Effect of chairside surface treatments on biaxial flexural strength and subsurface damage in monolithic zirconia for dental applications

Transcription

1 University of Iowa Iowa Research Online Theses and Dissertations Spring 2016 Effect of chairside surface treatments on biaxial flexural strength and subsurface damage in monolithic zirconia for dental applications Kan Wongkamhaeng University of Iowa Copyright 2016 Kan Wongkamhaeng This thesis is available at Iowa Research Online: Recommended Citation Wongkamhaeng, Kan. "Effect of chairside surface treatments on biaxial flexural strength and subsurface damage in monolithic zirconia for dental applications." MS (Master of Science) thesis, University of Iowa, Follow this and additional works at: Part of the Oral Biology and Oral Pathology Commons

2 EFFECT OF CHAIRSIDE SURFACE TREATMENTS ON BIAXIAL FLEXURAL STRENGTH AND SUBSURFACE DAMAGE IN MONOLITHIC ZIRCONIA FOR DENTAL APPLICATIONS by Kan Wongkamhaeng A thesis submitted in partial fulfillment of the requirements for the Master of Science degree in Oral Science in the Graduate College of The University of Iowa May 2016 Thesis Supervisor: Professor Isabelle L. Denry

3 Graduate College The University of Iowa Iowa City, Iowa CERTIFICATE OF APPROVAL MASTER S THESIS This is to certify that the Master s thesis of Kan Wongkamhaeng has been approved by the Examining Committee for the thesis requirement for the Master of Science degree in Oral Science at the May 2016 graduation. Thesis Committee: Isabelle L. Denry, Thesis Supervisor Deborah V. Dawson Julie A. Holloway

4 ACKNOWLEDGEMENTS First, I would like to thank my thesis advisor, Dr. Isabelle Denry. Thank you for your support and encouragement through out the years as well as the knowledge and techniques that you have shared with me. Without your unwavering patience, this thesis would not have been possible. Thank you Dr. Julie Holloway and Dr. Deborah Dawson, for always sharing ideas and guiding me. Thank you for your motivation and inspiration. Thank you Dr. Ghadeer Thalji, who have always supported me throughout the residency at The University of Iowa. Thank you all my co-residents, Dr. Ahmed Mahrous, Dr. Wissanee Jia-Mahasap, Dr.Jose Miguel Garcia Loera, Dr. Salahaldeen Abuhammoud, Dr. Alessandro Milani, Dr. Maged Abdelaal and Dr. Jorge Garaicoa Pazmino, who have always been there for me and understood me when I was down. You know who I am. I would like to acknowledge Thai Royal scholarship that allowed me to study the Master degree in Oral Science at The University of Iowa. In addition, I would like to thank the American College of Prosthodontics Education Foundation for their financial support for this project. Lastly and the most importantly, I would like to thank all my family who have always supported me and been there for me. Your unconditional love and care are my power that strengthens me to move forward. ii

5 ABSTRACT Objective: The goal of the present study was to investigate the effect of chairside surface treatments on biaxial flexural strength and subsurface damage of monolithic zirconia ceramics. Methods: Specimens (15x15x1.2 mm 3 ) were prepared by sectioning from commercially available zirconia blanks (BruxZir TM ) and sintering according to manufacturer s recommendations. Fully dense specimens were randomly divided into five groups (n=30) and treated as follows; 1) as-sintered (AS) 2) air abraded with 50 µm alumina fine particles (AAF), 3) air abraded with 250 µm coarse alumina particles (AAC), 4) ground (G), and 5) ground and polished (GP) to mimic chairside and dental laboratory treatments. Microstructural changes were thoroughly characterized by optical and scanning electron microscopy, surface profilometry and atomic force microscopy. Crystalline phases and their depth profile were investigated by x-ray diffraction (XRD) and grazing incidence x-ray diffraction (GIXRD). Results were analyzed by Kruskal- Wallis test and Tukey s adjustment for multiple comparisons. A 0.05 level of significance was used. Reliability was evaluated by Weibull analysis. Results: All treatment groups exhibited a significant difference in mean surface roughness (R q ) compared to the as-sintered group (p<0.05). The AAC group showed the highest surface roughness at 1.08 ± 0.17 µm, followed by the G, AAF and AS groups. The GP group exhibited the lowest surface roughness. The group air abraded with fine particles showed the highest mean biaxial flexural strength ( ± MPa), but was not different from the ground and polished group ( ± MPa). The groups air abraded with coarse particles or ground with diamond bur exhibited comparable mean biaxial flexural strength at ± MPa and ± MPa, respectively. The as-sintered group had the lowest mean biaxial flexural strength at ± MPa. The depth of compressive stress layer, measured by GIXRD was approximately 50 µm in the AAF group, followed by the AAC group with ~35 µm, ~10 µm for the ground group and ~5 µm for the ground and polished group. Deep subsurface cracks were observed in the AAC group (~80 µm in depth) and G group (~25 µm in depth), whereas shallower flaws were present in the AAF and GP groups at 10 and 3 µm, iii

6 respectively. Weibull analysis represented a greater reliability in zirconia specimens treated with air abrasion groups. Conclusions: Surface treatments induced the t-m transformation in 3Y-TZP and associated development of compressive stresses to a depth that varied with the severity of the treatment performed. GIXRD revealed that AAF led to the thickest compressive stress layer, followed by AAC, G and GP. SEM revealed that subsurface damage was most severe with AAC, followed by G, AAF and GP. We propose that the flexural strength results can be explained by the difference between the depth of the compressive stress layer induced by the transformation and the depth of the subsurface flaws. iv

7 PUBLIC ABSTRACT This study investigated the effect of chairside surface treatments on bending strength and subsurface damage of a commercially available zirconia ceramic used for all-ceramic dental restorations. Zirconia samples were grouped according to treatment modality as follows; 1) air abraded with fine particles, 2) air abraded with coarse particles, 3) ground and 4) ground plus polished. Untreated, as-sintered zirconia was used as control. Subsurface damage, toughening layer resulting from structural changes, and bending strength were investigated. Our results showed that groups air abraded with fine particles or ground plus polished had the highest bending strength. The group air abraded with fine particles led to the thickest toughening layer followed by groups air abraded with coarse particles, ground, and ground plus polished, respectively. The subsurface damage was most severe with the group air abraded with coarse particles followed by groups ground, air abraded with fine particles and ground plus polished, respectively. Overall this research indicates that the bending strength of zirconia can be explained based on the difference between the depth of the toughening layer and the deepest flaws from chairside surface treatments. Air abrasion with fine particles and grinding followed by fine polishing are acceptable chairside or laboratory treatments for the commercially available zirconia dental ceramic investigated in this study. v

8 TABLE OF CONTENTS LIST OF TABLES... vii LIST OF FIGURES... viii CHAPTER 1. INTRODUCTION... 1 CHAPTER 2. HYPOTHESIS AND SPECIFIC AIMS... 5 CHAPTER 3. MATERIALS AND METHODS Specimens preparation and sintered density measurements Microstructure and surface roughness parameter characterization Subsurface damage and defect characterization Biaxial flexural strength Statistical methods CHAPTER 4. RESULTS Density Microstructure and surface roughness parameter characterization Subsurface damage and defect characterization Crystalline phase characterization Biaxial flexural strength CHAPTER 5. DISCUSSION CHAPTER 6. CONCLUSIONS APPENDIX BIBLIOGRAPHY vi

9 LIST OF TABLES Table 1. Surface treatment procedures of experimental group...6 Table 2. Mean surface roughness (Rq) of the various treatment groups...14 Table 3. Volume fraction of monoclinic phase for specimens treated with various surface treatments obtained form either standard incidence or grazing incidence XRD (GIXRD)...25 Table 4. GIXRD data after air abrasion with fine particles...26 Table 5. GIXRD data after air abrasion with coarse particles...26 Table 6. GIXRD data for ground specimens...27 Table 7. GIXRD data for ground and polished specimens...27 Table 8. Mean flexural strength and Weibull analysis...32 Table A1. Flexural strength of as-sintered zirconia specimens...41 Table A2. Flexural strength of air abraded with fine particles zirconia specimens...41 Table A3. Flexural strength of air abraded with coarse particle zirconia specimens...42 Table A4. Flexural strength of ground zirconia specimens...43 Table A5. Flexural strength of ground and polished zirconia specimens...44 vii

10 LIST OF FIGURES Figure 1. Bonded interface configuration used to study the subsurface damage....8 Figure 2. Schematic of indent geometry used to calculate the depth of layer removed by polishing Figure 3. A). AFM micrograph of as-sintered 3Y-TZP (1580 o C ). B). SEM micrograph of as-sintered 3Y-TZP...13 Figure 4. Representative surface roughness profiles for the various treatment groups. A) As-sintered; B) Air abraded with fine particles; C) Air abraded with coarse particles; D) Ground; and E) Ground and polished Figure 5. Optical micrographs of subsurface damage induced by various surface treatments. A and B): Air abraded with fine particles; C and D): Air abraded with coarse particles; E and F): Ground; and G and H): Ground and polished.16 Figure 6. SEM micrographs of as-sintered zirconia specimens. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification Figure 7. SEM micrographs of specimens air abraded with 50 µm alumina particles. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification...19 Figure 8. SEM micrographs of specimens air abraded with 250 µm alumina particles. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification...20 Figure 9. SEM micrographs of zirconia specimens ground with diamond bur under water cooling. A) Cross section; B) Line angle between top surface and cross section; C) Top surface with 10,000X magnification; D) Top surface at 20,000X magnification...21 viii

11 LIST OF FIGURES Figure 10. SEM micrographs of zirconia specimens ground with diamond bur under water cooling and polished according to manufacturer s recommendations. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000 magnification...22 Figure 11. X-ray diffraction patterns of the various surface treatment groups A, F, and K: As-sintered specimen; B, G and L: Specimen air abraded with fine particles; C, H and M: Specimen air abraded with coarse particles; D, I and N: Ground specimen; E, J and O: Specimen ground and polished; A through E: standard incidence XRD; F though J: Grazing incidence XRD incidence angle: 1º; K through O: Grazing incidence XRD incidence angle: 9º...24 Figure 12. Monoclinic volume fraction as a function of depth for the various groups (GIXRD)...28 Figure 13. Intensity ratio for the t (002) and t (200) reflections (indicating ferroelastic domain switching) for the various groups...29 Figure 14. Peak position for the main tetragonal t (101 ) reflection as a function of depth..30 Figure 15. Full-Width at Half-Maximum for the main tetragonal t (101) reflection as a function of depth...31 Figure 16. Plots of Ln(Ln(1/(1-Pf)) as a function of Ln(strength) used to determine Weibull modulus, A) As-sintered; B) Air abraded with fine particles; C) Air abraded with coarse particles; D) Ground with diamond bur; and E) Ground and polished...33 Figure 17. Effect of surface treatments on subsurface damage in zirconia specimens, A) Air abraded with fine particles; B) Air abraded with coarse particles; C) Ground with diamond bur; and D) Ground and polish...38 ix

12 CHAPTER 1. INTRODUCTION Zirconia, one of the most promising candidate materials for all-ceramic dental restorations, exists in three crystallographic structures, monoclinic (m), tetragonal (t) and cubic (c). The monoclinic form is stable from room temperature to 1170 o C, then transforms to a tetragonal from 1170 o C to 2370 o C. The structure becomes cubic above 2370 o C and up to melting point (1). Upon cooling, the transformation from tetragonal to monoclinic (t-m) is accompanied by an increase in volume (~4.4%) and subsequently leading to shattering upon cooling (2, 3). In order to prevent this transformation upon cooling, pure zirconia is alloyed with stabilizing oxides, such as CaO, MgO, Y 2 O 3 or CeO 2, to retain the tetragonal form in a metastable state at room temperature (3-5). Yttrium oxide is the most widely used stabilizer in dental zirconia, at a concentration of 3 mol.% (3Y-TZP). One of the most prominent characteristics of zirconia ceramics is their martensitic stress-induced transformation, where the crystalline structure in solid state is changed through athermal, diffusionless motion of atomic boundaries, which proceeds at the speed of sound (4, 6, 7). The overall mechanism of martensitic transformation is the conversion of tetragonal phase to monoclinic phase caused by shearing displacement of zirconia ions, and the displacement of oxygen ions from tetragonal to monoclinic lattice along the c axis (8). The martensitic t-m transformation is induced either by temperature or external stress under isothermal conditions (4, 9). Both transformations are important, in which temperature-driven transformation will determine the amount of tetragonal phase remaining after thermal cycling, whereas stress-induced transformation will ensure the excellent fracture toughness of zirconia ceramics (10). The stress-induced phase transformation of zirconia ceramics is the process of phase transformation from metastable tetragonal to stable monoclinic in the tensile field around the tip of an advancing crack (12, 17). This transformation is associated with an increase in volume (9), leading to a net compressive stress field at the crack-tip. This stress-induced volumetric change reduces the local crack tip intensity, and increases the resistance of crack propagation with a toughening effect (14). 1

13 Due to their excellent mechanical properties, zirconia ceramics are not only used in industrial applications, but also in biomedical and dental applications (1). A major biomedical application is total hip or knee replacements, with excellent biocompatibility (11, 12). The clinical success of zirconia femoral heads led to the implantation of more than 600,000 prostheses worldwide (13). Biomedical grade 3Y-TZP exhibits optimal mechanical strength due to the stress-induced phase transformation toughening, as mentioned earlier. However, metastable 3Y-TZP is also susceptible to low temperature degradation (LTD) or aging in the presence of water (14, 15) or in vivo. An autocatalytic and progressive t-m transformation at the surface of the prosthesis leads to grain pull out, surface roughening and micro-cracking. The aging process is also associated with wear and may lead to crack propagation and catastrophic failure, as reported for a series of femoral heads in 2001 (13, 14). Metal-free dental restorations are becoming increasingly popular due to their superior esthetics and concerns with potential allergies to metals (16). All-ceramic dental restorations are custom-fabricated and are therefore likely to present some processing defects. For example, heat pressing induces the creation of porosity within leucite-based or lithium-disilicate-based restorations. Hard machining introduces residual stresses in leucite-based and lithium-disilicate based ceramics (17). This explains why hard machined all-ceramic restorations are limited to single unit or short span multi-unit restorations. Zirconia ceramics were introduced to dentistry in 2001 due to their excellent mechanical properties and acceptable esthetics, allowing their use for the fabrication of posterior fixed partial prostheses. Zirconia ceramics are also used as a core materials with reduced thickness for veneering, or full contour monolithic restorations (18). To date, zirconia dental ceramics have had an excellent track record in terms of clinical performance. Cumulative 5-year survival rates of more than 95% have been reported, with fractures of the veneering porcelain being the most common complication (19-21). The mechanism for chipping of the veneering porcelain has been linked to the difference between the coefficient thermal contraction between zirconia and veneering materials, and tempering stresses created during rapid cooling (17). Although these issues have been successfully addressed, monolithic zirconia dental restorations are becoming a popular alternative to bilayered zirconia-based dental restorations. Zirconia dental 2

14 restorations are fabricated through the soft machining process, using computer-aided design/computer-aided manufacture (CAD/CAM). The partially sintered 3 mol.% yttriastabilized zirconia (3Y-TZP) blocks are milled and sintered at high temperature to achieve full density as recommended by the manufacturers, e.g. BruxZir TM HT 2.0 is sintered at 1580 o C for 2.5 hours (22). The sintering temperature determines zirconia grain size, phase assemblage, and mechanical properties. Higher sintering temperatures lead to an increase in grain size, but a decrease in tetragonal content (23). Meanwhile, upon the sintering process, the pre-sintered zirconia blocks are subject to dimensional shrinkage ~ 20-25% (24). Although the dimensional shrinkage is compensated by the computer-aided design, the restorations often require internal adjustments. In addition, surface adjustments of the restorations are often necessary to achieve optimal occlusion, proper seating and improve bonding by creating micromechanical retentions between the restoration and luting cement. Grinding followed by polishing with polishing kits are recommended for occlusal adjustments (25), whereas air abrasion is performed to improve the mechanical retention between zirconia restorations and luting cements (26-29). These surface adjustments are likely to trigger the stress-induced phase transformation, which is accompanied by the formation of a compressive stress layer, leading to an increase in mechanical properties (30-33). However, these surface adjustments also induce surface and/or subsurface damage (34). The efficacy of the phase transformation is dependent on the mean zirconia grain size, where larger grains are more likely to transform to the monoclinic phase (14, 17, 35). The mechanical strength of zirconia is also directly dependent on the thickness of compressive stress zone (33, 36). X-ray diffraction (XRD) analysis also showed that residual stresses created by surface treatments are associated with alterations of the XRD-profile consisting in 1) shift in the peak position of the main tetragonal t (101) reflection (37), 2) an increase in the amount of monoclinic volume fraction (33), 3) an increase in the amount of ferroelastic domain switching (38) and 4) an increase in the Full-Width at Half-Maximum of the main tetragonal t (101) reflection (39). The effect of various surface treatments on the strength of zirconia ceramics has been studied extensively (30, 32, 40, 41). It was shown that air abrasion with fine particles was more efficient in promoting the formation of a thick compressive stress 3

15 layer, and was associated with an increase in mean biaxial flexural strength. Meanwhile, surface modifications created by air abrasion with coarse particles or grinding with diamond coated instruments led to mixed results, either increasing (32, 41) or decreasing the mean biaxial flexural strength (40, 42). Polishing after grinding was shown to decrease the thickness of the compressive stress layer as well as the size of the flaws created by grinding, and consequently associated with a decrease in strength (43). Understanding the effect of surface modifications, e.g. grinding, polishing and airabrasion on the thickness of the compressive stress layer and subsurface damage is therefore critical in the achievement of successful restorations. These modifications are likely to affect the long-term performance of zirconia dental restorations, but to date; assessment of subsurface damage has been limited. Thus, the objective of this study was to evaluate the effect of chairside surface treatments on strength and subsurface damage of monolithic zirconia. 4

16 CHAPTER 2. HYPOTHESIS AND SPECIFIC AIMS There is a significant amount of information available on the effect of surface treatments on the phase transformation and mechanical properties of zirconia-based dental ceramics (32, 44, 45). However, little is known about the actual impact of these surface treatments on subsurface damage, which strongly influences both short and longterm clinical performance (32, 34). These surface treatments may increase the mechanical strength, but they may also decrease reliability, depending on the flaw population. To date, no consensus has been reached regarding which surface treatment is the most effective. Thus, the purpose of this study was to investigate the effect of clinically relevant (chairside or laboratory-performed) surface treatments on the biaxial flexural strength and subsurface damage of a commercially available monolithic dental zirconia (BruxZir TM ) using a bonded interface configuration, in combination with state of the art characterization techniques. Microstructural and crystallographic subsurface changes were thoroughly characterized after grinding, fine or coarse air abrasion, and grinding followed by polishing. Grazing incidence x-ray diffraction (GIXRD) was used to evaluate the transformation and depth of the compressive stress layer. Subsurface damage was investigated by both optical and scanning electron microscopy (SEM) in a bonded interface configuration. The null hypothesis was that the depth of the compressive stress layer is independent of the type of surface treatment. 5

17 CHAPTER 3. MATERIALS AND METHODS 3.1 Specimens preparation and sintered density measurements Commercially available pre-sintered monolithic zirconia blanks (98.5 mm in diameter; 12 mm in height; (BruxZir TM, Glidewell laboratories, Newport Beach, CA, USA) were sectioned into squares (15 x 15 x 1.2 mm 3 ) with a high speed diamond saw (Buehler Isomet, Evanston, IL). Specimens (n=30 per group) were sintered at 1580 o C for 2.5 hours, and furnace-cooled according to the manufacturer s recommendations. The specimens were randomly assigned to various treatment groups (Table 1). Sintered specimens (n=3) were individually weighted on an analytical scale (Mettler, Toledo). Their density was measured with a Helium Pycnometer (Accupyc 1340, Micrometrics Instrument Corporation, Norcross, GA, USA) according to ASTM standard (46). Group AS AAF AAC G GP Surface treatment As-sintered (Control group) Air abraded with 50 µm alumina particles (fine) at a pressure of 4 bars Air abraded with 250 µm alumina particles (coarse) at a pressure of 4 bars (Negative control group) Ground with fine grit diamond bur (856DEF , Brasseler, USA) with water spray cooling Ground with fine grit diamond bur (856DEF , Brasseler, USA) with water spray cooling, and polished with recommended polishing kit (Dialite ZR, Brasseler, USA) without water spray cooling Table 1. Surface treatment procedures of experimental groups. 6

18 3.2 Microstructure and surface roughness parameter characterization The mean real grain size was determined from atomic force micrographs (AFM) acquired in contact mode (Veeco di Innova, Bruker AXS Inc., Madison, USA). Specimens were polished to a 0.5 µm finish using a series of abrasives ending with diamond polishing paste. The specimens were thermally etched at 1250 o C for 30 minutes and furnace-cooled. They were ultrasonically cleaned in ethanol prior to imaging by AFM. The mean grain size was determined by the linear intercept method as described in ASTM standard E The real grain size (D) was calculated using equation (1), (47) Where C is the length of a line drawn across the micrograph, N is the number of grain boundaries counted along line C. D = 1.56 C N (1) The surface roughness was analyzed by surface profilometry (Surftest SJ-210; Mitutoyo Corporation, Aurora, IL). The root mean square roughness (Rq) was measured on surface treated specimens (n=5 per group). The investigation was performed at five locations on each specimen with a transverse speed 0.5 mm/s and cutoff value of 0.8 as described in ASTM D (48). The microstructure was also characterized by scanning electron microscopy (SEM; Hitachi, S-4800) under secondary electron imaging and in ultra high-resolution mode. Specimens were subjected to various surface treatments as described previously, ultrasonically cleaned in ethanol, and gold-coated prior to SEM examination. 7

19 3.3 Subsurface damage and defect characterization Subsurface damage was assessed, using a bonded interface configuration (49). Bar-shaped specimens (10 x 4 x 6 mm 3 ) were polished to a 1-micrometer finish. Two polished surfaces were bonded together, using a cyanoacrylate-based adhesive (n=3 per group). The resultant top surface was treated according to the assigned groups (Figure 1). After separation of the polished surfaces by immersion in acetone, subsurface damage was characterized by both optical microscopy and scanning electron microscopy. Treated surface Polished surface 6 mm Y-TZP 4 mm 10 mm Bar-shaped specimen Polished surface Bonded-interface configuration Figure 1. Bonded interface configurations used to study subsurface damage 3.4 Crystalline phase characterization Surface crystalline phases were analyzed by both grazing incidence x-ray diffraction (GIXRD) and x-ray diffraction (XRD) on bulk specimens (n=3 per group). Scans were performed in the two-theta range 27 to 37 degrees at the scanning rate of 0.5 degree per minute (Rigaku diffractometer, λ Cu Kα = Å). Relative amounts of monoclinic and tetragonal phases were determined using the expression proposed by Garvie and Nicholson(3) in equation (2). X! = I!!111! + I! (111) I!!111! + I! (111) + I! (101) (2) respectively. Where I m and I t are the integrated intensities of monoclinic and tetragonal phases, 8

20 The monoclinic volume fraction, V m was calculated using the equation proposed by Toraya (50) in equation (3). V! = 1.311X m X m (3) Subsurface crystalline phases and stress state were measured by grazing incidence X-ray diffraction (GIXRD) on bulk specimens. A Vickers indentation was made at the center of the specimens under a 98 N load. GIXRD was performed in the two-theta range 27 to 37 degrees at the scanning rate of 0.5 degree per minute (λ Cu Kα = Å). GIXRD was performed at various incidence angles from 1 to 9 degrees. The x-ray penetration depth (D) was calculated according to equation (4), (51, 52) D = 2α μ (4) Where µ is the x-ray absorption coefficient of 3Y-TZP ( µm -1 ), and α is the angle of incidence. Vickers indentation A 135 o θ D Y-TZP A A! θ = 67.5 o tan θ = 2 D Figure 2. Schematic of indent geometry used to calculate the depth of the layer removed by polishing. Where A is the diagonal of the Vickers indentation, 2θ is the Vickers indenter angle; D is the depth of the indentation. 9

21 After GIXRD scanning, specimens were polished with 3-µm diamond paste for 30 seconds, and then, the diagonal of the Vickers indentation was re-measured using optical microscopy. The relative depth of the removed layer was calculated from the geometry of the resulting Vickers indentation as described in Figure 2. The process was repeated until the intensity of both the tetragonal t (002) and t (200) reflections returned to a base value of 0.66, indicating the absence of stresses. The depth of the compressive stress layer was assessed by 1), the relative monoclinic volume fraction, with a control volume fraction of approximately 0.05; 2), the presence or absence of ferroelastic domain switching (FDS), with a base value for the FDS ratio of about 0.66 for annealed zirconia; 3), the peak position shift for the main tetragonal reflection t (101), where the peak position in the absence of strain is at a d spacing of Å for 3Y-TZP; and 4), the Full-Width at Half-Maximum of the main tetragonal reflection, with a base FWHM value of about 0.33 in the absence of strain. 3.5 Biaxial flexural strength The mean biaxial flexural strength was measured using a ball-on-ring-of-balls fixture following the method developed by Wachtman et al. (53). Specimens (n=30 per group) were loaded at the center of the support circle, with the treated side in tension. Testing was performed at a crosshead speed of 0.5 mm/min with use of a universal testing machine (Instron 5965, Instron Corporation, Canton, MA). The mean flexural strength was calculated from the failure loads, according to ISO standard 6872 (54) in equation (5). Specimens were inspected after testing and rejected if the fracture path did not cross the specimen center. σ! = 3P(1 + ν) 4πt!!1 + 2 ln(a b) + (1 ν) (1 + ν) b! a!!1 2a!! R!! (5) Where P is the fracture load (N), t is the thickness of the specimen (mm), a is the radius of support circle (mm), b is the radius of uniform loading at center (mm), R is the disc radius (mm), and Poisson s ratio (ν) is equal to

22 3.6 Statistical methods Conformance to model assumptions was assessed using standard residual analyses including graphical displays, the Shapiro-Wilk test for assessment of normality, and the Brown-Forsythe test of variance homogeneity. In the assessment of surface roughness via the Rq measurement, five specimens were randomly allocated to each of five experimental groups based upon surface treatment, and five replicate surface roughness measurements were made within each specimen. Analysis was carried out using a mixed linear modeling approach, which allowed for the correlation of the repeated measures within specimen, with specimen being treated as a random effect. In consideration of biaxial flexural strength, evaluation of assumptions of the oneway analysis of variance indicated violation of the assumptions of normality (p=0.0067, Shapiro-Wilk test) and variance homogeneity (p=0.041, Brown and Forsythe test), leading to the use of the nonparametric analog of one-way ANOVA, the Kruskal-Wallis test. Estimates of reliability (Weibull modulus: m) were analyzed using linear regression according to the method described by Conover (55), which specifies that the estimate of the Weibull modulus is the slope derived from the linear regression of Ln(Ln(1/(1-(rank- 0.5/N))) on the natural logarithm of flexural strength. Rank refers to the rank order of the corresponding flexural strength value. In order to test whether the slopes differed among the five surface treatment groups, the fit of a joint linear model specifying a separate slope and intercept for each of the five treatment groups was compared to that of a model with different intercepts but only a single slope, i.e., specifying that the Weibull modulus was the same across all surface treatment groups; in addition, all pairwise comparisons of slopes among experiment groups were made. In all instances, adjustment was made for all pairwise comparisons among treatment groups, in conjunction with an overall (experiment-wise) 0.05 Type I error level. The Tukey method, which is appropriate for all pairwise comparisons of means, was used in the analysis of surface roughness and flexural strength; in the latter, the modification of the Tukey method in the context of Kruskal-Wallis analysis used was as described by 11

23 Conover (55). For pairwise comparisons of slopes, the Bonferroni method of adjustment was used, again specifying an experiment-wise Type I error level of

24 CHAPTER 4. RESULTS 4.1 Density The density of the sintered specimens was ± g/cm 3, which corresponds to 99.7 % of the theoretical density for 3Y-TZP (6.080 g/cm 3 ). 4.2 Microstructure and surface roughness parameter characterization The mean real grain size of as-sintered specimens (1580 ºC) was 1.21±0.16 µm. Representative AFM and SEM images are displayed in Figure 3A-3B. A B 1 µm Figure 3. A) AFM micrograph of as-sintered zirconia (1580 o C); B) SEM micrograph of as-sintered zirconia. 13

25 Surface roughness measurements are summarized in Table 2 and Figure 4. The data provided strong evidence of differences in surface roughness among the groups (p < ), and following Tukey adjustment for the ten pairwise multiple comparisons of treatment group means, it was found that the mean surface roughness of each group differed significantly from all of the others. The ground and air abraded in both either fine or coarse particles groups showed significant higher surface roughness than assintered group (0.42±0.06 µm). The group air abraded with coarse particles had the highest surface roughness (1.08±0.17 µm), followed by the ground group (0.91±0.14 µm) and air abraded with fine particles group (0.69±0.09 µm), respectively. The group ground and polished had the lowest surface roughness (0.18 ±0.03 µm). Experimental Groups AS: As-sintered 3Y-TZP AAF: Air abraded with fine particles (50 µm) with 4 bars AAC: Air abraded with coarse particles (250 µm) with 4 bars Surface roughness: Rq (SD) 0.42 (0.06) a 0.69 (0.09) b 1.08 (0.17) c G: Ground with diamond bur 0.91 (0.14) d GP: Ground with diamond bur, and polished with polishing kit 0.18 (0.03) e Identical letters denote no statistically significant difference among groups after Tukey adjustment for multiple comparisons in conjunction with an experiment-wise Type I error level of In this instance, the mean roughness of each group was significantly different from all of the others. Table 2. Mean surface roughness (Rq) of the various treatment groups. 14

26 Surface Roughness (µm ) A AS Surface Roughness (µm ) B AAF Surface Roughness (µm ) C AAC Surface Roughness (µm ) D G Surface Roughness (µm ) E GP Figure 4. Representative surface roughness profiles for the various treatment groups. A) As-sintered; B) Air abraded with fine particles; C) Air abraded with coarse particles; D) Ground; and E) Ground and polished. 15

27 A. Treated surface B. Polished bonded interface configurations ~10 µm AAF 30 µm AAF 30 µm C. D. ~80 µm AAC 30 µm AAC 30 µm E. F. ~25 µm G 30 µm G 30 µm G. H. ~3 µm GP 30 µm GP 30 µm Figure 5. Optical micrographs of subsurface damage induced by various surface treatments. A and B): Air abraded with fine particles; C and D): Air abraded with coarse particles; E and F): Ground; and G and H): Ground and polished 16

28 4.3 Subsurface damage and defect characterization Optical micrographs of the subsurface damage induced by various surface treatments are displayed in Figure 5 (A-H). Significant subsurface damage can be clearly seen in both AAC and G groups, as indicated by arrows. In the AAC group, flaws propagated parallel to the surface, extending approximately 80 µm in depth (Figure. 5C- 5D). Ground specimens exhibited relatively straight flaws, which propagated parallel to the top surface (Figure. 5E-5F) and extended to a maximum depth of 25 µm. Only minor subsurface damage was presented in both AAF (Figure.5A-5B), and GP (Figure.5G-5H) groups, with a maximum depth of 10 and 3 µm, respectively. SEM micrograph displayed in Figures 6 through 10 show the topography and microstructural details of surface and subsurface damages for the various groups. The assintered specimens exhibited remnant grooves from diamond blade sectioning on the surface and no apparent subsurface damage (Figure 6A). Tilted SEM image showed a clearly defined line-angle between surface and subsurface aspects (Figure 6B). Zirconia grains can be seen at higher magnification (Figures 6C and 6D). Deeper surface defects, and uniformly rough surface damage with evidence of plastic deformation and grain pullout were observed after air abrasion with fine particles (Figure 7). Higher magnification of surface damage confirmed extensive plastic deformation (Figures 7C and 7D). Air abrasion with coarser particles produced extensive surface damage with deep gouges and substantial plastic deformation on the treated surface (Figures 8C and 8D). This translated into curving subsurface cracks extending parallel to the treated surface at a depth of 80 µm (Figures 8A and 8B). Remaining grinding grooves were present on the ground zirconia surface (Figures 9B and 9C). Shallower damage was observed on the GP specimens (Figures 10A and 10B). Higher magnification examination of the surface revealed plastically deformed and aligned zirconia grains, together with some remaining surface damage (Figures 10C and 10D). 17

29 Treated surface A Polished surface Z Bonded interface configuration X Y B C D Figure 6. SEM micrographs of as-sintered zirconia specimens. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification. 18

30 A Treated surface Polished surface Z Bonded interface configuration X B Y C D Figure 7. SEM micrographs of specimens air abraded with 50 µm alumina particles. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification 19

31 A Treated surface Polished surface Z Bonded interface configuration X B Y C D Figure 8. SEM micrographs of specimens air abraded with 250 µm alumina particles. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification 20

32 A Treated surface Polished surface Z Bonded interface configuration X Y B C D Figure 9. SEM micrographs of zirconia specimens ground with diamond bur under water cooling. A) Cross section; B) Line angle between top surface and cross section; C) Top surface at 10,000X magnification; D) Top surface at 20,000X magnification 21

33 A Treated surface Polished surface Z Bonded interface configuration X Y B C D Figure10. SEM micrographs of zirconia specimens ground with diamond bur under water cooling and polished according to manufacturer s recommendations. A) Cross section; B) Line angle between top surface and cross section; C) Top surface with 10,000X magnification; D) Top surface with 20,000X magnification. 22

34 4.4 Crystalline phase characterization X-ray diffraction patterns for various surface treatment groups are displayed in Figure 11. As-sintered zirconia contained only tetragonal phase (Figure.11A). The XRD patterns for the air abraded and ground were characterized by significant broadening of the main reflection for tetragonal zirconia, and a development of a left shoulder corresponding to the rhombohedral or pseudo-cubic phase, indicating the present of a significant amount of residual stress (Figure. 11B-11D). Peak broadening was associated with various amounts of monoclinic phase (Table 3.) and reversal of intensities for the t (002) and t (200) reflections of the tetragonal phase, indicating the occurrence of ferroelastic domain switching (Figure. 11B-11D). The XRD pattern for zirconia specimens ground and polished revealed the presence of the tetragonal phase only, similar to the as-sintered group (Figure. 11E). The volume fraction of monoclinic phase analyzed by standard incidence XRD and GIXRD at incident angles of 1 o (1 µm deep) and 9º (5 µm deep) is summarized in Table 3 and Figure 11. The AAF and AAC groups presented the highest volume fraction of monoclinic phase at the top surface at 8%, whereas the G, GP and AS showed relatively low volume fraction of monoclinic phase at 2%, 1% and 1%, respectively. GIXRD detected more volume fraction of monoclinic phase at 1 µm, when compared to the top surface; they were decreased corresponding to the depth of the specimens. The alteration of FWHM, peak position and ferroelastic domain switching were observed in GIXRD at incidence angle of both 1 o and 9 o (Figure 11F-11O). The AAF, AAC, G and GP groups presented the alteration of the GIXRD profiles at the depth of 1 µm, which were significantly difference from the XRD profile at the top surfaces. The GIXRD profile was returned to their normal value in the G group at the depth of 5 µm (Figure 11O), and the FDS ratio for the AAF group was slightly reversed at the depth of 5 µm (Figure 11L). 23

35 Figure 11. X-ray diffraction patterns of the various surface treatment groups. A, F, and K: As-sintered specimen; B, G and L: Specimen air abraded with fine particles C, H and M: Specimen air abraded with coarse particles; D, I and N: Ground specimen E, J and O: Specimen ground and polished A through E: standard incidence XRD F though J: Grazing incidence XRD incidence angle: 1º K through O: Grazing incidence XRD incidence angle: 9º 24

36 Measurements of the XRD distinctive features are summarized in Table 3 through 7 for the various treatments, and displayed in Figure 12 through 15. Air abrasion with either fine or coarse alumina particles showed the relatively thick compressive stress layer at 50 and 35 µm, respectively. For the AAF group, the relative monoclinic phase volume fraction and the peak position of tetragonal t (101) reflection returned to their base values at a depth of 37 µm, whereas the FWHM went back to its unstrained value at a depth of 76 µm. The FDS ratio returned to its unstrained value at the depth of 52 µm (Table 4). Similarly, the AAC group presented the reversal of both the volume fraction of monoclinic phase and FDS ratio at a depth of 20 µm. The peak position and FWHM were reversed at the depth of 37 and 42 µm, respectively (Table 5). Zirconia specimens treated by grinding with diamond bur showed the reversal of the GIXRD profile at the depth of 12 µm, except for the volume fraction of monoclinic phase and the peak position, which reversed at a depth of 5 and 2 µm (Table 6). The GP group showed the reversal of volume fraction of monoclinic and peak position at a depth of 3 µm. The FDS ratio and FWHM were reversed at a depth of 12 µm (Table 7). According to the XRD distinctive features, the AAF group exhibited the thickness compressive stress layer (~50 µm), followed by the AAC (~35 µm), G (~10 µm) and GP (~5 µm), respectively. Groups XRD measurements Surface monoclinic phase volume fraction GIXRD measurements Monoclinic phase volume fraction at 1 µm depth GIXRD measurements Monoclinic phase volume fraction at 5 µm depth AS AAF AAC G GP Table 3. Volume fraction of monoclinic phase for specimens treated with various surface treatments obtained form either standard incidence or grazing incidence XRD (GIXRD). 25

37 Depth (µm) Volume fraction of Monoclinic phase Ferroelastic Domain Switching ratio Tetragonal t (101) reflection FWHM Peak position (A ) Asymmetry factor (control) Table 4. GIXRD data after air abrasion with fine particles Depth (µm) Volume fraction of Monoclinic phase Ferroelastic Domain Switching ratio Tetragonal t (101) reflection FWHM Peak position (A ) Asymmetry factor (control) Table 5. GIXRD data after air abrasion with coarse particles 26

38 Depth (µm) Volume fraction of Monoclinic phase Ferroelastic Domain Switching ratio Tetragonal t (101) reflection FWHM Peak position (A ) Asymmetry factor (control) Table 6. GIXRD data for ground specimens. Depth (µm) Volume fraction of Monoclinic phase Ferroelastic Domain Switching ratio Tetragonal t (101) reflection FWHM Peak position (A ) Asymmetry factor (control) Table 7. GIXRD data for ground and polished specimens. 27

39 Volume fraction of monoclinic phase Figure 12. Monoclinic volume fraction as a function of depth for the various groups (GIXRD) Depth (micrometers) 28

40 Figure 13. Intensity ratio for the t (002) and t (200) reflections (indicating ferroelastic domain switching) for the various groups. Depth (micrometers) 29

41 Peak position of tetragonal t (101) reflection Depth (micrometers) Figure 14. Peak position for the main tetragonal t (011) reflection as a function of depth. 30

42 Figure 15. Full-Width at Half-Maximum for the main tetragonal t (011) reflection as a function of depth. Depth (micrometers) 31

43 4.5 Biaxial flexural strength The mean biaxial flexural strength data is summarized in Table 8. The as-sintered control group showed the lowest mean biaxial flexural strength ( ± MPa), significantly lower compared to all other surface treatment groups after adjustment for multiple comparisons. The AAF group had the highest mean biaxial flexural strength at ± MPa, but was not statistically different from the GP group, although AAF and GP had significantly greater flexural strength than all other groups. There was no significant difference between AAC and G groups, which were intermediate in terms of the distribution of flexural strength. Groups n Flexural strength: MPa (SD) Weibull Analysis: m (SD) R 2 AS (141.92) c 9.04 (0.58) b AAF (202.58) a 9.80 (0.32) a,b AAC (147.62) b (0.37) a G (196.77) b 8.33 (0.29) b GP (209.76) a 8.45 (0.26) b Identical letters denote no statistically significant difference among groups after adjustment for multiple comparisons in conjunction with an experiment wise Type I error level of Table 8. Mean flexural strength and Weibull analysis Three specimens were excluded from this study since they did not fracture through the center of the specimens (Table 8). The biaxial flexural strength for all treated groups are presented in Weibull plots in Figure 16. The Weibull modulus (m) for assintered zirconia specimen was 9.04 ± The Weibull modulus slightly decreased in zirconia specimens treated by grinding and/or polishing at 8.33 ± 0.29, and 8.45 ± 0.26, respectively. However, zirconia specimen treated by air abrasion both either fine or coarse particles showed significantly difference Weibull modulus when compared to other groups (9.80 ± 0.32 and ± 0.36, respectively) (Table 8). When pairwise comparisons of Weibull modulus were carried out according to the standard Bonferroni method in conjunction with an overall 0.05 level of Type I error; results showed that the Weibull modulus from the AS, G and GP differed significantly from that of the AAC group. No other pairwise comparisons of slopes could be said to differ significantly after adjustment for multiple comparisons. 32

44 Figure 16. Plots of Ln(Ln(1/(1-Pf)) as a function of Ln(strength) used to determine Weibull modulus, A) As-sintered; B) Air abraded with fine particles; C) Air abraded with coarse particles; D) Ground with diamond bur; E) Ground and polished 33

45 CHAPTER 5. DISCUSSION The purpose of this study was to investigate the effect of chairside surface treatments on biaxial flexural strength and subsurface damage of a commercially available dental zirconia. The results of the present study showed that the density of sintered-zirconia specimens (1580 o C for 2.5 hours) exceeded 99.7 % of the theoretical density for 3Y-TZP. The density is a primary indicator for controlling the quality of sintered ceramics. Inadequate sintering indicates that residual porosity is present. This is likely to affect the mechanical performance of zirconia as pores can affect the hardness and elastic modulus (56). Grain size increases as a function of sintering temperature, and directly influences the stability and mechanical properties of zirconia. Grain size is determined by the temperature and duration of the sintering process (57). Higher sintering temperatures result in larger grain sizes and enhance the transformability of tetragonal phase leading to higher fracture toughness. For commercially available zirconia ceramics for dental restorations, the sintering temperature ranges between 1350 o C to 1580 o C, and the duration from 2 to 6 hours depending on manufacturer s recommendations (23, 57). Our results showed that the mean real grain size of Bruxzir TM zirconia was 1.21 ± 0.16 µm, indicative of a zirconia more susceptible to transform than zirconia sintered at lower temperature. Surface roughness measurements showed that the AAC group had the greatest surface roughness, followed by G, AAF, AS and GP, respectively. The results are similar to other published studies, where polishing after surface modifications leads to a smoother surface of the zirconia restorations (58, 59). Surface roughness also plays a major role on wear behavior against different antagonist materials. Several in vitro studies confirmed a positive correlation between surface roughness of zirconia specimens and wear rate against different antagonists materials in two-body wear test (58-60). However, our results showed that surface roughness might be associated with extensive subsurface damage, as depicted in Figure 5. For example, the AAC group showed the greatest roughness of 1.08 ± 0.17 µm, and the subsurface damage extended to the depth of approximately 80 µm. 34

46 Externally applied stresses, e.g. grinding, air abrasion and fractures, triggered phase transformation from metastable tetragonal (t) to monoclinic phase (m), giving rise to volume expansion and subsequently the formation of compressive stress fields around the crack tip (4, 35, 61). Studies have shown that Vickers indentations in zirconia ceramics induce the formation of a zone of plastic deformation (62), and an elastic zone surrounding the plastic zone. After the indenter is removed, the recovery of the elastic zone is inhibited completely by the plastic zone, resulting in the formation of a residual stress field in associated to plastic/elastic contact. The residual stress field is compressive inside the plastic zone, and tensile in tangential directions in the elastic zone. The magnitude of the tensile stress is dependent on the applied load (33). For this study, the depth of compressive stress layers was analyzed by grazing incidence x-ray diffraction on surface-indented specimens, and showed that the residual stress from surface treatments leads to the alteration of the peak position of tetragonal t (101) reflection, the ferroelastic domain switching, the Full-Width at Half-Maximum of tetragonal t (101) phase, and the volume fraction of monoclinic phase. GIXRD profile revealed that the peak position of tetragonal t (101) reflection in the treated groups shifted toward a lower diffraction angle, indicating a shifting of the d spacing to the higher value. The d spacing, or space between two lattice (37) in the absence of strain is Å for 3Y-TZP, whereas the d spacing for tetragonal t (101) refection in as-sintered zirconia specimen was approximately Å. The shift of d spacing toward a higher value reveals the presence of uniform crystal lattice strain (37, 63). Our results showed that shifting was more pronounced for the air abraded groups, indicating a larger amount of uniform strain in these groups, as opposed to the G and GP groups, which showed no uniform strain. For example, the d spacing of the AAF group at the treated surface is Å, which returned to unstrained value at a depth of 37 µm ( Å). Ferroelastic domain switching was observed in all treated groups. Under stress, domains within each zirconia grain switched to a different orientation, associated with twinning and resulting to the changes in the intensity of t (200) and t (002) reflections in GIXRD profile (35, 38), and consequently strengthening zirconia (38, 64). Our results showed that the ferroelastic domain switching ratio in the as-sintered specimens was 35

47 approximately 0.6. For the AAF and AAC groups, the FDS ratio returned to a base value at the depth of 52 µm, and 20 µm respectively. The G and GP showed relatively shallow FDS, where they returned to a base value at a depth of 5 and 3 µm respectively. The broadening of the FWHM indicates a non-uniform strain, plastic deformation and possibly grain size refinement (65) In the present study, the FWHM of as-sintered main tetragonal refection is 0.33 in the absence of strain. The GP group showed the highest broadening FWHM (0.88), followed by AAF, G and AAC respectively. The FWHM and asymmetric factor of the main tetragonal t (101) reflection are associated with the formation of rhombohedral or pseudo-cubic phase shoulder in the XRD profile and strained tetragonal phase (23, 32, 66, 67). In this study, the GP group showed the highest asymmetry factor, which may be caused by the formation of textured layer from polishing. The relationship between the asymmetry factor and the depth of zirconia specimen may not be a linear relationship, as indicated by fluctuation of the asymmetric factor in the AAF and AAC groups. The relationship of the volume fraction of monoclinic phase and the depth of compressive stress layer was observed in this study. We propose that the depth of stressinduced phase transformation layer is dependent on the magnitude of applied-stresses. However, it may not match the compressive stress level. From example, the compressive stress layer of the AAF groups was approximately 50 µm, which is thicker than the level of monoclinic phase (37 µm). It is obvious that the peak position of the main tetragonal t (101) reflection, the ferroelastic domain switching ratio, the Full-Width at Half-Maximum, the asymmetry factors of the main tetragonal t (101) reflection, and the volume fraction of monoclinic phase were affected by the residual stress-produced surface modifications. The depth of the compressive stress layer was determined as the depth at which the XRD profile features retuned to a base value corresponding to unstrained control specimens. Our study revealed that biaxial flexural strength depends on the relationship between compressive stress layers and subsurface damage. For zirconia air abraded with fine particles group, subsurface damage extended approximately 10 µm, whereas the compressive stress layer was approximately 50 µm. The adverse effect from the flaws was overcome by the benefit from compressive stress layer, resulting in an increase in 36

48 mean flexural strength ( ± MPa) when compared to the as-sintered group. This was also true for the ground and polished group, where the depth of the subsurface damage (~3 µm) was less than the thickness of the compressive stress layer (~5 µm). Even though the compressive stress layer was quite thin, it still enhanced the mean biaxial flexural strength for the GP group ( ± MPa). On the other hand, the beneficial aspects of the compressive stress layer were overcome by the extent of the damage in the AAC and G groups. Cracks obtained from air abrasion with coarse particles extended up to ~ 80 µm, and decreased the mean biaxial flexural strength ( ± MPa). Similar to the G group, where the damage depth (~25 µm) was greater than the depth of the compressive stress layer (~10 µm), resulting to a decrease in the mean biaxial flexural strength ( ± MPa). The AAC group had the highest Weibull modulus, associated with a significant decrease in mean biaxial flexural strength, which would be expected for the negative control group. The AAF group had a significantly higher Weibull modulus and flexural strength than the control group, indicating that the AAF has better reliability with a smaller population of flaws and is less likely to break at a stress much lower than the mean strength value. In contrast, even though the Weibull modulus of the G group was not significant from the others, it had the lowest modulus value, suggesting the presence of a wider population of flaws (Figure 17). Within the limitations of this study, according to the manufacturer s recommendations, it is clinically acceptable to air abrade the intaglio surface with fine alumina particles. If occlusal adjustments are required, gently grinding with diamond burs and careful polishing with recommended polishing kits for zirconia is also an acceptable procedure. These two treatments led to the highest strength value, without being detrimental to the reliability. 37

49 Figure17. Effect of surface treatments on subsurface damage in zirconia specimens; A) Air abraded with fine particles; B) Air abraded with coarse particles; C) Ground with diamond bur; and D) Ground and polished 38