|Year : 2019 | Volume
| Issue : 3 | Page : 149-162
Reinforcing experimental resin-composites with synthesized zirconia and alumina nanofibers: evaluation of cuspal flexure, flexural strength, flexural modulus and fracture toughness
AbdAllah F. El-Saadany1, Samy M. El-Safty BSc 2, Usama M. Abdel Karim2, El-Refaie S. Kenawy3
1 Department of Prosthodontist, Monoufia University Hospital, Shebeen El-Koom, Egypt
2 Department of Dental Biomaterials, Faculty of Dentistry, Tanta University, Tanta, Egypt
3 Department of Chemistry, Faculty of Science, Tanta University, Tanta, Egypt
|Date of Submission||10-Jun-2019|
|Date of Acceptance||29-Aug-2019|
|Date of Web Publication||14-Jan-2020|
Samy M. El-Safty
Department of Dental Biomaterials, Faculty of Dentistry, Tanta University, Tanta
Source of Support: None, Conflict of Interest: None
Aim: This study evaluated cuspal flexure, flexural strength, flexural modulus and fracture toughness of experimental resin-composites reinforced with zirconia and alumina nanofibers. Zirconia and alumina nanofibers were synthesized by wet electrospinning technique. Characterization of these nanofibers was carried out by scanning electron microscopy, X-ray diffraction and Fourier transform infrared spectroscopy that confirmed pure forms of zirconia nanofibers (50–100 nm in diameter) and alumina nanofibers (70–100 nm in diameter).
Materials and methods: The study was divided into seven groups (n = 10/group) according to the type and percentage of added nanofibers: control group, three groups reinforced with 2, 4 and 6 wt% zirconia nanofibers and more three groups reinforced with 2, 4 and 6 wt% alumina nanofibers.
Results: A one-way analysis of variance revealed significant differences between the studied groups (P = 0.000) for all investigated properties. Mean cuspal flexure ranged between 0.13 and 0.36% with markedly lower values for groups with higher percentages of nanofibers. For flexural strength, data ranged between 98 and 168 MPa, for flexural modulus between 7.6 and 17.04 GPa, and for fracture toughness between 1.16 and 3.51 MPa m1/2 with greater values recorded for nanofibers-reinforced groups.
Conclusion: Zirconia nanofibers were superior to alumina nanofibers in reducing cuspal flexure and improving the studied mechanical properties.
Keywords: alumina and zirconia nanofibers, electrospinning, experimental resin-composites
|How to cite this article:|
El-Saadany AF, El-Safty SM, Abdel Karim UM, Kenawy ERS. Reinforcing experimental resin-composites with synthesized zirconia and alumina nanofibers: evaluation of cuspal flexure, flexural strength, flexural modulus and fracture toughness. Tanta Dent J 2019;16:149-62
|How to cite this URL:|
El-Saadany AF, El-Safty SM, Abdel Karim UM, Kenawy ERS. Reinforcing experimental resin-composites with synthesized zirconia and alumina nanofibers: evaluation of cuspal flexure, flexural strength, flexural modulus and fracture toughness. Tanta Dent J [serial online] 2019 [cited 2020 Sep 24];16:149-62. Available from: http://www.tmj.eg.net/text.asp?2019/16/3/149/275935
| Introduction|| |
Resin-based composites (RBCs) have been increasingly used in restorative dentistry as an alternative to dental amalgam because of their color-matching to the tooth structure . Unlike dental amalgam, resin-composite restorations are more conservative because they bond micromechanically to the tooth structure. This prevents unnecessary removal of sound tooth material . Biologically, resin-composites are safer than dental amalgam due to lack of mercury toxicity .
However, technique sensitivity, relative inferior mechanical properties, decreased wear resistance, polymerization shrinkage and higher coefficient of thermal expansion have been reported as some of their shortcomings,. Many efforts have been performed to overcome these undesirable characteristics. Low-shrink resins were manufactured to substitute the conventional dimethacrylate resins to minimize polymerization shrinkage . Potent photointiators were added to improve degree and depth of cure . Increasing filler loading has been considered one of the important strategies applied to enhance mechanical properties and wear resistance of the resin-composite,.
Micro and nano particulate fillers have been the most commonly used filler type as a reinforcing mean of resin-composite materials. However, there are some advantages anticipated for nanofiber fillers over particulate fillers. These include: (i) their large specific surface area, high aspect-ratio (length/diameter) and unique structure which lead to high interfacial bonding force between fibers and resin , (ii) stress distribution and inhibition of crack initiation and propagation , and (iii) capacity to improve some mechanical properties such as flexural strength (FS) and fracture toughness .
Electrospinning is the most commonly applied and controlled method to fabricate nanofibers. It uses an electrical charge to draw nano-sized fibers from a solution. On the contrary to other nanofibers fabrication techniques – that can be applied to a limited number of polymeric materials – electrospinning has been applied to produce polymeric, ceramic and metallic nanofibers from their precursor materials,.
Polymerization shrinkage is an inherent property of dental resin-composites, whether they are self-cured or light-cured. The conversion of monomer to polymer results in a closer and tighter arrangement of molecules leading to shrinkage. Intermolecular distance changes from 0.3–0.4 to 0.15 nm on polymerization of resin-composite material . This polymerization shrinkage builds up stresses that might cause cuspal flexure and bonding failure between resin-composite filling and tooth structure. Cuspal flexure is a biomechanical phenomenon that results from the interaction between polymerization shrinkage stress of the resin-composite and the compliance of the cavity wall. Methacrylate-based polymerizations are particularly susceptible to stress development owing to both the low gel-point conversion as well as the large amount of volume shrinkage per double bond that occurs .
FS is an important mechanical property, particularly for brittle materials,. FS testing includes compressive, tensile and shear stresses testing that the restoration will be mostly subjected to in the oral cavity . Flexural modulus is an intrinsic material property which is directly linked to its composition  and the bonding between atoms . It is considered as a function of many factors such as filler content , monomer chemistry , monomer structure  and filler/matrix interactions . Together with adhesive properties, flexural modulus is a critical factor in microleakage, secondary caries and filling dislodgement .
Fracture toughness can be defined as the stress that a material can withstand prior to failure and represents its ability to resist crack propagation from an existing flaw . Fracture toughness is an important property to evaluate performance of a dental resin-composite in the oral cavity. Dental resin-composites are brittle materials, so that any flaw or crack fairly affects their stress tolerance which leads to catastrophic fracture at stresses below their proportional limit .
The main objectives of this study were to synthesize zirconia and alumina nanofibers by electrospinning technique for the reinforcement of an experimental resin-nano-composite, to characterize the synthesized nanofibers by scanning electron microscopy (SEM), X-ray diffraction (XRD) and Fourier-transform infrared spectroscopy (FTIR), to compare the effect of added zirconia nanofibers with that of alumina nanofibers on some properties of experimental resin-composite; cuspal flexure, FS, flexural modulus and fracture toughness. Our null hypotheses were: (i) adding the electrospun nanofibers to the experimental resin-composite will show no significant effect on the studied properties of resin-composites, and (ii) there will be no significant difference between the effect of zirconia nanofibers and that of alumina nanofibers on examined properties of experimental resin-composites.
| Materials and Methods|| |
Materials used in this study are listed in [Table 1].
|Table 1 Information and percentages of chemical ingredients used in preparing the experimental resin-composite and electrospun zirconia and alumina nanofibers|
Click here to view
Preparation of nanofibers by electrospinning
The instruments necessary for electrospinning include a high voltage supplier, a capillary tube with a small-diameter needle, and a metal collecting screen. An electric field is applied to the end of the capillary tube that contains the polymer solution held by its surface tension and forms a charge on the surface of the solution. As the intensity of the electric field increases, the hemispherical surface of the solution at the tip of the capillary tube elongates to form a conical shape known as the Taylor cone. Upon further increase in the electric field in which the repulsive electrostatic force overcomes the surface tension, the charged jet of the solution is ejected from the tip of the Taylor cone to become very long and thin electrospun nanofibers. Randomly-oriented charged polymer fibers solidify and are collected on top of a positively charged metallic surface,.
Preparation of polyvinyl alcohol (PVA) solution: PVA solution (12 wt%) was prepared by dissolving 12 g PVA granules in 100 ml distilled water at 80–90°C for 2 h in paraffin oil bath with stirring by a magnetic stirrer. (DAIHAN MaXtir 500S High Performance Digital Magnetic Stirrers; SRICO, South Korea). The resulting mixture was further stirred for 24 h at room temperature.
Preparation of zirconia nanofibers: in this study, the method described by Shao et al.  was applied for the synthesis of zirconia nanofibers from zirconium oxychloride octahydrade precursors and PVA as a polymer vehicle. An electrospinning machine (Nano-01A Electrospinning setup; MEC C Co. Ltd, Tokyo, Japan) was used to prepare the nanofibers. For synthesis of zirconia nanofibers, a solution was prepared in the proportions of '20: 1: 2' wt% of PVA solution, zirconium oxychloride octahydrate (ZrOCl·8H2O) and H2O, respectively. Twenty grams of PVA solution were dropped slowly into a solution of 1 g ZrOCl2·8H2O and 2 g of distilled water. The mixing was performed in oil bath at 60°C for 6 h in the magnetic stirrer. A capillary tube carrying this solution with a small-diameter needle (18 G) was connected to the positive terminal of a high-voltage supply (Spellman SL30) that generates 24 kV DC at a flow rate of 0.1 ml/h. The distance between the tip of the needle that ejects the electrospun nanofibers and the aluminum collector was set at 10 cm. The freshly produced nanofibers were dried for 12 h at 70°C in a vacuum dry oven (Vacuum Drying Chambers; Binder, Bohemia, North America) The nanofibers were then calcinated at a rate of 240°C/h until reaching 500°C. Upon reaching 500°C, the nanofibers were left under this temperature for 5 h. The temperature was then increased with the same rate (240°C/h) until reaching 1000°C and then kept for 2 h at this temperature for complete removal of PVA solvent polymer.
Preparation of alumina nanofibers: synthesis of alumina nanofibers from aluminum acetate using PVA as a polymer vehicle is consistent with the method described by Panda and Ramakrishna  who applied electrospinning for the synthesis of Al2O3 nanofibers using aluminum acetate and aluminum nitrate in combination with PVA and polyethylene oxide. Two grams of aluminum acetate (C6H9ALO6) salt was added to 20 g of PVA solution (proportion of 2: 20, respectively). The composition was mixed by a magnetic stirrer in oil bath at 45°C for 12 h. Once again, a capillary tube carrying the solution with a small-diameter needle (18 G) was connected to the positive terminal of a high-voltage supply (Spellman SL30) that generates 20 kV DC at a flow rate of 0.4 ml/h. Nanofibers drying was carried out exactly as in the case of zirconia nanofibers. Calcination was performed at a rate of 240°C/h until reaching a temperature of 700°C then kept constant for 5 h. Then the temperature was raised with the same rate until reaching 1300°C then held constant for 2 h.
Characterization of prepared nanofibers
Characterization of the synthesized nanofibers was carried out by three methods: SEM, XRD and FTIR.
SEM: Nanofiber specimens were covered with gold coating (SPI-Modules Vac/Sputter Coater, Tokyo, Japan) and scanned by an electron microscope (JEOL-JSM-5200LV). Scanning of nanofibers was carried out after calcination at 500°C for zirconia nanofibers and 750°C for alumina nanofibers at a magnification of × 7500 and after calcination at 1000°C for zirconia nanofibers and 1300°C for alumina nanofibers at a magnification of × 20 000.
XRD: The XRD patterns of the ZrO2 and Al2O3 nanofibers were recorded three times: (i) immediately after preparation (before calcination), (ii) after calcination at 500°C for zirconia nanofibers and 750°C for alumina nanofibers, and (iii) after 1000°C for zirconia nanofibers and 1300°C for alumina nanofibers. Each diffraction pattern of these nanofibers was carried out using an XRD diffractometer (XRD, Diffractometer, Model: GNR APD-2000 PRO, EA Almelo, The Netherlands). A rotating x-ray generator (40 kW, 40 mA) with CuKα radiation (wavelength λ=1.540598 Š) was used. XRD profiles of nanofibers were recorded from 4° to 90° at a scanning speed of 2°/min. The peaks obtained from ZrO2 nanofibers were compared to standard references in Joint Committee for Powder Diffraction Standards (JCPDS) file available in software (card no. 89-2843)  for zirconium oxide nanocrystals and films. The peaks obtained from Al2O3 nanofibers were compared to standard references according to (JCPDS card no. 42-1468)  for aluminum oxide nanocrystals and films.
FTIR: Using the FTIR method (FTIR, model: EQUINO X55; Bruker, Germany) the functional groups of the PVA, ZrO2 and Al2O3 in nanofibers were identified three times: firstly, immediately after preparation, secondly, after calcination at 500°C for zirconia nanofibers and 750°C for alumina nanofibers and thirdly, after calcination at 1000°C for zirconia nanofibers and 1300°C for alumina nanofibers. Nanofibers were milled in mortar and pestle then added to potassium bromide powder at a ratio of 1: 80, respectively. The mixture was pressed under hydraulic press to form a tablet. Ten scans were recorded for each tablet between wave numbers of 5000–200/cm with resolution of 1/cm.
Ball-milling of nanofibers
Nanofibers were produced in the form of sheets. To obtain the desired size (<100 nm), nanofibers were ball-milled. Ball-milling of nanofibers was conducted by a planetary photon grinder milling machine (Retsch – PM 400; Haan, Germany) with a ball size of 10 mm, at a speed of 350 rpm, for 7 h.
Silanization of nanoparticles and nanofibers
Inorganic nanofibers (ZrO2 and Al2O3 nanofibers) and inorganic fillers (silica and zirconia nanoparticles) were silanized with 5 wt% silane coupling agent '3-methacryloxypropyltrimethoxysilane' and 95 wt% acetone solvent . Firstly, each type of nanofibers and nanoparticles was dispersed in acetone in a separate bottle. Then, the silane coupling agent was added to each solution at a percent of 5 wt% at room temperature. The mixture was continuously stirred at a speed of 150 rpm for 15 h in the magnetic stirrer. Each solution was then filtered in order to collect the modified nanofibers and nanoparticles. Secondly, the mixture was stored for 24 h at 37°C to assure the complete removal of the solvent. After storage, the mixtures were sieved through a 300 nm then 100 nm sieve and kept in a sonication device for 10 min (Power sonic 405; Hwashin Technology Co., Korea). Finally, the obtained nano-sized fibers and particles were dried in an oven at 110°C for 3 h under vacuum.
Preparation of an experimental resin-composite
Preparation of the resin matrix of the experimental resin-composite was carried out according to the method and proportions described by Asmussen and Peutzfeldt . The organic matrix that forms 27 wt% of the resin-composite was prepared by mixing bisphenol A diglycidyl methacrylate (Bis-GMA), urethane dimethacrylate (UDMA) and bisphenol A glycidyl methacrylate ethoxylated (Bis-EMA) as the forming monomers. Triethylene glycol dimethacrylate (TEGDMA) was added as a diluent comonomer and polyethylene glycol dimethacrylate (PEGDMA) as a cross-linking agent. Firstly, TEGDMA and PEGDMA were mixed in proportions of 50: 50 (wt%) to form a diluent solution. Mixing process was performed by Vortex mixer (Vortex, Ika; Sigma Aldrich, St Louis, Missouri, USA) at a speed of 100 rpm for 6 h at room temperature. Then, TEGDMA/PEGDMA mixture was mixed with UDMA and Bis-EMA in proportions of 1: 1: 1 wt% at 100 rpm for 24 h at room temperature. Finally, 30 wt% of Bis-GMA was added to 70 wt% of the previously prepared solution (TEGDMA/PEGDMA)/UDMA/Bis-EMA) and the solution was mixed at 200 rpm for 24 h at room temperature.
The filler component of the experimental nanocomposite was prepared according to previous published works,. Silanized ZrO2 nanoclusters (20 wt%) were mixed to 80 wt% of silanized ZrO2 nanoparticles in a mechanical stirrer with a vertical blade (5040001 RW28; Atlanta, USA) at a speed of 20 rpm for 2 h at room temperature to form ZrO2 nanofiller mixture. This process was repeated for silica nanoclusters/nanoparticles in the same proportions and conditions to form silica nanofiller mixture. Twenty percent of ZrO2 nanofillers mixture and 80 wt% of SiO2 nanofillers mixture were mixed in the mechanical stirrer at 20 rpm for 5 h at room temperature to form the filler component of the experimental resin-composite.
Then the mixture was divided to seven groups; one group without nanofibers. Other six groups were formed by adding zirconia or alumina nanofibers with 2, 4 and 6% at the expense of the nanoparticles. The added nanofibers were thoroughly mixed with the nanoparticles in the mechanical stirrer at 50 rpm for 6 h at room temperature to ensure even distribution of the nanofibers throughout the fillers.
The photoinitiator (camphorquinone) and coinitiator [ethyl-4-(N, N′-dimethylamino) benzoate] were mixed in proportions of 1: 1 wt% in the mechanical stirrer at 10 rpm for 1 h at room temperature to form the photo-activation system. Finally, each group of the experimental resin-composite was prepared in proportions of 27 wt% organic matrix, 72 wt% fillers and 1% photo-activation system. Mixing process was achieved in a centrifugal mixing device (Speed-Mixer, DAC 150 FVZK; Hauschild Engineering, Hamm, Germany).
Grouping: Experimental resin-composite groups were formulated and classified according to the type and percentage of added inorganic nanofibers as follows:
- Group I: Experimental resin-composite without nanofibers (control group).
- Group II: Experimental resin-composite reinforced with 2 wt% Al2O3 nanofibers.
- Group III: Experimental resin-composite reinforced with 4 wt% Al2O3 nanofibers.
- Group IV: Experimental resin-composite reinforced with 6 wt% Al2O3 nanofibers.
- Group V: Experimental resin-composite reinforced with 2 wt% ZrO2 nanofibers.
- Group VI: Experimental resin-composite reinforced with 4 wt% ZrO2 nanofibers.
- Group VII: Experimental resin-composite reinforced with 6 wt% ZrO2 nanofibers.
The investigated properties were: cuspal flexure, FS, flexural modulus and fracture toughness. For each property, 10 specimens (n = 10) were prepared and tested for each group. The curing unit intensity output was monitored using a light meter after curing each group for each property (Cromalux, MEGA-PHYSIK; GmbH and COKG, Germany).
Cuspal flexure: Approval for this research was obtained from Research Ethics Committee, Faculty of Dentistry, Tanta University. The purpose of this study was explained to the patients and informed consents were obtained to use their extracted teeth in the research according to guidelines on human research adopted by Research Ethics Committee, Faculty of Dentistry, Tanta University.
A total of 70 maxillary premolars were collected from Oral surgery and Orthodontic Departments. The teeth were thoroughly cleaned and free of caries, hypoplastic defects and cracks. The maximum bucco-palatal width of each maxillary premolar was nearly equal. The maxillary premolars were collected and stored in 0.5% chloramine solution at 23 ± 1°C until they were submitted to cavity preparation. Each tooth was prepared using a high speed handpiece with water cooling for a large standardized class II mesio-occlusal-distal cavity using a parallel diamond bur (Miltex RA #556; Miltenberg Inc., New York, New York, USA). The width of the proximal box was two-thirds of the bucco-palatal width. A gingival seat of nearly 1 mm was prepared and extended above the cemento-enamel junction at the cervical aspect. The occlusal isthmus was prepared to the half of the bucco-palatal width and the cavity depth at the occlusal isthmus was standardized to 3 mm. The width of occlusal box was 4/5 of the intercuspal width. The cavo-surface margins were prepared at 90° and all the internal line angles were rounded. Periodontal probe was used to verify preparation dimensions during cavity fabrication,.
After cavity preparation, the teeth were mounted in a self-curing acrylic resin (PMMA; Esschem Co., Linwood, Pennsylvania, USA). The digital image correlation method was used for measuring cuspal flexure which analyzes the displacement of an object using a USB digital microscope with a built-in camera (COMET xS; Steinbichler Optotechnik GmbH, Neubeuern, Germany). Two reference points were selected on buccal and palatal cusps of each tooth. After that, the mounted teeth were scanned using an optical scanner (COMET xS; Steinbichler Optotechnik GmbH) before restoration and the initial length between the reference points (L0) was measured.
The 70 maxillary premolars were divided into seven groups (n = 10); control group and six experimental groups according to the type and percentage of the added nanofibers. The tooth surfaces were etched and prepared for bonding by the bonding system (Adper Single Bond Plus Adhesive Refill, 51102, 3 M; ESPE, USA) according to the manufacturer's instructions. The resin-composite material was packed into the cavity in horizontal increments with 2 mm thickness for each. Each increment was light-cured for 20 s using a light curing unit (Elipar TM Deep Cure, 3 M; ESPE) with light irradiance of 1200 mW/cm2. The restored teeth were scanned 10 min after restoration. The distance between the reference points (Lfinal) after curing restoration (deformation; deformed image) was recorded. The cuspal flexure was computed as the difference between the final and initial measurements, ΔL = Lfinal − L0. Percentage of cuspal flexure was calculated by: cuspal flexure%=(ΔL/L0)×100.
FS:FS (Ef) testing was performed according to ISO 4049:2009 by three-point bending test . Seventy bar-shaped specimens (25 mm length × 2 mm width × 2 mm height) were prepared in half-split stainless steel molds. The mold was adapted to a glass slide covered with Mylar strip and a separating medium was applied to the mold walls with a brush. The resin-composite material was packed to the mold cavity by a plastic instrument. After that, another glass slide covered with Mylar strip was applied on top of the mold. Curing was performed on three overlapping points (20 s each) on top and bottom of the specimen to ensure complete curing of the whole specimen using the same light curing unit described above. After polymerization, specimens were removed from the mold and stored in distilled water at 37°C for 24 h before being submitted to testing.
FS testing was performed in a Universal Testing Machine (model 3365; Instron, High Wycombe, UK). The test assembly consisted of two supporting wedges placed 20 mm apart and a loading wedge that applied the load at a cross-head speed of 0.75 mm/min. The applied force and strain during the bending was measured as a function of deflection. The FS in MPa was calculated according to the following formula: FS = 3 Fd/2wh2, where F = maximum force (N), d = distance between the two supports, w = width of the specimen and h = height of the specimen (mm).
Flexural modulus calculation:flexural elastic modulus (E) in GPa was calculated after recording the FS by the following equation: E = FL3/ 4BH3d, where F is the maximum force (N), L is the distance between the supports, B is the width of the specimen, H is the height of the specimen, and d is the deflection (mm).
Fracture toughness: Fracture toughness testing was conducted according to ISO/FDIS 6872:2007 . Seventy bar-shaped specimens (25 mm length × 5 mm width × 2 mm depth) were prepared in half-split stainless steel molds. Packing of the resin-composite material, curing and storage of specimens were done as described for the FS specimens. A single edge V-notched beam at midspan was prepared in each specimen with a disc 0.5 mm width (3014; Patterson Dental, New York, New York, USA). The notch depth was about one-third the specimen depth. Ruler and pen marker were used to standardize the position, direction and length of the notch.
Fracture toughness was determined on a single-edge notch specimen using the three-point bending method according to the procedures outlined in ISO/FDIS 6872, 2007. A three-point bending test (span length = 20 mm with the notch centrally located on the tensile side) was done in a Universal Testing Machine (model 3600; Instron Industrial Products, Norwood, Massachusetts, USA) with a load cell of 5 kN. Specimens were loaded until fracture at a cross-head speed of 0.5 mm/min. The load-deflection curves were recorded using computer software. (Instron Bluehill Lite software) Calculation of fracture toughness (Griffith = MPa.m1/2) by the single-edge notched method and was done according to the following equation:
where F = maximum load, B = specimen width, S = supporting span, w = specimen height, c = notch length, f (c/w)=a function of c and w.
Statistical analysis: Data were collected and tabulated and statistically analyzed by an IBM compatible personal computer with SPSS version 20 (released 2011; SPSS Inc., Armonk, New York, USA). Data for cuspal flexure, FS, flexural modulus and fracture toughness of all studied groups were analyzed using a one-way analysis of variance with the significance level established at (P ≤ 0.05). The post-hoc Tukey test was used to determine differences between groups for each studied property.
| Results|| |
Scanning electron microscopy
Zirconia (ZrO2) nanofibers: SEM image [Figure 1]a, at a magnification of × 7500, shows smooth uniform surfaces of the as-synthesized zirconium oxide/PVA fibers that were calcinated at 500°C. The diameters of the fibers were in the range of 200–300 nm. ×At 20 000 magnification, ZrO2 nanofibers obtained after calcination at 1000°C [Figure 1]b appeared to have rough surface because of crystallization of ZrO2 at this high temperature. Calcination at this temperature caused the decomposition and complete removal of PVA that played the role of a template for fiber formation during electrospinning process. The diameters of the produced ZrO2 nanofibers ranged between 50 and 100 nm.
|Figure 1: Scanning electron microscopy (SEM) image (×7500) showing zirconium oxide/polyvinyl alcohol nanofibers after calcination at 500°C (a) and SEM image (×20 000) showing pure zirconia (ZrO2) nanofibers after calcination at 1000°C (b).|
Click here to view
Alumina (Al2O3) nanofibers: At a magnification of × 7500, SEM [Figure 2]a shows the as-synthesized aluminum acetate/PVA fibers that were calcinated at 700°C having smooth uniform surfaces with varying diameters that are in the range of 250–400 nm. Upon calcination at 1300°C, the Al2O3 nanofibers appeared with a beaded structure due to loss of the PVA organic polymers leaving pure alumina nanofibers. The diameters of the nanofibers were reduced due to calcination to be in the range of 70–100 nm [Figure 2]b.
|Figure 2: Scanning electron microscopy (SEM) image (×7500) showing aluminum acetate/polyvinyl alcohol nanofibers after calcination at 700°C (a) and SEM image (×20 000) showing pure alumina (Al2O3) nanofibers after calcination at 1300°C (b).|
Click here to view
At a magnification of × 20 000, [Figure 3]a shows SEM of nanofibers mat. [Figure 3]b shows SEM of ball-milled nanofibers at a magnification of × 50 000.
|Figure 3: Scanning electron microscopy image showing nanofibers mat (a) and ball-milled nanofibers (b).|
Click here to view
ZrO2 nanofibers: An amorphous behavior was observed in the XRD of the as-synthesized ZrO2/PVA fibers before calcination [Figure 4]a. This means that nanofibers contain crystalline (ZrO2) and noncrystalline (PVA) structures. The samples that were obtained after calcination at 500°C exhibited diffraction peaks centered at 30.358, 35.201, 50.508 and 64.408, that are characteristic of the tetragonal ZrO2 suggesting that the sample present mainly in the metastable tetragonal phase [Figure 4]b. This result is in accordance with the data reported in case of nano-crystalline zirconia in JCPDS (card no. 89-2843) . However, peaks that were centered at 24.650 and 28.405 revealed the co-existence of some content of monoclinic crystalline phase as well. Upon raising the calcination temperature to 1000°C, diffraction peaks corresponding to the monoclinic zirconium oxide phase became sharper and more intense, emphasizing the presence of monoclinic phase [Figure 4]c. This indicates the gradual transformation of the metastable tetragonal phase to the monoclinic one. Diffraction peaks corresponding to the tetragonal and monoclinic phases were well-matched with database in JCPDS (card no. 89-2843) for ZrO2 nanocrystals and films .
|Figure 4: X-ray diffraction peaks of: (a) electrospun ZrO2/polyvinyl alcohol nanofibers before calcination, (b) electrospun ZrO2 nanofibers calcinated at 500°C and (c) electrospun ZrO2 nanofibers calcinated at 1000°C.|
Click here to view
Al2O3 nanofibers: Before calcination, an amorphous behavior was observed in the XRD of the as-synthesized Al2O3/PVA fibers [Figure 5]a. At 700°C, the peak of γ-Al2O3 was displayed with very poor diffraction intensity [Figure 5]b. At 1300°C [Figure 5]c, α-alumina phase was formed. This was indicated by diffraction peaks 45.24 and 68.26 according to (JCPDS card no. 42-1468) . From the above results, the phase transition of alumina nanofibers in this study can be described as follows: amorphous/microcrystalline→γ-Al2O3→α-Al2O3 due to the process of heat treatment.
|Figure 5: X-ray diffraction peaks of: (a) electrospun Al2O3/polyvinyl alcohol nanofibers before calcination, (b) electrospun Al2O3 nanofibers calcinated at 700°C and (c) electrospun Al2O3 nanofibers calcinated at 1300°C.|
Click here to view
Fourier transform infrared spectroscopy
ZrO2 nanofibers: In order to confirm the formation of pure zirconium oxide nanofibers after calcination at high temperature with the complete removal of PVA, FTIR spectrum of the various nanofibers samples at different calcination temperatures was recorded. The FTIR spectrum of the as-synthesized inorganic/organic hybrid nanofibers [Figure 6]a showed some strong absorption of the infrared waves in the region of 2000–1000/cm, which corresponds to stretching and bending vibrations of PVA molecule. After calcination at 500°C, the intensity of those peaks was decreased, indicating the beginning of removal of most PVA molecules from the fibers, and different peaks appeared around 750 and 550/cm due to the formation of zirconium oxide [Figure 6]b. The formation of pure zirconium oxide was indicated by the IR spectra of the samples calcinated at 1000°C, which displayed intense peaks at 750 and at 520/cm [Figure 6]c due to Zr-O bond stretching. In addition, disappearance of the absorptions corresponding to the PVA molecule indicated the complete removal of them at this temperature and the nanofibers formed were exclusively consisting of zirconium oxide.
|Figure 6: Fourier transform infrared spectroscopy spectra of: (a) electrospun ZrO2/polyvinyl alcohol nanofibers before calcination, (b) electrospun ZrO2 nanofibers calcinated at 500°C and (c) electrospun ZrO2 nanofibers calcinated at 1000°C.|
Click here to view
Al2O3 nanofibers: As in the case of ZrO2, FTIR spectrum of the as-synthesized inorganic/organic hybrid nanofibers [Figure 7]a was in the region of 2000–1000/cm. At 700°C, some of Al2O3 fibers appeared [Figure 7]b. Three characteristic peaks at 634, 581, and 777/cm for alumina nanofibers calcinated at 1300°C were confirmed. Observation of α-phase alumina indicates Al-O bond bending and stretching [Figure 7]c.
|Figure 7: Fourier transform infrared spectroscopy spectra of: (a) electrospun Al2O3/polyvinyl alcohol nanofibers before calcination, (b) electrospun Al2O3 nanofibers calcinated at 700°C and (c) electrospun Al2O3 nanofibers calcinated at 1300°C.|
Click here to view
Results of investigated properties: Mean data and SD of cuspal flexure, FS, flexural modulus and fracture toughness of all investigated groups are listed in [Table 2]. The experimental resin-composite group without nanofibers (control group) exhibited the greatest cuspal flexure (0.36%). The reinforced groups showed lower values compared to the control group with the lowest values recorded for reinforced groups with 6% (0.13%), 4% (0.21%) zirconia nanofibers and 6% alumina nanofibers (0.24%) followed by 2% zirconia (0.26%), 4% alumina (0.28%) and 2% alumina (0.31%) nanofibers. For FS, data ranged between 98 MPa for control group and 168 MPa for resin-composite reinforced with 6% zirconia nanofibers. For a given percentage of nanofibers, reinforcement with zirconia nanofibers revealed greater FS values than reinforcement with alumina nanofibers. This was clear with 2% nanofibers for both zirconia and alumina (126 and 122 MPa, respectively), with 4% (146 and 138 MPa, respectively) and with 6% (168 and 154 MPa, respectively).
|Table 2 Mean and SD of cuspal flexure (%), flexural strength (MPa), flexural modulus (GPa) and fracture toughness (MPa.m1/2) for all studied groups|
Click here to view
The results for flexural modulus were exactly the same for the FS in terms of ranking of studied groups where the greatest value was recorded by resin-composite reinforced with 6% zirconia nanofibers (17.04 GPa) and the lowest value was shown by the control group (7.60 GPa). Similarly, for the same percentage of reinforcing nanofibers, zirconia nanofibers recorded greater results than those recorded by alumina nanofibers as can be seen in [Table 2].
For fracture toughness, the lowest mean values were recorded by the control group (1.16 MPa m1/2) followed by groups reinforced with 2% alumina (1.81 MPa m1/2), 2% zirconia (1.99 MPa m1/2), 4% alumina (2.42 MPa m1/2), and 4% (2.78 MPa m1/2) zirconia nanofibers. The highest fracture toughness was recorded by the group reinforced with 6% zirconia nanofibers (3.51 MPa m1/2) followed by that reinforced with 6% alumina nanofibers (2.95 MPa m1/2).
One-way analysis of variance revealed significant differences between the studied groups; for cuspal flexure (P = 0.000), for FS (P = 0.000), for flexural modulus (P = 0.000) and for fracture toughness (P = 0.000).
Post-hoc Tukey test revealed significant differences between all studied groups for cuspal flexure except between the control group and that reinforced with 2% alumina nanofibers (P = 0.102), between groups reinforced with 2 and 4% alumina nanofibers (P = 0.630), between groups reinforced with 4 and 6% alumina nanofibers (P = 0.243), between groups reinforced with 4% alumina and 2% zirconia nanofibers (P = 0.769), between groups reinforced with 6% alumina nanofibers and both groups reinforced with 2 (P = 0.974) and 4% (P = 0.370) zirconia nanofibers and between groups reinforced with 2 and 4% zirconia nanofibers (P = 0.060).
For FS, there were significant differences between all studied groups except between the groups reinforced with 2% alumina and 2% zirconia nanofibers (P = 0.486). For flexural modulus, there were significant differences between all investigated groups. For fracture toughness, there were significant differences between all studied groups except between groups reinforced with 2% alumina and 2% zirconia nanofibers (P = 0.211) and between groups reinforced with 6% alumina and 4% zirconia nanofibers (P = 0.252).
| Discussion|| |
Majority of research studies employ commercial RBCs for comparative evaluation. However, interpretation of significant findings from such studies is mostly handicapped both by many variations in the formulations of commercial RBCs, and the fact that manufacturers do not reveal precise details of formulation differences in their products. In such a case, it may be difficult to identify the most specific component causing variation in the studied properties. In the current study, the use of experimental RBCs with controlled formulations is very critical as this allows systematic investigation of variables controlling RBC behavior, thus allowing hypothesis testing of fundamental concepts,,. Besides, with commercial formulations, substitution of an amount of nanoparticles with nanofibers is not possible. Moreover, addition of nanofibers to commercial formulations could not be possible as their consistency will be changed and become difficult in manipulation.
SEM is an essential investigation to show the morphology of nanofibers and their diameters . Scanning was carried out at two levels of magnification; ×7500 and × 20 000. These two levels of magnification were selected for morphological clarification of the synthesized nanofibers. A relatively clear form of a mixture of PVA/nanofibers can be identified at a magnification of × 7500. Pure forms of the synthesized nanofibers at the nano-scale can be clearly identified at a magnification of × 20 000 . For each scanning, XRD was used to measure structural make-up, content and crystallinity of structures . FTIR is an excellent test to investigate presence of target materials and check out complete removal of the polymer solvent .
Zirconia and alumina are bioinert ceramics. Zirconia is characterized by high strength properties and fracture toughness  and can be found in three crystalline forms: monoclinic (M), tetragonal (T) and cubic (C). During heating, zirconia undergoes a transformation process. The monoclinic form is stable from room temperature up to 1170°C. Above this temperature, it transforms into a denser tetragonal phase with a 5% volume decrease. The tetragonal form is stable between 1170 and 2370°C. At temperatures from 2370 to 2680°C, zirconia takes a cubic crystal structure. A reversible transformation from the tetragonal to the monoclinic form occurs during cooling at a temperature of about 100°C under 1070°C, with 3–4% volumetric expansion and generation of forces that cause cracks closure in the ZrO2 ceramics . Alumina is characterized by having high abrasion resistance and chemical inertness . Alumina, with either nanometer-sized or micrometer-sized grains is promising as a potential substitute as a strengthening agent. In this study, selection of aluminum acetate over aluminum nitrate was because it was found that in case of aluminum nitrate, the resulting nanofibers were highly hygroscopic due to the presence of nitrate anions with some difficulty in collection due to the strong repulsion the nitrate anions .
Investigated properties and tested hypotheses
As the addition of zirconia and alumina nanofibers showed significant effects on the studied resin-composite properties, the first null hypothesis was rejected. For cuspal flexure, the control group recorded significantly greater percent than all groups reinforced with both ZrO2 and Al2O3 nanofibers (P = 0.000) except that reinforced with 2% Al2O3 nanofibers (P = 0.102). Cuspal flexure% was clearly minimized with increasing the percent of added nanofibers. This reduction might be attributed to the strong adhesion between resin and silanated nanofibers in the resin-composite because of the large surface area and high aspect ratio . These porous, nonwoven networks have high fiber interconnectivity. Moreover, presence of nanofibers in the molecular level increases their reactivity .
The cuspal flexure exhibited by all groups (0.13–0.36%) was lower than the acceptable percent for a resin-composite material (1–2%) as reported in the literature,. Reduction of polymerization shrinkage has been reported to be accomplished by: (i) addition of nanofillers (particles and fibers) that increase the filler loading in the prepared resin-based materials , (ii) presence of aliphatic UDMA in resin composition which has some advantages over Bis-GMA including its lower molecular weight as well as lighter viscosity due to the presence of flexible urethane or carbamate linkages and absence of aromatic groups  and (iii) PEGDMA was added instead of a portion of the TEGDMA resin to moderate the shrinkage resulting from the very low molecular weight TEGDMA . In addition, it can be thought that the decreased polymerization shrinkage for nanofibers-reinforced groups is due to the wide distribution of the nanofibers that effectively fill the interstitial spaces between polymer chains reducing the shrinkage. The enhanced bonding that exists between silanized nanofibers and resin may be another cause of this reduced shrinkage.
The decreased cuspal flexure% of nanofibers-reinforced resin-composites in this study is consistent with the study conducted by Garoushi et al.  who reported a significant reduction in polymerization shrinkage of microfibers-reinforced resin-composites (posterior ever × composite) in comparison with different commercial posterior resin-composites (Alert, Tetric Evo Ceram Bulk Fill, Voco X-tra base, SDR, Venus Bulk Fill, SonicFill, Filtek Bulk Fill, Filtek Superme, and Filtek Z250). The authors attributed this reduction to the strong bonding between resin and silanated fibers. In another study , it was reported that the decreased shrinkage stress of microfibers-reinforced resin-composite was due to strong bonding between resin and silanized fibers as well as the horizontal direction of fibers during condensation.
For FS and modulus, the nanofibers-reinforced groups recorded significantly greater mean values than the control group (P = 0.000). FS and modulus mean data for all reinforced groups showed direct positive correlation [Table 2] between the recorded values and the percent of added nanofibers. Enhanced FS and modulus for the nanofibers-reinforced groups can be reasoned, in addition to the characteristics of reinforcing nanofibers mentioned earlier, by the outstanding mechanical performance and stabilization of ceramic nanofibers . The small size and excellent structure of ceramic nanofibers could provide a good distribution in the resin matrix (vertical stirring). As the diameter of fibers is reduced, most of the ions, molecules and functional groups will be available on the outmost layer which can grant high reactivity to the nanofibers that are not found in their traditional bulk counterparts. Nano-scaled fibers strength is over ten times as high as that of most of micro-scaled fibers. Because of such advantages of ceramic nanofibers, when a dental resin-composite undergoes a load or pressure, ceramic nanofibers have a great potential to inhibit the emergence of microcracks in the material and prevent their enlargement .
In addition, the higher FS values in ZrO2 nanofibers-reinforced groups than those reinforced with Al2O3 nanofibers can be attributed to the fact that ZrO2 has FS (900–1400 MPa) that is nearly twice that of Al2O3 (450 MPa) . Also, this finding can be supported by the microstructural differences between ZrO2 and Al2O3 where the former has a higher density (6 g/cm3) than the latter (3.9 g/cm3) . FS and modulus SDs of all reinforced groups in the study were small. This could be attributed to the even distribution of nanofibers throughout resin matrix due to proper mixing in a mechanical stirrer with a vertical blade. A uniform dispersion of the fibers throughout the matrices is very critical. This is because agglomeration of the nanofibers forms a stress concentration area which is the main reason for creating small cracks in the matrix, thus damaging the properties of the material,. Therefore, to produce the biggest effect of chemical bonding, nanofibers should be dispersed as uniform as possible.
Findings of this study for FS and modulus are in agreement with many studies that applied different nanofibers to reinforce dental resin-composites. Tian et al.  reported that the addition of 3% fibrillar silicate into nylon 6 nanofibers could enhance mechanical performances of dental resin-composites, while large amount of fibers showed no improvement because of lack of interfacial linking at the interface between nanofibers and the dental matrix. This can be supported by findings of a study conducted by Chen et al.  who reported that with 10 wt% hydroxyapatite 'HAP' nanofibers, biaxial FS of dental resin-composites raised by 22.2%, while with 40 wt% HAP nanofibers, the biaxial FS decreased by 60%. Dodiuk-Kenig et al.  reported that addition of PVA 'PV-OH' nanofibers could enhance the resin-composite mechanical properties. They thought that the effective reinforcement of resin-composites with nanofibers mostly lies in the close linking between resin matrix and nanofibers. This was proved by comparing the results recorded by treated nanofibers and those of untreated ones.
It was reported  that the addition of 1 wt% glass nanofibers (with diameters of about 400 nm) to a dental resin-composite, the tensile strength, elastic modulus, and work of fracture raised by 12, 33, and 52%, respectively. Moreover, Tian et al.  reported that with 1 wt% nano-scaled fibrillar silicate, the modulus of elasticity, bending strength, and work of fracture increased by 16.7, 40, and 78.4%, respectively. Guo et al.  used zirconia–silica and zirconia–yttria–silica nanofibers for reinforcing resin-composites. The percents of fibers added were 2.5, 5 and 7.5 wt%. They found that reinforcement with zirconia–silica or zirconia–yttria–silica nanofibers of 2.5% or 5.0 wt% significantly increased the FS, flexural modulus and energy at break over the control. Further increase in the added nanofibers (7.5 wt%), however, recorded lower results than those reported by 2.5% and 5.0 wt%, although they were still higher than those of the control.
With regard to fracture toughness, all nanofibers-reinforced groups exhibited significantly greater mean values than the control group, (P = 0.000), with better results for the groups reinforced with ZrO2 nanofibers than those reinforced with Al2O3 nanofibers. This can be attributed to that zirconia has higher fracture toughness (7–10 MPa.m1/2) than alumina (5–6 MPa.m1/2) . Similar to the FS and modulus, the results of fracture toughness [Table 2] showed direct positive correlation with the added nanofibers of both types. The enhanced fracture toughness of the reinforced groups over the control one can be explained on the basis that nanofibers can provide a larger area for load transfer and promote toughening mechanisms such as fiber bridging and fiber pullout . Under high external pressure, microcracks tend to be created in the body of resin-composites. However, the room of cracks is not empty, nanofibers still exist in the middle of cracks plane and function across planes. Thereby, cracks are prevented via the remaining fillers. Hence, much of the stress is transmitted through the fillers into the resin, and that is the reason why the nanofibers must be largely tougher than the resin matrices .
Findings of our study regarding fracture toughness are in consistence with a study  reporting higher fracture toughness of nylon 6 nanofibers-reinforced resin-composites compared to the control group. This outcome could be attributed to the high specific surface area and high aspect ratio of nanofibers which increased the interfacial adhesion between nylon 6 nanofibers and the resin matrix. The study showed that microcracks could be efficiently prevented in the presence of nanofibers. Also, in the study mentioned earlier , the fracture toughness was significantly increased by addition of 2.5–5 wt% of zirconia–silica and zirconia–yttria–silica nanofibers but decreased with 7.5 wt%. In this study, choosing the percentages of 2, 4, and 6% of zirconia and alumina nanofibers was based on the findings of the previous study , where addition of 2.5–5% of nanofibers did significant improvement in composite properties, while increasing the percentages of these fibers up to 7.5% resulted in an adverse effect on the studied properties.
As there were significant differences between the ZrO2 nanofibers-reinforced groups and Al2O3 nanofibers-reinforced groups in most of studied properties, even with equal percentage of added nanofibers, the second null hypothesis was rejected as well. Superiority of ZrO2 nanofibers-reinforced groups to that of Al2O3 nanofibers-reinforced groups can be attributed to the superior characteristics that ZrO2 possesses such as great FS and modulus, high density and enhanced fracture toughness than Al2O3. The metastable tetragonal crystalline structure at room temperature is considered the main reason for the superior fracture strength of ZrO2. This structure represents an efficient mechanism against flaw propagation and has a strong impact against subcritical crack growth,.
Monoclinic phase is a more stable phase. One of the most important properties is a remarkable increase in fracture toughness of the material by hindering, but not preventing, the propagation of a crack; compressive stress concentration converts the transformation from the tetragonal phase to the monoclinic one. Increasing the crystal volume, constrained by the surrounding ones, leads to a favorable compressive stress that limits growth of cracks. It was reported that phase transformation toughening is the most probable mechanism for exceptional FS and fracture toughness of ZrO2 among all other ceramics,.
| Conclusion|| |
Synthesized zirconia and alumina nanofibers proved to be effective reinforcing fillers to the dental resin-composites. Compared to the control group, these fibers reduced the cuspal flexure and enhanced the FS, flexural modulus and fracture toughness. Within the range of added nanofibers (2–6 wt%) of both types, there was a strong correlation between percent of added nanofibers and increase of mechanical properties as well as reduction of cuspal flexure of experimental resin-composites. When used in equal percent, there was an obvious superiority of ZrO2 compared to Al2O3 nanofibers in improving the resin-composite properties.
Financial support and sponsorship
Conflicts of interest
There are no conflicts of interest.
| References|| |
Opdam N, Bronkhorst E, Loomans B, Huysmans M. 12-year survival of composite vs. amalgam restorations. J Dent Res 2010; 89:1063–1067.
Manhart J, Chan H, Hamm G, Hickel R. Review of the clinical survival of direct and indirect restorations in posterior teeth and permanent dentition. Oper Dent 2004; 29:481–508.
Leistevuo J, Leistevuo T, Helenius H. Dental amalgam fillings and the amount of organic mercury in human saliva. Caries Res 2001; 35:163–166.
Ferracane JL. Resin composite-state of the art. Dent Mater 2011; 27:29–38.
Jun SK, Kim DA, Goo HJ, Lee HH. Investigation of the correlation between the different mechanical properties of resin composites. Dent Mater J 2013; 32:48–57.
Tantbirojn D, Pfeifer CS, Braga RR, Versluis A. Do low-shrink composites reduce polymerization shrinkage effects? J Dent Res 2011; 90:596–601.
Moszner N, Salz U. Recent developments of new components for dental adhesives and composites. Macro Mater Eng 2007; 3:245–271.
Kim H, Ong J, Okuno O. The effect of filler loading and morphology on the mechanical properties of contemporary composites. J Prosthet Dent 2002; 87:642–649.
Ruddell D, Maloney M, Thompson J. Effect of novel filler particles on the mechanical and wear properties of dental composites. Dent Mater 2002; 18:72–80.
Venugopal J, Ramakrishna S. Applications of polymer nanofibers in biomedicine and biotechnology. Appl Biochem Biotechnol 2005; 125:147–157.
Cipitria A, Skelton A, Dargaville T, Dalton P, Hutmacher D. Design, fabrication and characterization of PCL electrospun scaffolds – a review. J Mater Chem 2011; 21:9419–9453.
Tian M, Gao Y, Liu Y, Yiliang L, Xu R, Hedin N, et al
. Bis-GMA/TEGDMA dental composites reinforced with electrospun nylon 6 nanocomposite nanofibers containing highly aligned fibrillar silicate single crystals. Polymer (Guildf) 2007; 48:2720–2728.
Wang S, Fu D, Li X. Functional polymeric nanofibers from electrospinning. Recent Pat Nanotechnol 2009; 1:21–31.
Reneker D, Yarin A, Zussman E, Xu H. Electrospinning of nanofibers from polymer solutions and melts. Advan App Mech 2007; 41:43–65.
Ferracane JL. Developing a more complete understanding of stresses produced in dental composites during polymerization. Dent Mater 2005; 21:36–42.
Kleverlaan CJ, Feilzer AJ. Polymerization shrinkage and contraction stress of dental resin composites. Dent Mater 2005; 21:1150–1157.
Palin W, Fleming G, Marquis P. The reliability of standardized flexure strength testing procedures for a light-activated resin-based composite. Dent Mater 2005; 21:911–919.
Junior R, Adalberto S, Zanchi CH, Carvalho RVD, Demarco FF. Flexural strength and modulus of elasticity of different types of resin-based composites. Braz Oral Res 2007; 21:16–21.
Feilzer A, Dauvillier B. Effect of TEGDMA/BisGMA ratio on stress development and viscoelastic properties of experimental two-paste composites. J Dent Res 2003; 82:824–828.
Mesquita RV, Axmann D, Geis-Gerstorfer J. Dynamic visco-elastic properties of dental composite resins. Dent Mater 2006; 22:258–267.
Chung S, Yap A, Chandra S, Lim C. Flexural strength of dental composite restoratives: comparison of biaxial and three-point bending test. J Biomed Mater Res B Appl Biomater 2004; 71:278–283.
Vasudeva G. Monomer systems for dental composites and their future: a review. J Calif Dent Assoc 2009; 37:389–398.
Hahnel S, Dowling AH, El-Safty S, Fleming GJ. The influence of monomeric resin and filler characteristics on the performance of experimental resin-based composites (RBCs) derived from a commercial formulation. Dent Mater 2012; 28:416–423.
Irie M, Tjandrawinata R, Lihua E, Yamashiro T, Suzuki K. Flexural performance of flowable versus conventional light-cured composite resins in a long-term in vitro
study. Dent Mater J 2008; 27:300–309.
Arenas G, Cho SD, Bulpakdi P, Matis BA. Effect of bleaching on fracture toughness of resin composites. Oper Dent 2009; 34:703–708.
Elbishari H, Silikas N, Satterthwaite J. Filler size of resin composites, percentage of voids and fracture toughness: is there a correlation? Dent Mater J 2012; 31:523.
Yarin A, Koombhongse S, Reneker H. Taylor cone and jetting from liquid droplets in electrospinning of nanofibers. J App Phys 2001;9:4836–4846.
Ji W, Sun Y, Yang F, van den Beucken J, Fan M, Chen Z, Jansen A. Bioactive electrospun scaffolds delivering growth factors and genes for tissue engineering applications. Pharm Res 2011; 28:1259–1272.
Shao C, Guan H, Liu Y, Gong J, Yu N, Yang X. A novel method for making ZrO2
nanofibers via an electrospinning technique. J Cryst Grow 2004; 267:380–384.
Panda PK, Ramakrishna S. Electrospinning of alumina nanofibers using different precursors. J Mater Sci 2007; 42:2189–2193.
Marlene C, Howard F, McMurdie H, Evans B, Boris P, Harry S, Nicolas C. International Centre for Diffraction Data. National Bureau of Standards Monograph Data for 58 Substance. Joint Committee for Powder Diffraction Standards (JCPDS) Card No 89-2843, for zirconium oxide nanocrystals and films 1981; 25:1–110.
Zanchi C, Ogliari F, Silva R, Lund R, Machado H, Prati C, Carreño N, Piva E. Effect of the silane concentration on the selected properties of an experimental microfilled composite resin. App Adhes Sci 2015; 3:27.
Asmussen E, Peutzfeldt A. Influence of UEDMA, BisGMA and TEGDMA on selected mechanical properties of experimental resin composites. Dent Mater 1998; 14:51–56.
Ilie N, Fleming GJ. In vitro
comparison of polymerisation kinetics and the micro-mechanical properties of low and high viscosity giomers and rbc materials. J Dent 2015; 43:814–822.
Ilie N. Comparative effect of self- or dual-curing on polymerization kinetics and mechanical properties in a novel, dental-resin-based composite with alkaline filler. Materials (Basel) 2018; 11:108–120.
Do T, Church B, Veríssimo C, Hackmyer SP, Tantbirojn D, Simon JF, et al
. Cuspal flexure, depth-of-cure, and bond integrity of bulk-fill composites. Pediatr Dent 2014; 36:468–473.
Moorthy A, Hogg CH, Dowling AH, Grufferty BF, Benetti AR, Fleming GJ. Cuspal deflection and microleakage in premolar teeth restored with bulk-fill flowable resin-based composite base materials. J Dent 2012; 40:500–505.
International Organization for Standardization. Dentistry-polymer-based filling, restorative and luting materials 2009; ISO 4049:2009 (1).
|39.|International Organization for Standardization
. Switzerland: Dentistry – Ceramic Materials; 2007.
Kumar N, Shortall A. Standardization of mixing methods for experimental resin-based composites. J Pak Dent Ass 2011; 20:131–143.
Chiari MDS, Rodrigues MC, Xavier TA, de Souza EMN, Arana-Chavez VE, Braga RR. Mechanical properties and ion release from bioactive restorative composites containing glass fillers and calcium phosphate nano-structured particles. Dent Mater 2015; 31:726–733.
Saghiri MA, Asgar K, Lotfi M, Saghiri AM, Neelakantan P, et al
. Back-scattered and secondary electron images of scanning electron microscopy in dentistry: a new method for surface analysis. Acta Odont Scand 2012; 18:1–7.
Paradella T, Bottino M. Scanning electron microscopy in modern dentistry research. Braz Dent Sci 2012; 15:43–48.
Vogt FG, Williams GR. Advanced approaches to effective solid-state analysis: X-ray diffraction, vibrational spectroscopy, and solid-state NMR. Amer Pharmaceut Rev 2010; 13:58–65.
Stansbury JW, Dickens SH. Determination of double bond conversion in dental resins by near infrared spectroscopy. Dent Mater 2001; 17:71–79.
Scarano A, Carlo FD, Quaranta M, Piattelli A. Bone response to zirconia ceramic implants: an experimental study in rabbits. J Oral Implantol 2003; 29:8–12.
Fadda G, Colombo L, Zanzotto G. First-principles study of the structural and elastic properties of zirconia. Amer Phys Soc 2009; 79:1–13.
Price RL Haberstroh KM, Webster TJ. Enhanced functions of osteoblasts on nanostructured surfaces of carbon and alumina. Med Biol Eng 2003; 41:372–375.
Azad AM. Fabrication of transparent alumina (Al2
) nanofibers by electrospinning. Mater Sci Eng Appl 2006; 435:468–473.
Chia-Ling P, Boyce MC, Rutledge GC. Morphology of porous and wrinkled fibers of polystyrene electrospun from dimethyl formamide. Macromolecules 2009; 42:2102–2114.
Zussman E, Chen X, Ding W, Calabri L, Dikin A, Quintana J, Ruoff S. Mechanical and structural characterization of electrospun PAN-derived carbon nanofibers. Carbon N
Y 2005; 43:2175–2185.
Mucci V, Arenas G, Duchowicz R, Cook WD, Vallo C. Influence of thermal expansion on shrinkage during photopolymerization of dental resins based on bis-GMA/TEGDMA. Dent Mater 2009; 25:103–114.
Lee IB, Min SH, Seo DG. A new method to measure the polymerization shrinkage kinetics of composites using a particle tracking method with computer vision. Dent Mater 2012; 28:212–218.
Cook W, Forrest M, Goodwin A. A simple method for the measurement of polymerization shrinkage in dental composites. Dent Mater 1999; 15:447–449.
Garoushi S, Säilynoja E, Vallittu PK, Lassila L. Physical properties and depth of cure of a new short fiber reinforced composite. Dent Mater 2013; 29:835–841.
Garoushi S, Hatem M, Lassila LVJ, Vallittu PK. The effect of short fiber composite base on microleakage and load-bearing capacity of posterior restorations. Act Biomat Odont Scand 2015; 1:6–12.
Garoushi SK, Lassila LV, Tezvergil A, Vallittu PK. Fiber-reinforced composite substructure: Load-bearing capacity of an onlay restoration and flexural properties of the material. J Contemp Dent Pract 2006; 7:1–8.
Garoushi S, Vallittu PK, Lassila LV. Fracture resistance of short, randomly oriented, glass fiber-reinforced composite premolar crowns. Acta Biomater 2007; 3:779–784.
Cehreli M, Kokat A, Akca K. CAD/CAM Zirconia vs. slip-cast glass-infiltrated alumina/zirconia all-ceramic crowns: 2-year results of a randomized controlled clinical trial. J Appl Oral Sci 2009; 17:49–55.
Sailer I, Pjetursson BE, Zwahlen M, Hammerle C. A systematic review of the survival and complication rates of all-ceramic and metal-ceramic reconstructions after an observation period of at least 3 years. Part II: Fixed dental prostheses. Clin Oral Impl Res 2007; 18:86–96.
Garoushi S, Lassila L, Tezvergil A, Vallittu PK. Load bearing capacity of fibers-reinforced and particulate filler composite resin combination. J Dent 2006; 34:179–184.
Nagata K, Wakabayashi N, Takahashi H, Vallittu PK, Lassila L, Lippo V. Fracture resistance of CAD/CAM-fabricated fiber-reinforced composite denture retainers. Int J Prosthod 2013; 26:381–383.
Chen L, Yu Q, Wang Y, Li H. BisGMA/TEGDMA dental composite containing high aspect-ratio hydroxyapatite nanofibers. Dent Mater J 2011; 7:1187–1195.
Dodiuk-Kenig H, Lizenboim K, Roth S. Performance enhancement of dental composites using electrospun nanofibers. J Nanomater 2008; 8:6–12.
Chen Q, Zhang L, Yoon M, Wu X, Arefin R, Fong H. Preparation and evaluation of nano-epoxy composite resins containing electrospun glass nanofibers. J Appl Polym Sci 2012; 124:444–451.
Tian M, Gao Y, Liu Y, Liao Y, Hedin N, Fong H. Fabrication and evaluation of Bis-GMA/TEGDMA dental resins/composites containing nano fibrillar silicate. Dent Mater J 2008; 24:235–243.
Guo G, Fan Y, Zhang J, Hagan J, Xu X. Novel dental composites reinforced with zirconia-silica ceramic nanofibers. Dent Mater 2012; 28:360–368.
Dyer SR, Lassila LV, Jokinen M, Vallittu PK. Effect of fiber position and orientation on fracture load of fiber-reinforced composite. Dent Mater 2004; 20:947–955.
Garoushi S, Vallittu PK, Lassila LVJ. Short glass fiber reinforced restorative composite resin with semi-interpenetrating polymer network matrix. Dent Mater 2007; 23:1356–1362.
Yildirim M, Edelhoff D, Hanisch O, Spiekermann H. Ceramic abutments-a new era in achieving optimal esthetics in implant dentistry. Int J Periodontal Rest Dent 2000; 20:81–91.
Wen MY, Mueller HJ, CHai J, Woniak WT. Comparative mechanical property of three all-ceramic core materials. Int J Prosthodont 2000; 12:534–541.
Stuart AR, Filser F, Kocher P, Gauckler LJ. Fatigue of zirconia under cyclic loading in water and its implications for the design of dental bridges. Dent Mater J 2007; 23:106–114.
Sundh A, Sjoren G. Fracture resistance of all-ceramic zirconia bridges with differing phase stabilizers and quality of sintering. Dent Mater 2006; 22:778–784.
[Figure 1], [Figure 2], [Figure 3], [Figure 4], [Figure 5], [Figure 6], [Figure 7]
[Table 1], [Table 2]