

The authors declare the following financial interests/personal relationships which may be considered as potential

ro of

competing interests:

Acknowledgements

This work was supported by National Key Research and Development Program of China (No. 2018YFB1106600), the

Jo

ur

na

lP

re

-p

Chinese National Foundation for Natural Sciences under Contracts (No. 51672217).

Reference [1] Zhao H, Ye C, Fan Z, et al. 3D printing of CaO-based ceramic core using nanozirconia suspension as a binder, J. Eur. Ceram. Soc. 37 (2017) 5119-5125. [2] Wereszczak A A, Breder K, Ferber M K, et al. Dimensional changes and creep of silica core ceramics used in investment casting of superalloys, J. Mater. Sci. 37 (2002) 4235-4245. [3] Liu F, Fan Z, Liu X, et al. Aqueous gel casting of water-soluble calcia-based ceramic core for investment casting using epoxy resin as a binder, Int. J. Adv. Manuf. Tech. 86 (2016) 1235-1242. [4] Pollock T M. Alloy design for aircraft engines, Nat. Mater. 15 (2016) 809.

ro of

[5] Wang B S. Core-removing equipment and process of alumina-based ceramic core, Xian: Northwestern Polytechnical University, 2005.

[6] Lu Z L, Fan Y X, Miao K, et al. Effects of adding aluminum oxide or zirconium oxide fibers on ceramic molds for casting hollow turbine blades, Int. J. Adv. Manuf. Tech. 72 (2014) 873-880.

-p

[7] Bae C J, Halloran J W. Integrally cored ceramic mold fabricated by ceramic stereolithography, Int. J. Appl. Ceram. Tec. 8 (2011) 1255-1262.

re

[8] Zanchetta E, Cattaldo M, Franchin G, et al. Stereolithography of SiOC ceramic microcomponents, Adv. Mater. 28 (2016) 370-376.

lP

[9] Johansson E, Lidström O, Johansson J, et al. Influence of resin composition on the defect formation in alumina manufactured by stereolithography, Materials. 10 (2017) 138.

(2017) 4761-4767.

na

[10] Li K, Zhao Z. The effect of the surfactants on the formulation of UV-curable SLA alumina suspension, Ceram. Int. 43

[11] Zhang K, Xie C, Wang G, et al. High solid loading, low viscosity photosensitive Al2O3 slurry for stereolithography

ur

based additive manufacturing[J]. Ceram. Int. 45 (2019) 203-208. [12] Azarmi F, Amiri A. Microstructural evolution during fabrication of alumina via laser stereolithography technique,

Jo

Ceram. Int.45 (2019) 271-278.

[13] Xing Z, Liu W, Chen Y, et al. Effect of plasticizer on the fabrication and properties of alumina ceramic by stereolithography-based additive manufacturing, Ceram. Int. 44 (2018) 19939-19944. [14] Zhou M, Liu W, Wu H, et al. Preparation of a defect-free alumina cutting tool via additive manufacturing based on stereolithography–Optimization of the drying and debinding processes, Ceram. Int. 42 (2016) 11598-11602. [15] Gonzalez P, Schwarzer E, Scheithauer U, et al. Additive Manufacturing of Functionally Graded Ceramic Materials by Stereolithography, JoVE-J. Vis. Exp. 143 (2019) e57943. [16] Hu K, Wei Y, Lu Z, et al. Design of a Shaping System for Stereolithography with High Solid Loading Ceramic

Suspensions, 3D Print Addit Manuf, 5 (2018) 311-318. [17] Li H, Liu Y, Liu Y, et al. Influence of sintering temperature on microstructure and mechanical properties of Al2O3 ceramic via 3D stereolithography, Acta Metallurgica Sinica English Letters. 2019. [18] Sun Y, Shimai S, Peng X, et al. Gelcasting and vacuum sintering of translucent alumina ceramics with high transparency, J. Alloy. Compd. 641 (2015) 75-79. [19] Chen Z, Liu C, Li J, et al. Mechanical properties and microstructures of 3D printed bulk cordierite parts, Ceram. Int. 2019. [20] Zhou W, Li D, Wang H. A novel aqueous ceramic suspension for ceramic stereolithography[J]. Rapid. Prototyping. J. 16 (2010) 29-35.

Taguchi method[J]. P. I. Mech. Eng. B.-J. Eng. 226 (2012) 1249-1258.

ro of

[21] Chen Z, Li D, Zhou W. Process parameters appraisal of fabricating ceramic parts based on stereolithography using the

[22] He Z, Ma J. Grain-growth rate constant of hot-pressed alumina ceramics, Mater. Lett. 44 (2000) 14-18.

[23] Kuczynski G C. Self-diffusion in sintering of metallic particles[M]//Sintering Key Papers. Springer, Dordrecht, 1990

-p

509-527.

[24] Hillert M. On the theory of normal and abnormal grain growth, Acta metal. 13 (1965) 227-238.

re

[25] Li J, Ye Y. Densification and grain growth of Al2O3 nanoceramics during pressureless sintering, J. Am. Ceram. Soc. 89 (2006) 139-143.

559-572.

lP

[26] Karagedov G R, Lyakhov N Z. Preparation and sintering of nanosized α-Al2O3 powder, Nanostruct. Mater. 11 (1999)

[27] Morrissey K J, Czanderna K K, Carter C B, et al. Growth of α‐ Al2O3 Within a Transition Alumina Matrix, J. Am.

na

Ceram. Soc. 67(1984) c88-c90.

[28] Gao Z. Tabular microstructure of alumina, Refract. Mater. 28 (1994) 293-299.

ur

[29] Kondrashova D, Valiullin R. Freezing and melting transitions under mesoscalic confinement: Application of the Kossel–Stranski crystal-growth model, J. Phys Chem C 119 (2015) 4312-4323.

Jo

[30] Chen M, He J, Zhang Y, et al. High purity and fully dense MgO ceramic by conventional and vacuum sintering, Chinese Journal of Rare Metals, 42 (2018): 402-407. [31] Montero C, Ochoa A, Castaño P, et al. Monitoring Ni0 and coke evolution during the deactivation of a Ni/La2O3– αAl2O3 catalyst in ethanol steam reforming in a fluidized bed, J. Catal. 331 (2015) 181-192. [32] Gupta M, Singh V, Katyal P. Synthesis and Structural Characterization of Al2O3 nanofluids, Mater. Today. 5 (2018) 27989-27997. [33] Shibata N, Chisholm M F, Nakamura A, et al. Nonstoichiometric dislocation cores in α-alumina, Science. 316 (2007) 82-85. [34] Kadleıková M, Breza J, Veselý M. Raman spectra of synthetic sapphire, Microelectron. J. 32 (2001) 955-958.

[35] Kingery W D, Berg M. Study of the initial stages of sintering solids by viscous flow, evaporation‐ condensation, and self‐ diffusion, J. Appl. Phys. 26 (1955) 1205-1212. [36] Owen D, Hickey J, Cusson A, et al. 3D printing of ceramic components using a customized 3D ceramic printer, Prog. Addit. Manuf. 3 (2018) 3-9. [37] Dehurtevent M, Robberecht L, Hornez J C, et al. Stereolithography: a new method for processing dental ceramics by additive computer-aided manufacturing, Dent. Mater. 33(2017) 477-485. [38] Braga R R, Ballester R Y, Ferracane J L. Factors involved in the development of polymerization shrinkage stress in resin-composites: a systematic review, Dent. Mater. 21(2005) 962-970. [39] Li Y Y, Li L T, Li B. Direct ink writing of 3–3 piezoelectric composite, J. Alloy. Compd. 620 (2015) 125-128. [40] Zhang K, Xie C, Wang G, et al. High solid loading, low viscosity photosensitive Al2O3 slurry for stereolithography

ro of

based additive manufacturing, Ceram. Int. 45(2019) 203-208.

[41] Zhao H, Ye C, Fan Z, et al. 3D printing of CaO-based ceramic core using nanozirconia suspension as a binder, J. Eur. Ceram. Soc. 37(2017) 5119-5125.

316L stainless steel, Appl. Surf. Sci. 431 (2018) 112-121.

-p

[42] Yu J, Wang Y, Zhou F, et al. Laser remelting of plasma-sprayed nanostructured Al2O3–20 wt.% ZrO2 coatings onto

re

[43] Jiang W, Li K, Xiao J, et al. Effect of silica fiber on the mechanical and chemical behavior of alumina-based ceramic core material, J. Asian. Ceram. Soc. 5 (2017) 410-417.

lP

[44] Manshor H, Aris S M, Azhar A Z A, et al. Effects of TiO2 addition on the phase, mechanical properties, and microstructure of zirconia-toughened alumina ceramic composite, Ceram. Int. 41 (2015) 3961-3967. [45] Xing X, Wang H, Xiao G, et al. Assessment of titanium metallization thin film deposited on alumina substrate:

na

Microstructure and nano-indentation, J. Mech. Behav. Biomed. 80 (2018) 235-245. [46] Mata-Osoro G, Moya J S, Pecharroman C. Transparent alumina by vacuum sintering[J]. J. Eur. Ceram. Soc. 32(2012)

ur

2925-2933.

[47] Srdić V V, Winterer M, Hahn H. Sintering behavior of nanocrystalline zirconia doped with alumina prepared by

Jo

chemical vapor synthesis[J]. J. Am. Ceram. So. 83 (2000) 1853-1860. [48] Hao H S, Xu L H, Huang Y, et al. Kinetics mechanism of microwave sintering in ceramic materials[J]. Sci. China. Ser. E: Technol. Sci. 52(2009) 2727-2731.