Abstract
Objectives
In this brief and selective overview, basic factors contributing to aqueous induced dissolution, phase changes of zirconia and mechanical strength degradation of dental ceramics are considered.
Methods
Dissolution is important for porcelain and glass-ceramic materials as they predominantly contain a silica glass structure with various other cations incorporated which disrupt the silica network. In glass-ceramic materials and porcelains containing crystalline components the situation is more complex as the difference in the dissolution rate may increase the surface roughness and contribute to abrasion of the opposing definition.
Results
Factors contributing to the aqueous dissolution of silicate based dental ceramics and induced destabilization of Y-TZP zirconia or low temperature degradation (LTD) are considered. It is also noted that there have been a number of reports of spontaneous fracture of 3Y-TZP specimens because of LTD in the absence of external stress. A simple analysis of this situation is presented which indicates a strong specimen-size effect and also implicates the role of the remnant “pseudo-grain” structure resulting from the pressing of spray-dried powder agglomerates.
Significance
The final section of this review addresses the consequences of various environments on crack growth and strength degradation. In some situations it is suggested that the aqueous environment may enhance the strength of restorative materials. In all instances discussed above the clinical consequences are pointed out.
1
Introduction
The oral cavity is a potentially very hostile environment despite the critical role of the salivary glands in buffering this situation. The saliva not only buffers but also contains Ca 2+ and PO 4 3− ions that mediate dissolution and facilitate deposition of apatite on enamel. The pH range experienced in the oral cavity ranges from 1 to 10 with the highly acidic conditions associated with reflux of gastric fluid . More typically the pH of ingested beverages ranges from 3 to 8 with carbonated acidic beverages including sport drinks and energy drinks that have stimulants such as guaronic acid with low pH values.
The bacteria naturally present in the oral cavity are also highly effective in developing low pH conditions. A typical measurement of the pH changes with time of inoculated bacteria from a patient, that developed into a plaque biofilm on enamel, fed with pulses of sucrose in an artificial mouth model varies from 3 to 6 . In the above case the essentially static situation coupled with limited amount of sucrose, plus the buffering by enamel dissolution leads to a gradually rising pH following each sucrose addition. The low pH (3.5) that develops can result in substantial dissolution of enamel and onset of caries.
The major difference between external ingested beverages and caries beneath biofilms is that the latter are static adjacent to the adherent biofilm whereas the former are more general and react with the entire exposed teeth surfaces. The recent widespread use of energy and sports drinks is contributing to the world-wide incidence of tooth erosion (corrosion), especially of young people. Apart from acidic ingested fluids, Ccahuana et al. have also shown that clinically applied acid fluoride varnishes have low pH (3.6–3.9) and contain HF. To date the emphasis has been on the role of low pH but a number of common foods (Lima and soy beans, spinach) have pH values of 8 and higher .
Apart from the range of pH, the oral cavity experiences a substantial temperature range (0–67 °C ). The higher temperatures would accelerate the kinetics of any reactions between the liquids and the ceramic surfaces present. This has tended to influence the current ISO Standard 6872 test for evaluation of the solubility of dental ceramics, which measures the weight loss after various time periods at 80 °C. There have been a range of studies that have also investigated the influence of the pH on the release of ions from various dental ceramics . The most recent study by Esquivel-Upshaw et al. suggests that high pH conditions are most deleterious to the mass loss from glass based systems because the high pH is more effective in breaking up the silica glass framework.
The other major factor contributing to aqueous degradation of dental ceramic is the presence of stress on these structures placed in the oral cavity. There are three major sources of stress on or within materials in the mouth: (i) associated with mastication loads developed by the jaw muscles that may generate substantial contact and flexural stresses, (ii) associated with residual stresses that are present upon cooling after fabrication because of thermal expansion mismatch or temperature gradients, and (iii) localized stresses in materials because of their microstructure and presence of multiple components . For example, in the glass matrix of veneering porcelains containing leucite crystals because of the thermal expansion mismatch, at grain boundaries in single-phase non-cubic structured materials such as alumina (hexagonal) and zirconia (tetragonal) because of thermal expansion anisotropy (TEA) of these structures . These various sources of stress, especially in the presence of a corrosive medium, may initiate enhanced dissolution and/or local cracking. The extent of any such crack extension will be highly dependent on the size of the feature or body responsible for the stress and its magnitude. For instance in the case of TEA induced stresses these will be of short range typically less than the grain size but which can assist with the initiation of cracking and phase changes. These stresses combine with the known propensity for brittle ceramics to undergo stress corrosion cracking potentially resulting in time and fatigue dependent strength degradation .
The dental ceramic material systems considered in this article are: (i) feldspathic glass based porcelains, (ii) glass-ceramic systems, and (iii) sintered crystalline ceramics with a focus on zirconia. While in some situations the surface roughness developed upon exposure to a hostile environment is seen as detrimental to the wear of the opposing antagonist, it is also considered as beneficial to assist with the adhesive bonding of such etched and pitted materials. In this article some of the basic issues that determine the dissolution of ceramic materials in liquid environments will be considered. In the following section the role of liquids and temperature on the phase changes of zirconia are considered. The final section will consider the additional role of stresses on the mechanical response.
2
Dissolution of glass and ceramic materials
The focus of this section will be on the dissolution of dental ceramic materials in aqueous environments.
The dissolution of glass is an active area of research not only for dental applications but diverse fields ranging from nuclear waste containment, preservation of stained glass in the polluted urban environment to domestic kitchen ware. The dissolution of glass depends strongly on the composition of the glass and the leaching environment. The environment may also change with time depending upon whether it is static or flowing. For instance in a static environment the leaching out of alkali elements will increase the pH of the solution, which may then lead to a greater rate of dissolution of the silica matrix. Hench and Clark , based upon the chemistry of a glass and associated (effective) pH, identified 5 types of dissolution of glass: Type 1, pure silica in neutral pH resulting in a very thin (<5 nm) hydrated layer. Type 11 and 111 are durable glasses (low alkali content with pH <9) that develop a hydrated layer (<200 nm) denuded of alkali elements (see Fig. 1 ). Type IV glasses have a high alkali content where the hydration layer does not limit continued dissolution. Type V corresponds to continuous non-selective dissolution as is typical of all glasses with an effective pH in an aqueous environment of above 9. There is also a difference in the dissolution rate and mechanism of some glasses, which are phase separated, such as borosilicate glasses. Where the dissolution in such glasses proceeds faster in one phase than the other and whether the phases are truly separate or there is a percolation (interpenetration) of the two phases is important.
White critically reviewed the basic theories of corrosion of glasses and ceramics. He considered that the major factor controlling the dissolution rate of materials of general utility, namely relative low solubility, is usually an acid–base reaction which is surface reaction controlled. That is reactions occur between H + , H 3 O + or H 2 O molecules with cations on the surface. If a reaction product forms then the dissolution rate can be slowed down or almost stopped. For example the dissolution of enamel (apatite) in the presence of F ions resulting in a very low solubility CaFPO 4 layer formation . In the case of glasses the dissolution appears to be controlled by the ion exchange mitigated by diffusion of H + and H 3 O + into the glass with the counter diffusion of leachable ions out of the glass. There are two major differences between a crystalline ceramic and glass with the same composition: (i) The coordination and bonding energy of the composing elements is stronger for the crystalline phase compared to the glass and this reduces the dissolution rate of the crystalline material, (ii) a crystalline ceramic however has dislocations, precipitates and grain boundaries that generate residual stresses not present in a simple glass. These localized stress sites in crystalline materials enhance local etch type dissolution such as occurs in acidic dissolution of the cores of the enamel apatite crystallites. This difference in dissolution rate for crystalline versus glass is clearly demonstrated by the difference in the dissolution rate of albite (NaAlSi 3 O 8 ) in its crystalline or glassy state as shown in Fig. 2 . This figure shows that leaching of Na and Al is faster in the glass than crystalline form while that of the Si is not affected.
Devreux et al. consider that the durable glasses (Types I and II in Hench and Clarke’s classification) may be further subdivided into three classes depending upon the included oxides in the glass. These are Class A which consists of the silica matrix, Class B the water soluble species namely the alkali, boron and phosphorous oxides, and Class C the less soluble oxides such as Al, Zr, Re, Ti and Sn. These three classes of components are distributed randomly in the structure with oxygen atoms as the bonding between the elements. While it is acknowledged that the valence of the element constituents is important so is the coordination number they occupy in the silica matrix, for instance Zr is in octahedral while Al is in a tetrahedral environment. The classic model of glass solubility in water suggests the diffusion of hydronium and alkali ions occurs in parallel with dissolution of the silica network. At pH conditions <9 the dissolution rate of silica bonds is very slow. On the basis that in most instances different species such as alkali and boron atoms dissolve together it is argued that Class B atoms only dissolve when they are present at a free surface. If the volume of solution is large or continuously replaced then condensation of species on the surface does not occur.
Milleding et al. investigated various glass based dental ceramics for ion leaching in water (RT) and acetic acid (pH 4 at 80 °C). All these materials had a relatively narrow range of silicon content, although some contained leucite (K 2 O·Al 2 O 3 ·SiO 2 ) crystals whereas others were primarily glass. They found that the lowest fusing ceramic, which had the highest ion release rate, had the highest Na and K content in order to reduce its melting temperature. Most of the glasses had relatively low release rates of Si ions suggesting that the leaching of Na + and K + ions would create a porous hydrated silica layer. The leucite containing materials all had high Na and K release rates with the ratio of Na to K almost constant and also relatively high Si and Al release rates indicative of the dissolution of the leucite crystals in the glass ceramic. A typical example of a leucite containing glass ceramic after acid exposure is shown in Fig. 3 a . It is clear that the leucite crystals have dissolved somewhat more rapidly than the glass matrix, note also the rounding (blunting) of the cracks in the glass surrounding the leucite. In Fig. 3 b is an example of HF etched lithium disilicate containing glass ceramic where the surrounding glass matrix, which has reduced silicon content with the formation of the high volume fraction of lithium disilicate crystalline phase, has more rapidly dissolved .
What are the clinical implications? As pointed out by Oh et al. the surface dissolution and the presence of crystalline phases within porcelain can have an important factor in influencing the wear of enamel and dental ceramics. Glass based porcelains, usually but not exclusively feldspathic, and glass-ceramic materials will undergo relatively low dissolution rates with the rate typically being determined by the amount of alkali (Na, K and Li) plus other glass formers namely B and P present. These glasses will generally develop a slightly hydrated porous layer depleted of the soluble species and will be softer than the underlying glass. These materials, especially the slightly more soluble ones, would thus be expected to wear uniformly and as such be less aggressive toward the antagonist surface. The presence of crystalline species within the glass, as with glass ceramics or blended porcelains to achieve opacity, is likely to show differential dissolution of the glass versus the crystalline phase. In the case of leucite containing glass systems the leucite may dissolve faster than the supporting glass. This may be assisted by the residual tensile stress and twinning deformation of the crystallites as a consequence of the cubic to tetragonal transformation at 650 °C and associated shear deformation plus the higher CTE of this phase than the glass matrix. On the other hand, for lithium disilicate containing glass ceramics the crystalline Li 2 SiO 3 phase appears to dissolve more slowly than the glass matrix resulting in the development of a rough surface that enhances the wear rate of an antagonist . Heintze et al. in their systematic review of in vitro wear tests compared the results of 3 glass ceramic materials, a glass glaze material, a leucite and lithium disilicate glass ceramic with the latter showing higher antagonist wear. With respect to crystalline alumina and zirconia, there are minimal indications of dissolution in aqueous environments . The consequences of aqueous exposure for mechanical integrity of zirconia will be considered in the following sections.
2
Dissolution of glass and ceramic materials
The focus of this section will be on the dissolution of dental ceramic materials in aqueous environments.
The dissolution of glass is an active area of research not only for dental applications but diverse fields ranging from nuclear waste containment, preservation of stained glass in the polluted urban environment to domestic kitchen ware. The dissolution of glass depends strongly on the composition of the glass and the leaching environment. The environment may also change with time depending upon whether it is static or flowing. For instance in a static environment the leaching out of alkali elements will increase the pH of the solution, which may then lead to a greater rate of dissolution of the silica matrix. Hench and Clark , based upon the chemistry of a glass and associated (effective) pH, identified 5 types of dissolution of glass: Type 1, pure silica in neutral pH resulting in a very thin (<5 nm) hydrated layer. Type 11 and 111 are durable glasses (low alkali content with pH <9) that develop a hydrated layer (<200 nm) denuded of alkali elements (see Fig. 1 ). Type IV glasses have a high alkali content where the hydration layer does not limit continued dissolution. Type V corresponds to continuous non-selective dissolution as is typical of all glasses with an effective pH in an aqueous environment of above 9. There is also a difference in the dissolution rate and mechanism of some glasses, which are phase separated, such as borosilicate glasses. Where the dissolution in such glasses proceeds faster in one phase than the other and whether the phases are truly separate or there is a percolation (interpenetration) of the two phases is important.
White critically reviewed the basic theories of corrosion of glasses and ceramics. He considered that the major factor controlling the dissolution rate of materials of general utility, namely relative low solubility, is usually an acid–base reaction which is surface reaction controlled. That is reactions occur between H + , H 3 O + or H 2 O molecules with cations on the surface. If a reaction product forms then the dissolution rate can be slowed down or almost stopped. For example the dissolution of enamel (apatite) in the presence of F ions resulting in a very low solubility CaFPO 4 layer formation . In the case of glasses the dissolution appears to be controlled by the ion exchange mitigated by diffusion of H + and H 3 O + into the glass with the counter diffusion of leachable ions out of the glass. There are two major differences between a crystalline ceramic and glass with the same composition: (i) The coordination and bonding energy of the composing elements is stronger for the crystalline phase compared to the glass and this reduces the dissolution rate of the crystalline material, (ii) a crystalline ceramic however has dislocations, precipitates and grain boundaries that generate residual stresses not present in a simple glass. These localized stress sites in crystalline materials enhance local etch type dissolution such as occurs in acidic dissolution of the cores of the enamel apatite crystallites. This difference in dissolution rate for crystalline versus glass is clearly demonstrated by the difference in the dissolution rate of albite (NaAlSi 3 O 8 ) in its crystalline or glassy state as shown in Fig. 2 . This figure shows that leaching of Na and Al is faster in the glass than crystalline form while that of the Si is not affected.
Devreux et al. consider that the durable glasses (Types I and II in Hench and Clarke’s classification) may be further subdivided into three classes depending upon the included oxides in the glass. These are Class A which consists of the silica matrix, Class B the water soluble species namely the alkali, boron and phosphorous oxides, and Class C the less soluble oxides such as Al, Zr, Re, Ti and Sn. These three classes of components are distributed randomly in the structure with oxygen atoms as the bonding between the elements. While it is acknowledged that the valence of the element constituents is important so is the coordination number they occupy in the silica matrix, for instance Zr is in octahedral while Al is in a tetrahedral environment. The classic model of glass solubility in water suggests the diffusion of hydronium and alkali ions occurs in parallel with dissolution of the silica network. At pH conditions <9 the dissolution rate of silica bonds is very slow. On the basis that in most instances different species such as alkali and boron atoms dissolve together it is argued that Class B atoms only dissolve when they are present at a free surface. If the volume of solution is large or continuously replaced then condensation of species on the surface does not occur.
Milleding et al. investigated various glass based dental ceramics for ion leaching in water (RT) and acetic acid (pH 4 at 80 °C). All these materials had a relatively narrow range of silicon content, although some contained leucite (K 2 O·Al 2 O 3 ·SiO 2 ) crystals whereas others were primarily glass. They found that the lowest fusing ceramic, which had the highest ion release rate, had the highest Na and K content in order to reduce its melting temperature. Most of the glasses had relatively low release rates of Si ions suggesting that the leaching of Na + and K + ions would create a porous hydrated silica layer. The leucite containing materials all had high Na and K release rates with the ratio of Na to K almost constant and also relatively high Si and Al release rates indicative of the dissolution of the leucite crystals in the glass ceramic. A typical example of a leucite containing glass ceramic after acid exposure is shown in Fig. 3 a . It is clear that the leucite crystals have dissolved somewhat more rapidly than the glass matrix, note also the rounding (blunting) of the cracks in the glass surrounding the leucite. In Fig. 3 b is an example of HF etched lithium disilicate containing glass ceramic where the surrounding glass matrix, which has reduced silicon content with the formation of the high volume fraction of lithium disilicate crystalline phase, has more rapidly dissolved .
What are the clinical implications? As pointed out by Oh et al. the surface dissolution and the presence of crystalline phases within porcelain can have an important factor in influencing the wear of enamel and dental ceramics. Glass based porcelains, usually but not exclusively feldspathic, and glass-ceramic materials will undergo relatively low dissolution rates with the rate typically being determined by the amount of alkali (Na, K and Li) plus other glass formers namely B and P present. These glasses will generally develop a slightly hydrated porous layer depleted of the soluble species and will be softer than the underlying glass. These materials, especially the slightly more soluble ones, would thus be expected to wear uniformly and as such be less aggressive toward the antagonist surface. The presence of crystalline species within the glass, as with glass ceramics or blended porcelains to achieve opacity, is likely to show differential dissolution of the glass versus the crystalline phase. In the case of leucite containing glass systems the leucite may dissolve faster than the supporting glass. This may be assisted by the residual tensile stress and twinning deformation of the crystallites as a consequence of the cubic to tetragonal transformation at 650 °C and associated shear deformation plus the higher CTE of this phase than the glass matrix. On the other hand, for lithium disilicate containing glass ceramics the crystalline Li 2 SiO 3 phase appears to dissolve more slowly than the glass matrix resulting in the development of a rough surface that enhances the wear rate of an antagonist . Heintze et al. in their systematic review of in vitro wear tests compared the results of 3 glass ceramic materials, a glass glaze material, a leucite and lithium disilicate glass ceramic with the latter showing higher antagonist wear. With respect to crystalline alumina and zirconia, there are minimal indications of dissolution in aqueous environments . The consequences of aqueous exposure for mechanical integrity of zirconia will be considered in the following sections.
3
Zirconia in aqueous environments
The most common clinically used zirconia ceramic is 3Y-TZP (containing 3 mol% Y 2 O 3 ) with or without minor additions of silica and alumina. This material is typically sintered between 1350 and 1600 °C, resulting in grain sizes that range from 0.3 to 1 μm. Considering the composition and sintering temperatures the most recent Y 2 O 3 –ZrO 2 phase diagram suggests that at equilibrium typically 60–80% should be tetragonal with the remainder cubic phase. Subsequent more detailed analysis indicates a much more complex situation with the presence of 5 possible phases with 3 tetragonal compositions t, t′ and t″ (see Fig. 4 ). A major motivation for a better understanding of this situation arises from the thermal barrier coating community as long term high temperature exposure results in the formation of t phase that can transform to m on cooling and so jeopardize the lifetime of engine components. The presence of an initial t′ phase, a slightly Y 2 O 3 rich phase compared to t phase, which is dependent upon the initial composition of the sintered material is also significant. The partitioning of the t′ phase into the equilibrium t and c phases is very slow because of the slow high temperature diffusion of Y 3+ ions. The T o (t/m) and T o (c/t) boundaries, both of which are still relatively poorly defined, dictate whether t transforms to m and c to t″ upon cooling. For Y-TZP materials used in dentistry the fine grain size and dense structures constrain the t phase to RT but in the presence of stress or moisture the t to m phase transformation will occur. Lipkin et al. also clearly show that the partitioning of the t′ is dependent upon time and temperature and is predictable based upon the Holloman–Jaffe parameter. The major problem is that it is difficult to separate the 3 t phases and c phase by conventional XRD and careful deconvolution of the 400 diffraction peaks are required. This has clinical implications as the extent of the t and c phase partitioning has a strong bearing on the metastability of the t phase especially the low temperature degradation (LTD).
It has long been recognized that provided the grain sizes of the t phase are in the range 0.2–1.5 μm it is metastable at room temperature in densely sintered materials and able to be transformed by stress, thereby developing transformation toughening . However it also been appreciated that Y-TZP is susceptible to moisture induced t to m transformation that is maximized at 200–250 °C . The mechanism by which this occurs is associated with the diffusion of OH − ions into the Y-TZP lattice which assists with the nucleation of a critical m region in the lattice that can trigger part of the grain to transform. There is a considerable discussion on the mechanisms in two recent review papers . The destabilization is also assisted by the presence of stress, electrical fields and impurity segregation upon holding specimens at elevated temperatures . Other suggestions for the destabilization have been the moisture assisted release of Y stabilizing ions from the Y-TZP lattice .
The catastrophic consequences of LTD manifested in 2002 with the spontaneous shattering of more than 500 Y-TZP femoral hip prosthesis in patients less than 2 years following surgery. Since then there has been considerable awareness of this problem resulting in numerous studies . These have established general agreement as to the prevalence of this phenomenon but only recently has it been shown beyond doubt that the rate of transformation appears to be linear initially (<100 μm depth) with time and dependent upon temperature . These authors also show that moisture partial pressure seems to be less influential on the rate of transformation but has a significant influence on the incubation time for the onset of the transformation.
It has also been shown decades ago that surface treatments such as grinding can have an inoculating effect if the severely deformed surface is resintered above 1200 °C resulting in recrystalization and a finer grain size . More recent work suggests that lower temperature regeneration of the ground or damaged Y-TZP surface at 1000–1100 °C, which while reversing the m to t phase, appears to enhance the initial rate of subsequent LTD transformation (Keuper et al., private communication). In contrast severe grinding of Y-TZP without subsequent annealing treatments appears to initially inoculate Y-TZP from LTD. The as-sintered surface developed more m phase than the ground surfaces. However, longer term studies have shown that after the longer initiation time an accelerated rate of m formation follows.
It is also well known that moisture induced LTD leads to enhanced roughness of Y-TZP . This has detrimental consequences on the wear rate of polyethylene acetabular cups, increases friction and wear of the opposing counterface and in the dental clinical scenario would include the opposing dentition . LTD also led to the spontaneous fracture of hip balls in patients as it occurred in the area about the trepanned morse taper core of the ball adjacent to the Ti stem . There have also been numerous other reports of spontaneous fracture of samples during LTD testing .
For the spontaneous fracture of a Y-TZP sample due to LTD the internal stress must exceed the tensile strength of the material at the temperature considered. In Y-TZP, like all zirconia transformation toughened materials, the mechanical properties are strongly temperature dependent . If we consider a long rectangular beam or thin circular disk without externally applied stress then the only source for the internal tensile stresses is to compensate for the surface compressive stresses associated with the t to m transformation. A simple estimate of the thickness of the transformed zone required to cause spontaneous fracture is to consider a cross-section with the outer layer transformed m phase and an internal t phase. The compressive stresses developed in the outer layer must balance the tensile stresses developed internally. That is:
A t σ t + A u σ u = 0
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