Advertisement

Rare Metals

, Volume 38, Issue 11, pp 1003–1014 | Cite as

Polymer-infiltrated layered silicates for dental restorative materials

  • Ben-Cang Cui
  • Jing Li
  • Yuan-Hua LinEmail author
  • Yang Shen
  • Ming Li
  • Xu-Liang Deng
  • Ce-Wen Nan
Article
  • 104 Downloads

Abstract

Layered porous ceramic used for polymer-infiltrated-ceramic-network materials (PICNs) may be a promising candidate for dental restoration. The effect of sintering temperature of ceramic green bodies on mechanical and optical properties of PICNs is unclear. The purpose was to fabricate PICNs and evaluate their mechanical and optical properties. Polymer-infiltrated layered silicates for dental restorative materials were prepared via infiltrating polymerizable monomers into partially sintered porous silicates and thermo-curing. Bending samples for flexural strength and fracture toughness were fabricated (sample numbers of n = 15). Vickers hardness and elastic modulus were measured via nano-indentation (n = 10). One-way ANOVA and Weibull statistics were used for statistical analysis. Optical property was characterized by spectral reflectance. Brittleness index was used to characterize the machinability of the materials. Microstructures and phase structures were investigated using scanning electron microscopy (SEM) and X-ray diffractometer (XRD), respectively. Flexural strength of polymer-infiltrated layered silicates varied from 91.29 to 155.19 MPa, fracture toughness ranged from 1.186 to 1.782 MPa·m1/2, Vickers hardness ranged from 1.165 to 9.596 GPa, and elastic modulus ranged from 25.35 to 100.50 GPa. The formed glass phases at 1200 and 1300 °C showed influences on corresponding optical property, which could be observed from spectral reflectance. A kind of PICNs was fabricated by infiltrating polymerizable monomers into layered porous ceramic networks. Sintering temperature could have dramatic effects on the mechanical and optical properties of porous ceramics and PICNs. These kinds of materials possess similar properties to that of natural tooth and could be used for dental restoration.

Keywords

CAD/CAM blocks Dental composites Mechanical properties Polymer-infiltrated-ceramic-networks 

1 Introduction

Nowadays, there is a great development in dentistry, especially in dental restorative materials. Patients’ preference for aesthetic restoration brings about the studies on all-ceramics and composite restoration and the launches of corresponding commercial products [1]. However, high hardness of all-ceramics causes the ware of the opposite and adjacent natural tooth [2, 3, 4]. High modulus leads to the mismatch of force transfer during the chewing process [5, 6] and gives rise to the early failure of restoration. Because of the brittleness and susceptibility to micro-cracks of all-ceramics [6, 7], their performances are highly technically dependent [6, 8]. All of these restrict their utilization in dental restoration. In recent years, resin composites with high aesthetic performance and excellent mechanical properties find significant applications in dental restoration.

According to the processing technique, resin composites are classified into particle-filled composites and polymer-infiltrated-ceramic-network materials (PICNs) [9]. Particle-filled composites are fabricated through filling inorganic particles into polymerizable monomers. Particle-filled composites have been used for direct restoration for more than 50 years [10, 11]. Filling material and cement are the two most common clinical applications of the former. Nowadays, high-loading particle-filled composites are also being used for CAD/CAM blocks, with Lava Ultimate (3 M ESPE) and Cerasmart (GC Dental) as two representative commercial materials. For particle-filled composites, inorganic fillers are dispersed uniformly into the continuous phase polymer matrix, which is then cured by light or heat polymerization [12, 13]. Different from particle-filled composites, PICNs are fabricated through infiltrating polymerizable organic monomers into pre-sintered porous ceramic networks and curing.

The main objectives of dental restoration are to mimic the natural teeth at structure and physical properties. The three-dimensional interconnected structure of PICNs combines the ceramic phase and the polymer matrix together, through which, similar properties to natural tooth could be realized [14]. Each continuous phase in PICNs could contribute to the ultimate properties of composites. And therefore, both factors influencing the ceramic phase and the polymer phase could be considered to tailor the properties of PICNs. Precisely, ceramic particle size and shape, porous ceramic forming technology, polymerizable monomer systems and curing protocols could be manipulated for the improvement in properties. The green bodies can be prepared either via mold press [15] or via slip casting [16]. The polymerizable monomers are cured via photo-polymerization and heat polymerization. During polymerization, high temperature and high pressure could also be introduced. Okada K et al. [15] prepared a kind of PICNs at high compression molding pressure. Coldea et al. [17] fabricated polymer-infiltrated-ceramic-network materials with different ceramic densities (by manipulating ceramic particle size and sintering temperature) and evaluated the correlation between several mechanical properties and ceramic density. Nguyen et al. [18, 19, 20] obtained PICNs with superior mechanical properties through introducing high pressure and high temperature to polymerization.

In filled dental resin composite, particle shape has a significant influence on the ultimate properties [21, 22, 23, 24]. Mechanical properties of polymer-infiltrated nanoparticle porous ceramics [25] and nano-clusters porous ceramic networks [26, 27] have been examined. However, polymer-infiltrated layered porous ceramic network composite has rarely been investigated. The objectives of this paper were to fabricate polymer-infiltrated layered porous ceramic networks and characterize its properties. Sintering temperature could influence the characteristics of porous ceramics and mechanical and optical properties of the resultant composites. The flexural strength, fracture toughness, elastic modulus and Vickers hardness of composites and porous ceramics were compared. The optical properties were also evaluated.

2 Experimental

2.1 Characterization of inorganic powder

Particle size distributions of layered potassium aluminosilicate powder (Advantaged Chemical Co., Ltd., Taiping, Taiwan) were measured by a particle size analyzer (Mastersizer Hydro 2000MU). X-ray fluorescence (XRF, ARL Perform’X) was used for the chemical analysis.

2.2 Fabrication of samples

Layered potassium aluminosilicate was used to prepare porous layered ceramic networks. After fabrication of porous layered potassium aluminosilicate, PICNs were prepared by infiltrating monomer mixtures containing bisphenol-A-dimethacrylate, BisGMA, and triethylene glycol dimethacrylate, TEGDMA (Aladdin Reagents Company, Shanghai, China) into the partially sintered potassium aluminosilicate blocks and curing via heat polymerization. TEGDMA was used as a diluent with mass ratio of 50:50 to decrease the viscosity of BisGMA. Dibenzoyl peroxide, BPO (J & K Scientific LTD., Beijing, China), was used as thermo-initiator. The whole process includes three steps.

Firstly, ceramic networks were fabricated by partially sintering potassium aluminosilicate. PVA aqueous solution (3 wt%) adhesive was added into potassium aluminosilicate powder with a dosage of 0.05 g per gram of ceramic particles. Potassium aluminosilicate block green bodies were fabricated via mold press with a pressure holding time of 3 min. An isostatic cool pressing was followed with a pressure of 220 MPa and a holding time of 1.5 min. In order to obtain different porous structures, different temperatures (700–1300 °C) were used during the sintering process. Ceramic with a porous structure and nearly densely sintered ceramic were obtained through heating rate of 5 °C·min−1 and heat preservation of 120 min. In the second step, monomer mixtures were infiltrated into the partially sintered porous blocks and thermo-cured. The porous blocks were partially immersed into the polymerizable monomers containing 1 wt% dibenzoyl peroxide. Then, the porous blocks together with monomer mixture were put in a vacuum drying oven. The vacuum degree is 0.1 MPa. The complete infiltration of monomers was realized by vacuum capillary action at room temperature. The vacuum action could also eliminate bubbles in the monomer. After a dwell time of more than 24 h, monomers appeared on top of all pieces, indicating that the monomers were infiltrated through porous blocks. Then, the infiltrated pieces were heated up to 70 °C with a dwell time of 8 h. The cured pieces were continually heated up to 110 °C with a dwell time of 8 h to increase the degree of conversion. In the third step, PICNs and ceramic blocks were cut into testing bars according to corresponding testing standards. Pieces including PICNs and ceramics were sanded to 2.2 mm using 2000# waterproof abrasive paper. Bars with dimensions of 2.2 mm × 2.2 mm × 20.0 mm (sample numbers of n = 15), 4.2 mm × 2.2 mm × 20.0 mm (n = 15) and 4.2 mm × 2.2 mm × 10.0 mm (n = 1) of each material were cut from the above materials with a cutting machine. The rotating speed of the diamond saw was 3000 r·min−1, and feeding speed was 6 mm·min−1.

2.3 Flexural property tests

Testing bars with size of 2.2 mm × 2.2 mm × 20.0 mm (n = 15) were sanded to 2 mm × 2 mm × 20 mm (n = 15) for flexural samples. The bending bars were chamfered and polished thoroughly. The three-point bending testing was carried out using a universal testing machine (Shimadzu, EZ-100) according to ISO-4049. The flexural strength (σf) was calculated by Eq. (1):
$$\sigma_{\text{f}} = \frac{3Fl}{{2bd^{2} }}$$
(1)
where F is the maximum load value at fracture; l is the roller span distance (here 10 mm); b is the width and d is the height at the fracture point. The loading speed was 0.5 mm·min−1. The force was exerted onto the bars continuously until fracture.

2.4 Fracture toughness tests

Bending bars (n = 15) with dimensions of 4.2 mm × 2.2 mm × 20.0 mm (n = 15) were sanded to the final dimension of 4 mm × 2 mm × 20 mm (n = 15) for fracture toughness testing. The specimens were chamfered and polished before pre-notching. Then, the bars were pre-notched with a depth of 2 mm at the middle point with a high-speed cutting blade. An optical microscope (OM, Olympus BX50) was used to measure the exact depth value of each notch. The depth of each notch was measured twice from both sides. Then, the two values were averaged.

The fracture toughness (KIC) was measured with the standard single-edge-notched beam (SENB) method in a three-point bending format till fracture. The testing machine was the same as flexural property testing. The supporting span was 15 mm, and the loading speeding was 0.05 mm·min−1. KIC was calculated using the following equation [1],
$$K_{\text{IC}} = \frac{FL}{{BW^{3/2} }}f\left( {\frac{a}{W}} \right)$$
(2)
where B and W are the breadth and the width of the bar, respectively, F is the maximum load value at fracture, L is roller span distance (here 15 mm), and a is the notch depth. The function f(a/W) is the geometrical factor calculated using Eq. (3) [1]:
$$f\left( {\frac{a}{W}} \right) = 2.9\left( {\frac{a}{W}} \right)^{1/2} - 4.6\left( {\frac{a}{W}} \right)^{3/2} + 21.8\left( {\frac{a}{W}} \right)^{5/2} - 37.6\left( {\frac{a}{W}} \right)^{7/2} + 38.7\left( {\frac{a}{W}} \right)^{9/2}$$
(3)

2.5 Nano-indentation tests

Elastic modulus (E) and Vickers hardness (HV) were measured on thoroughly polished surfaces of each specimen. One specimen (4.2 mm × 2.2 mm × 10.0 mm) of each material was sanded and polished. The indentations were placed with a nano-indentation tester (MTS, XP). A triangular pyramid Berkovich tip was used to exert the load continuously on the polished surface of each specimen. Ten indentation points were placed on each specimen randomly. The nano-indentation depth was 1000 nm. The Oliver–Pharr method [28] was used to calculate elastic modulus (E) and Vickers hardness (HV).

2.6 Brittleness index calculation

Brittleness is an essential factor of CAD/CAM blocks which is used to quantify the machinability of dental restorative materials. Brittleness values (BI) and the standard deviations (SD) were calculated using Eqs. (4) [1] and (5):
$$BI = \frac{{\text{HV}}}{{K_{\text{IC}} }}$$
(4)
$$ \sigma_{BI} = \frac{\text{HV}}{{K_{{\text{IC}}} }}\sqrt {\left( {\frac{{\partial \ln {BI}}}{{\partial {\text{HV}}}}} \right)^{2} \left( {\sigma_{\text{HV}} } \right)^{2} + \left( {\frac{\partial \ln BI}{{\partial K_{{\text{IC}}} }}} \right)^{2} \left( {\sigma_{{K_{{\text{IC}}} }} } \right)^{2} } $$
(5)
where σBI, σHV and \(\sigma_{{K_{\text{IC}} }}\) are corresponding standard deviations (SD).

2.7 Spectral reflectance

Spectral reflectance value was used to characterize the optical properties of the materials. Samples with a thickness of 0.8 mm were thoroughly polished. Spectral reflectance of ceramics and composites was measured against white by a spectrophotometer (PerkinElmer Lambda 950, America).

2.8 Porosity, pore size distribution and density characterization

Mercury intrusion porosimetry (MIP) is a commonly used method to characterize the porosity and the pore size distribution of porous materials. Total porosity is obtained from the intrusion volume of mercury at a pressure of 414 MPa. The cumulative intrusion volume of mercury is measured at each pressure increment. The pore size distribution is calculated from corresponding intrusion volume increment versus each pore size interval. Bulk density is calculated via dividing bulk mass by bulk volume. The volume used to calculate the apparent (skeletal) density is the one with the removal of total opening volume.

2.9 SEM observation

Scanning electron microscope (SEM, MERLIN Compact, ZEISS) was used to investigate the microstructure of both PICNs and ceramics. The morphology of layered potassium aluminosilicate was also observed via SEM. The polished surfaces of both composites and ceramics were platinum coated and examined using a secondary electron detector.

2.10 Phase analysis

The phase structure was analyzed via X-ray diffractometer (XRD, Rigaku D/MAX-2550V). Effects of the sintering temperature of the green bodies on the phase deformation were evaluated based on XRD.

2.11 Statistical analysis

One-way ANOVA was used to analyze flexural strength, elastic modulus, Vickers hardness, fracture toughness and brittleness index. The significance level was set at p < 0.05. The flexural strength and fracture toughness of composites were subjected to Weibull statistics according to Eq. (6):
$$\ln \ln \left[ {1/\left( {1 - P_{n} } \right)} \right] = m\ln \sigma - m\ln \sigma_{0}$$
(6)
where σ is the tested flexural strength or fracture toughness value of each sample, Pn is the failure probability estimator, σ0 is the corresponding characteristic value and m is the Weibull modulus. Pn (probability of failure) was calculated according to Eq. (7):
$$P_{n} = {{\left( {n - 0.5} \right)} \mathord{\left/ {\vphantom {{\left( {n - 0.5} \right)} N}} \right. \kern-0pt} N}$$
(7)
where n refers to the nth bar tested, and N is the whole number of the samples tested.

3 Results and discussion

3.1 Characterization of layered potassium aluminosilicate

The most commonly used for characterizing particle size is particle size distribution curve. Figure 1a shows the particle size distribution of layered potassium aluminosilicate. The curve was plotted after taking a logarithm of 10 for the particle size. Particle sizes range from tens of nanometers to several microns. The average particle sizes based on surface area and volume are 0.49 and 2.56 μm, respectively. Figure 1b shows the lamellar microstructure of the inorganic particles. The flakes appear to be peeled off from the large particles.
Fig. 1

a Particle size distribution and b SEM analysis of layered potassium aluminosilicate

Table 1 presents the compositions of the inorganic part measured via XRF. The measured composition indicates that the powder is aluminum–potassium silicate rich in silica.
Table 1

Composition of ceramic part (wt%)

SiO2

Al2O3

K2O

Fe2O3

SO3

ZrO2

MgO

Na2O

TiO2

NiO

81.02

14.48

3.22

0.434

0.202

0.17

0.135

0.0839

0.0762

0.0589

CaO

Rb2O

Cl

P2O5

CeO2

MnO

Cr2O3

Y2O3

Ga2O3

 

0.0255

0.0191

0.0159

0.0159

0.0135

0.0125

0.0082

0.0061

0.0043

 

3.2 Mechanical properties

The mechanical property results including flexural strength, fracture toughness, Vickers hardness and elastic modulus of polymer-infiltrated layered silicates are shown in Table 2. Considering that nano-indentation data may be invalid for porous materials, only flexural strength and fracture toughness for porous layered potassium aluminosilicate are presented in Table 3. The flexural strength was unmeasurable for 700 °C sintered porous ceramics. As for fracture toughness of porous ceramics, it was unmeasurable for both 700 and 800 °C sintered groups. All the mechanical property values were tested by single-factor ANOVA.
Table 2

Measured mechanical properties of polymer-infiltrated layered silicates, where p value between groups is the difference between corresponding one and previous one

Mechanical properties

700 °C

800 °C

900 °C

1000 °C

1100 °C

1200 °C

1300 °C

Flexural strength (SD)/MPa

91.29 (10.80)

108.55 (7.77)

109.48 (7.14)

108.76 (13.83)

132.96 (11.65)

155.19 (16.29)

110.20 (10.90)

p value between groups

0.02

> 0.99

> 0.99

< 0.001

< 0.001

< 0.001

Fracture toughness (SD)/(MPa·m1/2)

1.250 (0.089)

1.186 (0.0690)

1.244 (0.104)

1.505 (0.155)

1.648 (0.115)

1.782 (0.174)

1.366 (0.223)

p value between groups

0.879

0.921

< 0.001

0.094

0.142

< 0.001

Vickers hardness (SD)/GPa

1.165 (0.119)

1.211 (0.112)

1.377 (0.165)

1.659 (0.138)

2.685 (0.244)

7.644 (0.762)

9.596 (0.512)

p value between groups

0.99

0.953

0.625

< 0.001

< 0.001

< 0.001

Elastic modulus (SD)/GPa

25.35 (1.15)

28.30 (1.31)

30.25 (1.63)

39.94 (1.96)

48.56 (3.02)

71.48 (3.15)

100.50 (4.31)

p value between groups

0.162

0.631

< 0.001

< 0.001

< 0.001

< 0.001

Brittleness index (SD)/(μm−1/2)

0.932 (0.100)

1.021 (0.097)

1.107 (0.141)

1.102 (0.103)

1.629 (0.156)

4.289 (0.468)

7.025 (0.563)

p value between groups

0.059

0.129

0.929

< 0.001

< 0.001

< 0.001

Table 3

Measured mechanical properties of porous layered silicates, where p value between groups is the difference between corresponding one and previous one

Mechanical properties

700 °C

800 °C

900 °C

1000 °C

1100 °C

1200 °C

1300 °C

Flexural strength (SD)/MPa

Unmeasurable

8.05 (1.85)

9.17 (1.02)

14.01 (1.12)

72.18 (5.21)

152.30 (12.67)

111.42 (10.50)

p value between groups

> 0.99

0.432

< 0.001

< 0.001

< 0.001

Fracture toughness (SD)/(MPa·m1/2)

Unmeasurable

Unmeasurable

0.27 (0.041)

0.25 (0.012)

0.80 (0.056)

1.80 (0.093)

1.38 (0.144)

p value between groups

0.920

< 0.001

< 0.001

< 0.001

With the increase in pre-sintering temperature, the flexural strengths of both porous silicates and composites increased and then decreased, with the highest values at 1200 °C. For composites, the flexural strength value increased from 91.29 to 155.19 MPa and decreased to 110.20 MPa. The flexural strengths of composites for 800, 900 and 1000 °C groups were not significantly different at 95% confidence, with p = 1. All other flexural strengths of composites were significantly different via one-way ANOVA. For porous silicates, the flexural strengths of 700 °C pre-sintered groups were unmeasurable. With the increase in pre-sintered temperature from 800 to 1300 °C, the flexural strength of porous silicates increased from 8.05 to 152.30 MPa and decreased to 111.42 MPa. The flexural strengths of porous silicates for 800, 900 and 1000 °C groups were not significantly different at 95% confidence, with p = 0.998 and p = 0.432, respectively. All other flexural strengths of porous silicates were significantly different via one-way ANOVA.

With the increase in pre-sintered temperature, fracture toughness of composites varied from 1.186 to 1.782 MPa·m1/2. The fracture toughness values of composites for all groups were not significantly different at 95% confidence, except for 900/1000 and 1200/1300 °C. For porous silicates, the fracture toughness values of 700 and 800 °C pre-sintered groups were unmeasurable. The fracture toughness values of porous silicates for 900 and 1000 °C groups were not significantly different at 95% confidence, with p = 0.920. All other fracture toughness values of porous silicates were significantly different via one-way ANOVA.

With the increase in pre-sintered temperature, Vickers hardness of composites varied from 1.165 to 9.596 GPa and showed a generally increasing trend (Table 2). The Vickers hardness values of composites were not significantly different at 95% confidence when the pre-sintered temperature increased from 700 to 1000 °C, while significantly different from 1000 to 1300 °C via one-way ANOVA. Elastic modulus values of PICNs are also listed in Table 2. The elastic modulus values of composites were significantly different except for 700, 800 and 900 °C groups via one-way ANOVA.

Brittleness index is presented in Table 2, as a parameter to evaluate the machinability of materials. Brittleness index values of composites were not significantly different at 95% confidence when the pre-sintered temperature increased from 700 to 1000 °C, while significantly different from 1000 to 1300 °C via one-way ANOVA. Brittleness index values of composites showed an increasing trend with pre-sintered temperature.

The goal of dental material research is to achieve its similarity to natural teeth. Flexural strength, elastic modulus hardness and fracture toughness of natural enamel were 60–90 MPa [29], 47–120 GPa [30], 2.7–6.4 GPa [30, 31], (1.3 ± 0.3) GPa [32], respectively. As shown in Table 2, the mechanical properties of the fabricated PICNs could mimic the natural enamel at a high level.

3.3 Weibull plots of flexural strength and fracture toughness

The Weibull modulus of both flexural strength and fracture toughness of composites is presented in Fig. 2a, b. High Weibull modulus indicates centralized and uniform data for flexural strength and fracture toughness.
Fig. 2

Weibull plots of flexural strength (σ) and fracture toughness (KIC): a Weibull plots of flexural strength and b Weibull plots of fracture toughness

Weibull modulus was commonly used as the statistical parameter to evaluate the dispersity and credibility of mechanical values [33]. Because of its success in describing the fracture data, Weibull distribution was often considered to be the first choice in the analysis procedure [34]. High Weibull modulus indicates the centralization of flexural strength and fracture toughness values. Weibull modulus values vary from 9.4 to 18.38 and from 7.04 to 20.37 for flexural strength and fracture toughness, respectively, much higher than those of zirconia dental restorative ceramics [35, 36].

3.4 Porosity, pore size distribution and density

As shown in Fig. 3a, with the increase in pre-sintered temperature, there is a decreasing trend of porosity for porous silicates, indicating the densification of ceramic networks. The open porosity varies from 40.84% to 0.62% for ceramic networks. After polymer infiltration, a significant reduction in porosity is observed. The obtained porosity via mercury intrusion porosimetry varies from 1.24% to 3.53%, indicating that most open pores are filled with polymers. Figure 3b presents pore size distribution of porous inorganic bulk at different sintering temperatures, where V represents cumulative intrusion volume and R is the pore size. With the increase in sintering temperature, the open pore size is observed to increase. Figure 3c shows the skeletal and bulk density of porous ceramic networks and PICNs. Samples 1–7 refer to porous ceramic networks with sintering temperature from 700 to 1300 °C, while Samples 8–12 refer to PICNs with pre-sintering temperature from 700 to 1100 °C. A significant increase in bulk density was observed after infiltration of polymer, indicating that polymer filled the open pores of ceramic networks. The skeletal density of porous ceramic was mainly dependent on the closed pores. The decrease in skeletal density with the increase in sintering temperature reflected the formation of closed pores. As the sintering temperature increased from 700 to 1200 °C, the densification of ceramic networks began, and thus, the bulk density increased. The skeletal density of PICNs varied from 2.08 to 2.13 g·cm−3. The density value is very close to that of natural dentin (1.79 to 2.12 g·cm−3) [37]. This similarity can be significant, as a density higher than that of natural teeth may cause additional falling force.
Fig. 3

a Porosity, b pore size distribution and c density of porous silicates and PICNs

3.5 Phase composition

Polymer and ceramic may exhibit different degrees of light and shade in SEM. This phenomenon can be attributed to the principle of atomic number contrast. Generally, the yields of both secondary electrons and backscattered electrons increase with the increase in atomic number. Therefore, in the analysis, the region with a higher atomic number can emit more secondary electrons and backscattered electrons than the region with a lower atomic number. Therefore, the image of the region with a higher atomic number is brighter. The polymer phase in PICNs is composed of light element carbon, hydrogen and oxygen. However, most of the elements in the ceramic phase are silicon, oxygen, aluminum, potassium, etc., with atomic number higher than that of polymer phase. To further identify the polymer and ceramic phase in PICNs, elemental distribution mapping (Fig. 4b, c) for the same region (Fig. 4a) of the polished surface was analyzed. Figure 4b, c shows the carbon and silicon distribution mapping over the total scanned area (Fig. 4a). It can be observed that the region enriched with carbon corresponded to the dark area, and the other to the bright one.
Fig. 4

Phase identification of PICNs: a SEM cross-sectional image of PICNs (700 °C group); elemental mapping of same region of b C and c Si

3.6 SEM results

Microstructural observation of porous silicates and corresponding composites via SEM is presented in Fig. 5. With the increase in pre-sintered temperature of the green bodies, the layered silicate disappeared and massive structures formed gradually. As the pre-sintering temperature increased from 700 to 1200 °C, a porous structure formed for ceramics. For 1300 °C pre-sintered groups, a nearly dense structure formed. Polymerizable monomers could be infiltrated into the porous ceramic networks thoroughly when the layered silicates were pre-sintered below 1100 °C. No polymer was found for 1200 and 1300 °C groups. When layered silicate blocks were pre-sintered at 1200 °C, a porous structure could also be formed. However, most of the pores were closed, and no polymerizable monomers could be infiltrated into the networks.
Fig. 5

SEM images of polished surfaces of porous silicates and composites: a1a7 porous silicates and b1b7 corresponding composites with pre-sintering temperature from 700 to 1300 °C

As illustrated in Fig. 5b1-b5, two continuous phases could be observed, i.e., ceramic network (the light areas) and polymer network (the gray areas). For Fig. 5b6, b7, no polymer phase could be found.

3.7 Phase structure of pre-sintered silicates

XRD patterns of pre-sintered silicates are presented in Fig. 6. Most of the diffraction peaks indicate that the main crystalline phase of the mineral is hexagonal quartz (as referred by the asterisk). Zeolite, sillimanite and sanidine may also be contained. All the pre-sintered silicates exhibited the same dominant crystalline phase. When the pre-sintered temperature reached 1200 °C, some low-temperature glass phases were observed. In particular, for the 1300 °C group, a large hump band was observed, indicating that glass phase spectrum tended to enhance.
Fig. 6

XRD patterns of pre-sintered silicates

3.8 Spectral reflectance and contrast ratio of composites

Spectral reflectance of porous silicates and composites against a white/black backgrounds and corresponding contrast ratio are presented in Fig. 7a, b, where B and W refer to the black and the white backgrounds, respectively. After infiltration of monomers and curing, the composites tended to have less spectral reflectance than corresponding porous silicates. For both black and white backgrounds, spectral reflectance of composites tended to increase with the increase in pre-sintered temperature from 700 to 1200 °C. When the pre-sintered temperature increased to 1300 °C, the spectral reflectance of composites decreased compared to that of 1200 °C groups, and was even less than that of 1100 °C groups in the region of medium and long wavelengths. The contrast ratio is commonly used to characterize the translucency of dental materials [38]. The value of contrast ratio was calculated from the ratio of spectral reflectance against black and white backgrounds. Contrast ratio varies from 0 to 1.0, where 0 refers to a transparent material, and 1.0 refers to a wholly opaque one [39, 40]. The results indicated that light translucency of PICNs was controlled by multiple factors. The highest translucency was observed for PICNs at a pre-sintered temperature of 900 °C. There was a general decreasing trend of translucency with the increase in pre-sintered temperature. This trend may be caused by the formation of closed pores, as indicated in Fig. 6.
Fig. 7

a Spectral reflectance of composites against white/black (W/B) backgrounds and b corresponding contrast ratio values

The aesthetic property is another essential performance of dental materials. Owing to superior mechanical properties and biocompatibility, titanium alloys are listed as important biomaterials [41, 42, 43]. However, they were currently often used as implants in dental restoration for aesthetic considerations. As the requirements of this property increased, more and more attentions were paid to aesthetics. Optical properties, such as spectral reflectance, could be used to evaluate the aesthetics of dental materials. The effect of pre-sintering temperature and polymer-ceramic interface played a significant role in the spectral reflectance of ceramics and composites.

After the densification of ceramics, a transformation to glass phase was observed in XRD patterns. The decrease in mechanical properties of 1300 °C groups and the spectral reflectance may be attributed to this transformation.

3.9 Comparison between mechanical properties

Dental materials are subjected to a complex mechanical environment. And therefore, the study of dental materials must deal with the resistance of materials to deformation and fracture. Flexural strength, fracture toughness, hardness and elastic modulus are the main mechanical properties. As for mechanical properties, it is the goal for dental material researchers to fabricate a kind of material with the corresponding performances comparable to that of the natural tooth. The natural tooth possesses two phases interconnecting with the other. Polymer-infiltrated ceramic networks possess the same structure and exhibit performances similar to that of the natural tooth.

Figure 8 presents the comparison of mechanical properties of composites to that of porous ceramics. When the pre-sintering temperature was below 1100 °C, the values of flexural strength and fracture toughness were significantly different. After infiltration of polymerizable monomers into the porous ceramic networks, enhancements were observed for both properties. Brittleness index was used to evaluate the machinability of dental materials. As the pre-sintering temperature increased, brittleness index value increased, indicating that densification of the ceramic parts would lead to the unmachinability of the materials. When the pre-sintering temperature increased to 1200 and 1300 °C, the “composite” group behaved the same as ceramic. Though they are termed “composites,” they are actually ceramics, as no infiltrated polymer was found.
Fig. 8

Comparison of mechanical properties of composites to that of porous ceramics: a flexural strength and b fracture toughness

For dental porcelains, there may be some correlations between fracture toughness and flexural strength [44]. However, the relationship may be complicated for PICNs. The characteristic strength (σ0) was obtained from the Weibull plots of flexural strength. To further examine the relationship between fracture toughness and flexural strength, the data were subjected to a polynomial fitting. Figure 9 illustrates the fitting polynomial relationship between KIC and σ0 for PICNs.
Fig. 9

Polynomial fitting polynomial relationship between K1C and characteristic strength (σ0) for PICNs

For dental porcelains, flexural strength was found to increase and then decrease with the increase in fracture, which is attributed to the microstructure of the materials [44]. As shown in Fig. 9, the characteristic strength increased with the increase in fracture toughness. This characteristic may be attributed to the particular structure of PICNs. These materials may contain many kinds of defects, such as micro-cracks in ceramic phase, the interface between polymer and ceramic and voids in polymer phase. PICNs are interpenetrating phase composites, i.e., the polymer phase and ceramic interconnecting with each other. The ductile polymer may contribute to distributing stress uniformly and avoiding stress concentration effectively.

4 Conclusion

The main objectives of this study were to fabricate composites with dual networks through infiltrating polymerizable monomers into porous layered silicates and curing and to characterize the properties (mechanical and optical properties). Based on the results of this study, the following conclusions may be obtained. The chemical composition and main crystal structure of the ceramic part were analyzed. The pre-sintering temperature of the green bodies affects the mechanical properties of porous ceramics and corresponding composites. Infiltration of polymerizable monomers into porous ceramics could contribute to the increase in mechanical properties. Spectral reflectance and transparency of composites and ceramic are influenced by the pre-sintering temperature and glass phase transformation. The mechanical properties of the fabricated PICNs could mimic the natural enamel at a high level. The density value of the fabricated PICNs is very close to that of the natural dentin. The characteristic strength increases with the increase in fracture toughness. The fabricated PICNs with mechanical properties similar to that of natural human enamel could be a promising candidate for dental CAD/CAM.

Notes

Acknowledgements

This work was financially supported by Beijing Municipal Science and Technology Commission (No. Z171100002017009) and the National Natural Science Foundation of China (Nos. 51532003, 51221291, 51328203 and 81671026).

References

  1. [1]
    He LH, Swain M. A novel polymer infiltrated ceramic dental material. Dent Mater. 2011;27(6):527.Google Scholar
  2. [2]
    Sripetchdanond J, Leevailoj C. Wear of human enamel opposing monolithic zirconia, glass ceramic, and composite resin: an in vitro study. J Prosthet Dent. 2014;112(5):1141.Google Scholar
  3. [3]
    Kim M, Oh S, Kim J, Ju S, Seo D, Jun S, Ahn J, Ryu J. Wear evaluation of the human enamel opposing different Y-TZP dental ceramics and other porcelains. J Dent. 2012;40(11):979.Google Scholar
  4. [4]
    Lawson NC, Janyavula S, Syklawer S, McLaren EA, Burgess JO. Wear of enamel opposing zirconia and lithium disilicate after adjustment, polishing and glazing. J Dent. 2014;42(12):1586.Google Scholar
  5. [5]
    Aboushelib MN, Kleverlaan CJ, Feilzer AJ. Microtensile bond strength of different components of core veneered all-ceramic restorations: Part II: zirconia veneering ceramics. Dent Mater. 2006;22(9):857.Google Scholar
  6. [6]
    Kelly JR. Clinically relevant approach to failure testing of all- ceramic restorations. J Prosthet Dent. 1999;81(6):652.Google Scholar
  7. [7]
    Gemalmaz D, Ergin Ş. Clinical evaluation of all-ceramic crowns. J Prosthet Dent. 2002;87(2):189.Google Scholar
  8. [8]
    Manicone PF, Rossi Iommetti P, Raffaelli L. An overview of zirconia ceramics: basic properties and clinical applications. J Dent. 2007;35(11):819.Google Scholar
  9. [9]
    Swain MV, Coldea A, Bilkhair A, Guess PC. Interpenetrating network ceramic-resin composite dental restorative materials. Dent Mater. 2016;32(1):34.Google Scholar
  10. [10]
    Chen M. Update on dental nanocomposites. J Dent Res. 2010;89(6):549.Google Scholar
  11. [11]
    Ferracane JL. Resin composite—state of the art. Dent Mater. 2011;27(1):29.Google Scholar
  12. [12]
    Moszner N, Salz U. New developments of polymeric dental composites. Prog Polym Sci. 2001;26(4):535.Google Scholar
  13. [13]
    Moszner N, Hirt T. New polymer-chemical developments in clinical dental polymer materials: enamel–dentin adhesives and restorative composites. J Polym Sci Part A: Polym Chem. 2012;50(21):4369.Google Scholar
  14. [14]
    Della Bona A, Corazza PH, Zhang Y. Characterization of a polymer-infiltrated ceramic-network material. Dent Mater. 2014;30(5):564.Google Scholar
  15. [15]
    Okada K, Kameya T, Ishino H, Hayakawa T. A novel technique for preparing dental CAD/CAM composite resin blocks using the filler press and monomer infiltration method. Dent Mater J. 2014;33(2):203.Google Scholar
  16. [16]
    Nguyen JF, Ruse D, Phan AC, Sadoun MJ. High-temperature-pressure polymerized resin-infiltrated ceramic networks. J Dent Res. 2014;93(1):62.Google Scholar
  17. [17]
    Coldea A, Swain MV, Thiel N. Mechanical properties of polymer-infiltrated-ceramic-network materials. Dent Mater. 2013;29(4):419.Google Scholar
  18. [18]
    Nguyen J, Migonney V, Ruse ND, Sadoun M. Resin composite blocks via high-pressure high-temperature polymerization. Dent Mater. 2012;28(5):529.Google Scholar
  19. [19]
    Nguyen J, Migonney V, Ruse ND, Sadoun M. Properties of experimental urethane dimethacrylate-based dental resin composite blocks obtained via thermo-polymerization under high pressure. Dent Mater. 2013;29(5):535.Google Scholar
  20. [20]
    Eldafrawy M, Nguyen JF, Mainjot AK, Sadoun MJ. A functionally graded picn material for biomimetic CAD/CAM blocks. J Dent Res. 2018;97(12):1324.Google Scholar
  21. [21]
    Satterthwaite JD, Maisuria A, Vogel K, Watts DC. Effect of resin-composite filler particle size and shape on shrinkage-stress. Dent Mater. 2012;28(6):609.Google Scholar
  22. [22]
    Suzuki S, Nagai E, Taira Y, Minesaki Y. In vitro wear of indirect composite restoratives. J Prosthet Dent. 2002;88(4):431.Google Scholar
  23. [23]
    Wille S, Hölken I, Haidarschin G, Adelung R, Kern M. Biaxial flexural strength of new Bis-GMA/TEGDMA based composites with different fillers for dental applications. Dent Mater. 2016;32(9):1073.Google Scholar
  24. [24]
    Randolph LD, Palin WM, Leloup G, Leprince JG. Filler characteristics of modern dental resin composites and their influence on physico-mechanical properties. Dent Mater. 2016;32(12):1586.Google Scholar
  25. [25]
    Li J, Zhang X, Cui B, Lin Y, Deng X, Li M, Nan C. Mechanical performance of polymer-infiltrated zirconia ceramics. J Dent. 2017;58:60.Google Scholar
  26. [26]
    Wang H, Cui B, Li J, Li S, Lin Y, Liu D, Li M. Mechanical properties and biocompatibility of polymer infiltrated sodium aluminum silicate restorative composites. J Adv Ceram. 2017;6(1):73.Google Scholar
  27. [27]
    Cui B, Li J, Wang H, Lin Y, Shen Y, Li M, Deng X, Nan C. Mechanical properties of polymer-infiltrated-ceramic (sodium aluminum silicate) composites for dental restoration. J Dent. 2017;62:91.Google Scholar
  28. [28]
    Lawn BR, Marshall DB. Hardness, toughness, and brittleness: an indentation analysis. J Am Ceram Soc. 1979;62(7–8):347.Google Scholar
  29. [29]
    Braga RR, Denry IL, Ferracane JL, Khajotia SS, Mahler DB, Marshall GW, Marshall SJ, Mitchell JC, Mitra SB, Muenchinger KL, Pfeifer CS, Powers JM, Sakaguchi RL. Craig’s restorative dental materials (thirteenth edition). Edited by Sakaguchi RL, Powers JM. Saint Louis: Mosby.2012.16.Google Scholar
  30. [30]
    Cuy JL, Mann AB, Livi KJ, Teaford MF, Weihs TP. Nanoindentation mapping of the mechanical properties of human molar tooth enamel. Arch Oral Biol. 2002;47(4):281.Google Scholar
  31. [31]
    Jeng Y, Lin T, Hsu H, Chang H, Shieh D. Human enamel rod presents anisotropic nanotribological properties. J Mech Behav Biomed. 2011;4(4):515.Google Scholar
  32. [32]
    White SN, Luo W, Paine ML, Fong H, Sarikaya M, Snead ML. Biological organization of hydroxyapatite crystallites into a fibrous continuum toughens and controls anisotropy in human enamel. J Dent Res. 2001;80(1):321.Google Scholar
  33. [33]
    Della Bona A, Anusavice KJ, DeHoff PH. Weibull analysis and flexural strength of hot-pressed core and veneered ceramic structures. Dent Mater. 2003;19(7):662.Google Scholar
  34. [34]
    Danzer R, Fischer FD, Lu C. Fracture statistics of brittle materials: Weibull or normal distribution. Phys Rev E. 2002;65(6):67102.Google Scholar
  35. [35]
    Ruales-Carrera E, Cesar PF, Henriques B, Fredel MC, Özcan M, Volpato CAM. Microtensile bond strength of zirconia after surface treatments and aging. Dent Mater. 2018;34:100.Google Scholar
  36. [36]
    Sen N, Us YO. Mechanical and optical properties of monolithic CAD–CAM restorative materials. J Prosthet Dent. 2018;119(4):593.Google Scholar
  37. [37]
    Gradl R, Zanette I, Ruiz-Yaniz M, Dierolf M, Rack A, Zaslansky P, Pfeiffer F. Mass density measurement of mineralized tissue with grating-based X-ray phase tomography. PLoS ONE. 2016;11(12):e167797.Google Scholar
  38. [38]
    Miyagawa Y, Powers JM, O’Brien WJ. Optical properties of direct restorative materials. J Dent Res. 1981;60(5):890.Google Scholar
  39. [39]
    Chimanski A, Cesar PF, Yoshimura HN. Effects of glass chemistry on the optical properties of highly translucent alumina-glass biocomposites for dental restorations. Ceram Int. 2017;43(16):13970.Google Scholar
  40. [40]
    Della Bona A, Nogueira AD, Pecho OE. Optical properties of CAD–CAM ceramic systems. J Dent. 2014;42(9):1202.Google Scholar
  41. [41]
    Jin J, Li XH, Wu JW, Lou BY. Improving tribological and corrosion resistance of Ti6Al4V alloy by hybrid microarc oxidation/enameling treatments. Rare Met. 2018;37(1):26.Google Scholar
  42. [42]
    Zhao DP, Tang JC, Nie HM, Zhang Y, Chen YK, Zhang X, Li HX, Yan M. Macro-micron-nano-featured surface topography of Ti–6Al–4V alloy for biomedical applications. Rare Met. 2018;37(12):1055.Google Scholar
  43. [43]
    Yuan BG, Yu HP, Li CF, Sun DL. Wear properties of nonhydrogenated, hydrogenated, and dehydrogenated Ti6Al4V alloy. Rare Met. 2018;37(7):574.Google Scholar
  44. [44]
    Cesar PF, Yoshimura HN, Miranda WG Jr, Miyazaki CL, Muta LM, Filho LER. Relationship between fracture toughness and flexural strength in dental porcelains. J Biomed Mater Res B Appl Biomater. 2006;78(2):265.Google Scholar

Copyright information

© The Nonferrous Metals Society of China and Springer-Verlag GmbH Germany, part of Springer Nature 2019

Authors and Affiliations

  1. 1.State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and EngineeringTsinghua UniversityBeijingChina
  2. 2.School and Hospital of StomatologyPeking UniversityBeijingChina

Personalised recommendations