ABSTRACT 
 
 
Title of Dissertation: EFFECT OF GLASS JOINS ON PERFORMANCE OF 
LAYERED DENTAL CERAMIC SYSTEMS 
 
Mey A.Saied, Doctorate of Philosophy 2009 
 
Dissertation Directed by: Associate Professor Isabel K. Lloyd 
Department of Materials Science and Engineering 
 
 
     
 
 
 Layered structures can be used to address the competing needs of systems like 
dental crown restorations where the exterior needs to be aesthetic and the interior 
needs to be strong and fatigue resistant.  Dental crowns typically have an aesthetic 
porcelain veneer layered on a strong, fatigue resistant ceramic or metallic core.  In 
current restorations, even when the core is shaped by a computer-aided design and 
manufacturing (CAD/CAM) or solid-freeform fabrication processes, the veneer is 
applied in sequential layers.  This process is labor intensive, time consuming and may 
not optimize the long-term performance properties of the veneer layer.  If the core 
and veneer layers were to be independently fabricated and then joined, their 
individual and the veneer-core system performance could be optimized.  Some groups 
have explored the possibility of joining with filled epoxies, which is easier, but may 
not be long-lasting.  In this project we explore the possibility of using more durable 
glassy joins.  Dense, thermal-expansion-matched (to the core and veneer glass) joins 
can be fired at temperatures far enough below the melting and/or slumping 
temperatures to join veneers to cores without degradation.  In this study, we design 
and fabricate joining glasses for bonding porcelain veneers to ceramic cores, 
specifically to dental aluminas and zirconias.  We study the chemical bonding and 
mechanical integrity of the resulting layers.  Finally, we assess the effects of glass 
joins on performance of layered dental ceramic systems. 
  
 
 
 
 
 
 
 
 
 
 
 
 
 
 
EFFECT OF GLASS JOINS ON PERFORMANCE 
OF LAYERED DENTAL CERAMIC SYSTEMS 
 
 
By 
 
Mey A. Saied 
 
 
 
 
Dissertation submitted to the Faculty of the Graduate School of the 
University of Maryland, College Park in partial fulfillment 
of the requirements for the degree of 
Doctor of Philosophy  
2009 
 
 
 
 
 
 
 
 
 
 
 
 
 
Advisory Committee: 
 Associate Professor Isabel K. Lloyd, Chair/Advisor 
 Dr. Brian R. Lawn, Co-advisor 
 Professor Lourdes Salamanca-Riba 
 Professor Manfred Wuttig 
 Associate Professor F. Patrick McCluskey 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
?Copyright by 
Mey A. Saied 
2009 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
 ii 
 
 
 
 
 
 
 
 
 
 
 
 
 
 
Dedication 
 
For my father, Abdullah, my mother, Fathia, my brother Hatim and especially my 
niece, Iman, the future. 
 
 
 
 
 
 
 
 iii 
 
Acknowledgements 
I would like to thank Dr. Isabel K. Lloyd  for supervising me throughout these 
very interesting and intellectually stimulating last five years.  She helped to inspire 
my curiosity for this subject matter.  In particular her collaboration with Dr. Brian 
Lawn on my behalf was a pivotal part of my completion.  Thank you for the late 
nights, and for your dedication to my progress.  
I would also like to thank Dr. Brian Lawn, NIST for being a wonderful co-
advisor, for his expert advice on mechanical properties and his outstanding logic.  He 
gave me unequivocal assistance, energy, and advice and treated like family, 
understanding my strengths and weaknesses, and having faith in me even when I 
faltered.  I owe him a great deal, and will miss him terribly. 
I wish to extend my deep gratitude to Dr. James Lee for his help and advice 
throughout.  He is a joy for anyone lucky enough to work with him.  In him I have a 
brother, a colleague and a friend.  At NIST, I would also like to thank Dr. Wolfgang 
Haller, Mr. George Quinn, and Mr. Ed Parry. 
To Dr. Dianne Rekow, New York University College of Dentistry, for taking 
a chance on me, pointing me in the right direction, and putting in so much effort and 
energy to help me succeed?many thanks.  I promise to pay it forward.  Also at NYU, 
I would also like to thank Dr. Van Thompson, Dr. Yu Zhang, Dr. Paolo Coelho and 
Ms. Elizabeth Clarke for their constant support, assistance and friendship. 
I would like to thank Dr. Otto Wilson of Catholic University for his advice, 
support, and friendship.  His input on sol-gel processing was invaluable, as were his 
 
 iv 
introductions to the Vitreous State Laboratory (VSL) group at Catholic University.  
Also at VSL, I would like to thank Dr. Isabelle Muller, Dr. Fernando Perez-Cardenas 
and the staff there who were essential for the glass processing segment of this project. 
At the University of Maryland, my gratitude is due to Dr. Kathleen Hart, who 
always listened and stepped in to make things happen, Dr. Robert Briber, for helping 
whenever funding hiccups occurred and Dr. Phil Piccoli, for his friendly assistance 
with electron microscopy. 
 I am thankful for financial support from the NRSA Minority pre-doctoral 
fellowship award grant NIDCR 5F31DE017297. 
Finally, I would like to thank my family for their continued patience, 
encouragement and faith even when they are so far away.  My father, who opened my 
eyes to the world, the possibilities, gave everything he could and helped me see that 
nothing could hold me back.  My mother, who withstood the storm of traditions, and 
opened her mind, and fought for me always.  My brother, who was my friend first, 
my protector second, and who always believed in me. I would be nothing without 
you. 
 
 
 
 v 
 
Table of Contents 
 
List of Tables .........................................................................................................vii 
 
List of Figures.......................................................................................................viii?
 
Chapter 1: ? Introduction.........................................................................................1?
1.1? Background and Approach ....................................................................1?
1.2? The Challenge .......................................................................................2?
1.3? Outline of Thesis...................................................................................3?
 
Chapter 2: ? Design & Processing of Tailored Glasses............................................6?
2.1? Introduction...........................................................................................6?
2.2? Background...........................................................................................7?
2.3? Materials Selection................................................................................8?
2.3.1? Choice of core and veneer materials ......................................................9?
2.3.2? Choice of glass compositions .............................................................. 12?
2.4? Preliminary Tests Using Industrial Glasses..........................................18?
2.5? Glass Preparation by Sol-Gel Processing .............................................18?
2.5.1? Sol-gel wet mixing.............................................................................. 18?
2.5.2? Melt and quench process ..................................................................... 25?
2.6? Material Characterization ....................................................................28?
 
Chapter 3:? Fused Joins: Laminar Ceramic Structures and Glass-Ceramic 
Interface Chemistry .......................................................................... 37?
3.1? Introduction.........................................................................................37?
3.2? Background.........................................................................................39?
3.2.1? Glass as a seal ..................................................................................... 39?
3.2.2? Traditional applications in glass/ceramic joining ................................. 41?
3.2.3? Dental all-ceramic restorations ............................................................ 42?
3.3? Experimental Methods.........................................................................44?
3.3.1? Preparation of bilayer specimens ......................................................... 44?
3.3.2? Characterization of layer interfaces ..................................................... 47?
3.4? Results ................................................................................................48?
3.4.1? Screening tests .................................................................................... 48?
3.4.2? Optical microscopy ............................................................................. 49?
3.4.3? Electron microprobe analysis .............................................................. 54?
3.5? Discussion and Summary ....................................................................56?
 
 vi 
 
Chapter 4: ? Mechanical Evaluation...................................................................... 73?
4.1? Introduction.........................................................................................73?
4.2? Background.........................................................................................74?
4.3? Experimental Methods.........................................................................77?
4.4? Results ................................................................................................80?
4.4.1? Qualitative observations ...................................................................... 80?
4.4.2? Quantitative analysis ........................................................................... 85?
4.5? Discussion and Summary ....................................................................92?
 
Chapter 5:? Summary and Future Work.............................................................. 96?
5.1? Summary.............................................................................................96?
5.2? Future Work........................................................................................97?
 
Appendix  ........................................................................................................... 99?
 
References  ......................................................................................................... 102?
 
Curriculum Vitae ................................................................................................ 110?
  
 
 
 vii 
 List of Tables 
Table 2.1? Properties of veneer and core materials........................15 
 
Table 2.2? Properties and compositions of Schott industrial lead-fre alumino-
borosilicate and bismuth glases.......................................16 
 
Table 2.3? Properties and compositions of Fero industrial lead-fre bismuth glases
 ......................................................17 
 
Table 2.4? Compositions of Winkelmann and Schott Factor-formulated lab glases 
for bonding electronic alumina substrates................................20 
 
Table 2.5? Compositions of Appen Factor-formulated lab-tailored alumina-
compatible and zirconia-compatible glases...............................21 
 
Table 2.6? Calculated properties of selected join Appen glas compositions.....22?
 
Table 3.1? Finger test, to determine ?pas? (P) or ?fail? (F) for bonded alumina- and 
zirconia-based specimens bonded with Winkleman?Schott (WS) and Appen (A) 
glases. Premature failures during prep indicated as FP......................50 
 
Table 3.2? Cutting test, to determine ?pas? (P) or ?fail? (F) for bonded alumina- and 
zirconia-based specimens bonded with Winkleman?Schott (WS) and Appen (A) 
glases . ......................................................51?
 
 
 vii 
 
 List of Figures 
 
Figure 2.1? Dental crowns; (a) Veneered zirconia crown, with cement application, 
prior to bonding to supporting structure (Courtesy P. Coehlo). (b) Failed al-ceramic 
(Alumina) crown, with veneer, core and tooth structure clearly visible (Courtesy 
Kenneth Malament).................................................10 
 
Figure 2.2? Traditional methods of making dental ceramic restorations (courtesy G. 
Zhang, E.D. Rekow, V.P. Thompson and Y. Wang).........................11 
 
Figure 2.3? Manufacturer recommended firing cycles for a) Nobel Rondo alumina-
veneering porcelain, and b) Sakura Interaction zirconia veneering porcelain.......14 
 
Figure 2.4? Lead-fre commercialy available glases from Schott (Bad Sackingen, 
Germany and Fero, USA). (a) Schott glas No. 8415, CTE 7.8x10
-6
C
-1
 on electronic 
alumina (EA) sintered at 800 C, showing piting; (b) Fero EG2964 CTE 8.5x10
-6
C
-1
 
on  alumina, showing poor weting; (c) Schott G01-8249 CTE 10.1x10
-6
C
-1
 on 
zirconia showing discoloration; (d) Schott G01-8421 CTE 9.7x10
-6
C
-1
 fired at 1000C 
on alumina, showing improved melting but crazing.........................23 
 
Figure 2.5? Sol-gel/wet mixing glas precursor preparation proces diagram.....26 
 
Figure 2.6? TGA of gels to ases minimum temperature for complete removal of 
volatiles. ......................................................27 
 
Figure 2.7? Diagram of glas preparation from sol-gel precursors.............30 
 
Figure 2.8? Crazing in Appen-formulated glas, AP650 (CTE = 6.5 x 10
-6
 C
-1
) fired 
at 800 C for 10 mins on Procera alumina substrate. No crazing, porosity is visible, 
and weting is good. Mate finish is due to surface roughnes..................31 
 
Figure 2.9? Thermomechanical Analysis (TMA) of alumina system with 
dilatometry data of tailored glases overlaid. The matching CTEs over the sintering 
temperature range are represented by the slopes of the graphs. Procera is the alumina 
core material, AP0700 is the glas join, and Rondo is the veneering porcelain.....32 
 
Figure 2.10? Thermomechanical Analysis (TMA) of zirconia systems with 
dilatometry data of tailored glases overlaid. The matching CTEs over the sintering 
temperature range are represented by the slopes of the graphs. Cyrtina is the zirconia 
core material, AP1040 is the glas join, and Sakura is the veneering porcelain.....33 
 
 
 
 
 
 ix 
Figure 2.11? Dilatometry of AP700 alumina systems joining glas. The red line 
represents a slope of y = mx, where m represents the coeficient of thermal expansion 
at the denoted temperature range and is calculated as 7.00 x 10
-6
 C
-1
 for this glas. T
g
 
and T
s
 (transition and softening temperatures respectively) are denoted by the 
intersection of the gren dotted lines representing change in slope, or transition of 
expansion behavior. The blue shaded elipse identifies the area where calculated and 
measured thermal expansion behaviors match..............................34 
 
Figure 2.12? Dilatometry of AP1040 zirconia systems joining glas. The red line 
represents a slope of y = mx, where m represents the coeficient of thermal expansion 
at the denoted temperature range and is calculated as 10.40 x 10
-6
 C
-1
 for this glas. T
g
 
and T
s
 (transition and softening temperatures respectively) are denoted by the 
intersection of the gren dotted lines representing change in slope, or transition of 
expansion behavior. The blue shaded elipse identifies the area where calculated and 
measured thermal expansion behaviors match..............................35?
 
Figure 3.1? Al-ceramic crown with join interlayer translated to a flat model crown.
 ......................................................40 
 
Figure 3.2? Firing cycles for alumina systems. a) Glas slurry was applied to both 
veneer and core layers, then biscuit-fired at 750 C (to promote adhesion of the glas 
to the substrates). b) Veneer and core blocks were then sintered at 830 C (to complete 
the join). ......................................................45 
 
Figure 3.3? Firing cycles for zirconia systems. a) Glas slurry was applied to both 
veneer and core layers, then biscuit-fired at 700 C (to promote adhesion of the glas 
to the substrates). b) Veneer and core blocks were then sintered at 800 C (to complete 
the join). ......................................................46 
 
Figure 3.4? Showing joins in glas-bonded Rondo porcelain and Procera alumina 
bi-layer.  Glas in (a) AP600, in (b) AP700. White horizontal lines are used to 
highlight the veneer/glas junction. Some porosity is evident in (a).............52 
 
Figure 3.5? Showing joins in glas-bonded Sakura porcelain and Cyrtina zirconia 
bi-layer.  Glas in (a) AP600, in (b) AP700. White horizontal lines are used to 
highlight the veneer/glas junction. Some porosity is evident in (a).............53?
 
Figure 3.6a? Microprobe scans of ion concentration across glas-bond interface for 
alumina-core system, for low anneal cycle.  Colored bands indicate interdifusion 
layers. Veneer (V), glas (G) and core (C) zones indicated...................57 
 
Figure 3.6b? Microprobe scans of ion concentration across glas-bond interface for 
alumina-core system, for intermediate anneal cycle.  Colored bands indicate 
interdifusion layers. Veneer (V), glas (G) and core (C) zones indicated.........58 
 
 
 x 
Figure 3.6c? Microprobe scans of ion concentration across glas-bond interface for 
alumina-core system, for high anneal cycle.  Colored bands indicate interdifusion 
layers. Veneer (V), glas (G) and core (C) zones indicated...................59?
 
Figure 3.7a? Microprobe scans of ion concentration across glas-bond interface for 
zirconia-core system, for low anneal cycle.  Colored bands indicate interdifusion 
layers. Veneer (V), glas (G) and core (C) zones indicated...................60 
 
Figure 3.7b? Microprobe scans of ion concentration across glas-bond interface for 
zirconia-core system, for intermediate anneal cycle.  Colored bands indicate 
interdifusion layers. Veneer (V), glas (G) and core (C) zones indicated.........61 
 
Figure 3.7c? Microprobe scans of ion concentration across glas-bond interface for 
zirconia-core system, for high anneal cycle.  Colored bands indicate interdifusion 
layers. Veneer (V), glas (G) and core (C) zones indicated...................62?
 
Figure 3.8a? Potasium difusion in alumina systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................63 
 
Figure 3.8b? Calcium difusion in alumina systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................64 
 
Figure 3.8c? Barium difusion in alumina systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................65 
 
Figure 3.8d? Difusion in alumina systems, a) silicon at veneer/glas (V/G) join 
interface, b) aluminum at glas/core (G/C) interface. For low (yelow), intermediate 
(red) and high (blue) anneal times......................................66?
 
Figure 3.9a? Potasium difusion in zirconia systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................67 
 
Figure 3.9b? Calcium difusion in zirconia systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................68 
 
Figure 3.9c? Barium difusion in zirconia systems, a) veneer/glas (V/G) join 
interface, b) glas/core (G/C) interface. For low (yelow), intermediate (red) and high 
(blue) anneal times..................................................69 
 
 
 xi 
Figure 3.9d? Difusion in zirconia systems, a) silicon at veneer/glas (V/G) join 
interface, b) zirconium at glas/core (G/C) interface. For low (yelow), intermediate 
(red) and high (blue) anneal times......................................70 
 
Figure 4.1? Schematic showing various cracks that can form in ceramic veneer/core 
layers joined by glas, and cemented onto a compliant support layer. At top surface, 
elastic contact can generate outer (O) and inner (I) cone cracks, and plastic contact 
can generate median (M) cracks.  At bottom surface, flexural streses can generate 
lateral cracks.  Al these cracks grow transversely though the layer thicknes, and 
ultimately intersect the join interface....................................75 
 
Figure 4.2? Schematic showing Vickers indentations in ceramic layers joined by 
glas.  Coordinate system shown for measuring lengths c of crack arms at diferent 
distances h from interface (taken at boundary betwen core and glas), for crack in (a) 
veneer and (b) core layers............................................78 
 
Figure 4.3? Vickers indentations in porcelain/glas/alumina layer structures, remote 
from bonding glas interface in (a) porcelain and (b) alumina.  Note esential 
symmetry of corner crack paterns in both materials, indicating absence of significant 
residual streses....................................................81 
 
Figure 4.4? Vickers indentations in porcelain/glas/zirconia layer structures, remote 
from bonding glas interface in (a) porcelain and (b) zirconia.  Note esential 
symmetry of corner crack paterns in both materials, indicating absence of significant 
residual streses....................................................82 
 
Figure 4.5? Cracks from Vickers indentation in glas-bonded porcelain/ alumina bi-
layer. Indentations in a) porcelain and b) alumina. Note crack arests at alumina 
traverses glas bonding layer and arests at alumina interface in a), and penetration 
across glas into porcelain in b). White horizontal lines are used to highlight the 
veneer/glas junction................................................83 
 
Figure 4.6? Cracks from Vickers indentation in glas-bonded porcelain/ zirconia bi-
layer. Indentations in a) porcelain and b) zirconia. Note crack traverses glas bonding 
layer and arests at zirconia interface in a), and penetration across glas into porcelain 
in b). Grey horizontal lines are used to highlight the veneer/glas junction........84 
 
Figure 4.7? Crack sizes c
1
, c
2
, c
3
 and c
4
 versus distance h of indentation center to 
interface in glas-bonded porcelain veneer/alumina core systems. Indentations at P = 
10 N in (a) porcelain and (b) alumina, orientation relative to interface indicated by 
inset. ......................................................86 
 
 
 
 
 
 
 xii 
Figure 4.8? Crack sizes c
1
, c
2
, c
3
 and c
4
 versus distance h of indentation center to 
interface in glas-bonded porcelain vener/ zirconia (Y-TZP) core systems. 
Indentations in (a) porcelain and (b) zirconia, orientation relative to interface 
indicated by inset. In zirconia graph, the red points represent glas AP1040, indents 
made at P = 40 N, while the blue points represent glas AP1020 indents made at P = 
35 N. ......................................................87 
 
Figure 4.9? Stres intensity factor as function at diferent location x = c? h from 
interface, for glas-bonded alumina vener/core system. Indentation corner cracks in 
(a) porcelain (x = h ? c
2
) and (b) alumina (x = c
1
 ? h), indicated by inset. Solid lines 
are empirical fits to data, dashed lines are asymptotic toughnes bounds.........90 
 
Figure 4.10? Stres intensity factor as function at diferent location x = c
1,2
 ? h 
from interface, for glas-bonded zirconia veneer/core system. Indentation corner 
cracks in (a) porcelain (x = h ? c
2
) and (b) zirconia (x = c
1
 ? h), indicated by inset. 
Solid lines are empirical fits to data, dashed lines are asymptotic toughnes bounds. In 
b) for zirconia, the red points represent glas AP1040, indents made at P = 40 N, 
while the blue points represent glas AP1020 indents made at P = 35 N. Note the 
toughnes values converging towards the same values.......................91?
 
 
 
 
 
 
 
1 
 
 Chapter 1:  Introduction 
 
 
 1.1 Background and Aproach 
 
 Dental restorations involve repairing or replacing dentition that has been lost 
due to disease or injury.  Several considerations must be taken into acount when 
such restorations are made. Mechanical properties must be sufficient to withstand the 
functional streses of chewing and, in extreme cases, to bruxing/grinding. Chewing 
may involve loads of several hundreds of N, over 10
6
 cycles/5 years, in a hostile oral 
environment, making restorations vulnerable to failure.  The hostile environment 
includes aqueous chemistry, saliva pH, thermal streses (hot coffe and ice-cream), 
and complex loading.  This requires high-strength materials for long life.  The 
hardnes and Young?s modulus of the replacement materials must be similar to 
natural dentition; any departure may adversely afect load transfer and lead to failure. 
And the materials should also be sufficiently tough to avoid fracture. In addition to 
these requirements are the needs for a good cementation of a crown replacement to 
the remaining tooth. Most important in the context of dentistry, there is the demand 
for aesthetics, requiring matching of natural teth in color, hue (tone), value 
(saturation) and chroma (lightnes), opacity, metamerism and fluorescence. Al this 
constitutes a significant materials chalenge.  
 Of primary interest to this study is the need to find new ways to fabricate 
dental crowns.  Currently, there is a trend toward al-ceramic systems, because of 
 
 
 
 
2 
aesthetics, bioinertnes and biocompatibility.  Traditional metal/aloy-based crowns 
are losing favor. But al-ceramic systems pose a problem. It is almost impossible to 
find a single ceramic that provides the esential combination of strength and 
aesthetics, so modern crowns are generaly fabricated by painting an aesthetic 
porcelain veneer onto a hard, stif and tough ceramic alumina or zirconia base. This 
approach has several drawbacks.  It is time- and labor-intensive, and demands a 
highly skiled technician.  Erors in sintering times and temperatures can introduce 
flaws (pores, cracks) that lead to premature failures. There is an interesting materials 
chalenge here, to produce an alternative economic procesing route. 
 Our approach is to introduce a diferent method of making al-ceramic 
crowns, whereby a porcelain veneer and ceramic core can be fabricated independently 
and subsequently fused with a strong join.  The prospect is of cheaper, beter and 
more durable crowns. We wil investigate this approach in this study using glas as a 
bonding agent. Glas joining has a solid history in several technologies, notably in 
glas seals and glazes, electronic packaging, fuels cel technology, etc.  Adjacent 
veneer and core ceramic layers could be fabricated independently via solid fre-form 
techniques or CAD/CAM methods, and sintered acording to material specifications. 
Then, the two layers could be fused using a thermaly sintered glas join, with melting 
and weting below the softening temperature of the veneer. 
 1.2 The Challenge 
 
 To our knowledge, the performance of glas joins has not been systematicaly 
evaluated in systems that relate to dental crown fabrication.  There are many 
 
 
 
 
3 
chalenges to putting such an approach into practice. (i) One must be able to prepare 
veneer and core ceramics routinely and independently. This is the realm of traditional 
ceramic science, and is wel documented. (i) Then one must demonstrate that these 
layers can be fused with an appropriate glas. The glas must wet the two ceramic 
components.  The glas join must have good thermal fit to both the veneer and the 
core, so as to avoid deleterious residual streses from thermal expansion mismatch 
and must be bioinert or biocompatible. (ii) One must understand the chemical nature 
of the join, and characterize it. Generaly, the join is expected to form an atomicaly 
mixed interdifusion layer. The degre of interdifusion, interdifusion chemistry and 
thermomechanical properties are al expected to influence the overal mechanical 
performance of the joins. (iv) Finaly, one must then test the strength of the join, to 
demonstrate its resilience under comparative oral operational conditions, 
incorporating a hostile environment and fatigue loading. Throughout the fabrication 
proces, it is important to remain aware of clinical relevance, to ensure that the 
technology is transferable to dental practice.  
 
 1.3 Outline of Thesis  
 
 The thesis wil folow the following route. Chapter 2 wil outline how suitable 
glases are chosen to match veneering porcelains and ceramic cores.  For this, 
matching of thermal expansion coeficients is vital, to avoid residual streses in the 
finished product.  The chapter wil describe how commercial glases were tested to 
met this need, and then rejected because of incompatibility concerns.  It wil then 
 
 
 
 
4 
describe how new glas compositions were designed and fabricated to overcome 
these incompatibility isues. Sol gel procesing wil form the underlying procesing 
route, with the prospective advantage of low-temperature manufacture. Various tests 
for asesing the glas properties wil then be described, including degre of weting, 
crazing and pore formation.  
 Chapter 3 wil deal with the methodology for joining veneer and core layers 
by glas fusion. Glases with a range of compositions wil be investigated, and tested 
to se how wel they bond the adjacent layers.  Diferent firing temperatures and 
times for glases of diferent compositions wil be explored, to determine optimum 
bonding conditions.  Simple mechanical screning tests for deciding whether any 
given layer combination is suitable for more detailed study wil be described. Optical 
microscopy of veneer/glas/core sections wil be used to determine the integrity of the 
fused interfaces.  Microprobe analysis for determining the nature of interdifusion 
layers at the interface boundaries wil be presented.  
 In Chapter 4, the mechanical integrity of the glas joins wil be examined in 
greater depth. For this, a Vickers indentation probe method developed at NIST wil 
be employed. This method introduces smal controlled cracks at prescribed locations 
in the veneer/glas/core sections, close to and away from the interfaces of interest. 
This is a simple but powerful way to examine local fracture properties of complex 
systems.  The indentations are aligned so that a lead crack is made to approach the 
glas join interface?if the crack arests at the boundary, or penetrates it, without 
deflecting along the interface, then the system is deemed to be adequately strong for 
the specific application in mind.  A quantitative analysis of this mechanical testing 
 
 
 
 
5 
procedure wil be presented.  
 Finaly, in Chapter 5, future directions wil be discussed. 
 
 
 
 
6 
 
 Chapter 2:  Design & Procesing of Tailored Glasses 
 
 
 2.1 Introduction 
 
 Dental crowns are used to repair badly damaged teth. They are traditionaly 
fabricated by fusing hand-layered porcelain veneer onto metal or, more recently, 
ceramic cores. In current restorations, even when the core is shaped by CAD/CAM 
(or solid freform fabrication proceses), the veneer is applied in sequential layers. 
Many problems are asociated with this method of veneer application, including high 
expense, labor-intensive multiple layers, which may degrade the long-term 
performance.  Fabricating veneers and core separately?e.g. computer-aided design 
and manufacturing, including robocasting by [1]?would appear to have some 
advantages. Some workers [2-5] have explored the possibility of joining with filed 
epoxy resins.  This is an easy solution, but the interfaces tend to be weak and 
susceptible to degradation over time.  
 In this study our aim is to explore the possibility of using more durable joins, 
specificaly glas interlayers.   This method requires higher temperatures, so more 
atention must be paid to residual stres isues.  However, it offers the prospect of 
much stronger and more chemicaly inert structures. There is a wide literature on the 
use of glases for joining, but dentistry imposes some stringent constraints.   For 
instance, it is important that the glas be matched in CTE (coeficient of thermal 
expansion) to that of the veneer and core, to minimize spurious streses. Also, there 
 
 
 
 
7 
is a limit to the operating temperatures for joining, to ensure that the integrity of the 
veneering porcelain is not compromised.  Additionaly, there are isues of toxicity. 
And finaly, the glas must resist large numbers of stres cycles under adverse 
chemical conditions (fatigue).  Thus lead-containing glases, while widely used as 
glas seals, are not a viable possibility for any form of human prostheses.  Al this 
means paying a great deal of atention to glas compositions and procesing 
methodologies. 
 In this chapter we wil lay out the procedure for selecting, designing and 
fabricating joining glases for porcelain veneers to ceramic cores, specificaly to 
dental aluminas and zirconias.  We wil provide a background outlining fabrication 
procedures, both general and specific.  Then we wil describe our experimental 
approach to the problem, with a detailed description of procesing in relation to the 
particular veneer/core combination under consideration.  Finaly, we wil describe 
various techniques used to characterize the materials produced, in preparation for 
layer joining in the next chapter. 
 
 2.2 Background 
 
 As indicated, layered structures can be used to addres the competing needs of 
systems like dental crown restorations where the exterior needs to be aesthetic and the 
interior needs to be strong and fatigue resistant.  Dental crowns typicaly have an 
aesthetic porcelain vener layered onto a strong, fatigue-resistant ceramic or metalic 
core [6], before being cemented onto the supporting tooth dentin structure layer, as 
 
 
 
 
8 
shown in Figs. 2.1 and 2.2.  The veneer-layering proces is labor intensive, time 
consuming and may not optimize the long-term performance properties of the veneer.  
If the core and veneer layers were independently fabricated and then joined, 
their individual performance could be optimized. Glas joining methods are widely 
used in other areas including lighting [7,8], fuel cels [9,10] and hermetic seals for 
electronic packaging
 
[11,12].  In principle, dense, CTE matched veneer/core joins 
could be fired at temperatures far enough below the softening temperature, T
s
, of the 
veneer without degradation.  Initial work in our laboratory has shown that this 
approach is feasible for ceramics used in dental restorations [13-16]. 
 Traditional glas procesing makes use of glas compositions to match 
requisite properties such as softening temperature, thermal expansion, electronic 
properties, opacity, etc. Examples of glas sealant use include hermetic sealing [7], 
electronic packaging [17] and solid oxide fuel cels [9,18].  The most widely used 
method is by mixing oxide powders with melting, fining, annealing and quenching 
[19,20].  Other methods include sol-gel and wet chemical procesing, with the 
advantage of lower-temperature manufacture (avoiding the need for platinum 
crucibles), greater control over material homogeneity and smal-batch  procesing 
[21-23]. This last approach has large appeal in the context of dental material joining, 
and so we adopt it here.  
 
 2.3 Materials Selection 
 
The first step was to determine which glas compositions we need to use.  We 
 
 
 
 
9 
need to match glas composition in terms of important properties, most importantly 
CTE and T
g 
 (glas transition temperature).   Our glases were formulated using 
property-dependent additivity factors developed by several early workers to tailor 
CTE [24-29], viscosity [30-34] and T
g
 (glas transition temperature) [35,36]. 
Compositions were then made by sol-gel procesing and glas-melting techniques, 
and tested for CTE and T
g
 using dilatometry and thermomechanical analysis (TMA), 
and then used to join layers of common veneer/ core combinations. 
 
 2.3.1 Choice of core and vener materials 
 For alumina-based core-veneer bilayers, experiments were caried out on core 
material from two sources:  one, an alumina used as substrates in the electronic 
industry (AD90, Coors, Golden, CO); and two, a dental alumina from Nobel BioCare 
(Procera, Nobel BioCare AB, Stockholm, Sweden).  The first of these was chosen 
because it was available cheaply in quantity, the second for its dental relevance. For 
zirconia-based bilayers, Cyrtina Dental zirconia, (Biozyram, Cyrtina Dental, 
Netherlands) was selected.  Materials were obtained as blocks, cut to size 
approximately 10 m square and 1?3 m thick. Blocks were ground to 0.2?0.7 m 
thicknes, and surfaces were polished flat and paralel to 1 ?m finish using diamond 
paste. Zirconia samples were finished with colloidal silica suspension to remove any 
strain-induced transformations.  
 Body porcelain vener powders were procured so as to match the CTE for 
the core materials. 
(a)
(b)
Figure 2.1 Dental crowns; (a) Veneered zirconia crown, with cement application, 
prior to bonding to supporting structure (Courtesy P. Coehlo). (b) Failed all-ceramic 
(Alumina) crown, with veneer, core and tooth structure clearly visible (Courtesy 
Kenneth Malament).
10
CORE
VENEER
1
m
m
v
e
n
e
e
r
/
0
.
5
m
m
c
o
r
e
Figure 2.2  Traditional methods of making dental ceramic restorations (courtesy G. Zhang, E.D.
Rekow
,
 V.P. Thompson and Y.
W
ang)
11
 
 
 
 
12 
 For alumina we used matching NobelRondo? Pres alumina porcelain 
(Nobel BioCare AB, Stockholm, Sweden), and for zirconia Sagkura Interaction 
porcelain (Elephant Dental, The Netherlands).  To make the porcelain, water-based 
slurries of the porcelain were poured into a mold, vibration was used to setle the 
porcelain and exces water was blotted off in sucesive layers. Then, the porcelain 
was dried and fired at the manufactures? recommended temperatures, as outlined in 
Fig. 2.3. Porcelain surfaces were then also ground and polished flat and paralel to 1 
?m finish using diamond paste.  
 Core and veneer properties relevant to our glas composition and properties 
include CTE of both core and veneer materials, and T
g
 and T
s
 (transition and 
softening temperature of the veneer materials).  T
s
 is the temperature at which a 
veneer can deform under its own weight. This would cause slumping and shape loss 
of the veneer, which must be avoided in dental restorations. Table 2.1 lists properties 
of veneer and core materials. 
 
 2.3.2 Choice of glass compositions 
 Initial investigations were performed on commercialy-available lead-fre 
industrial glas powders obtained commercialy (Schott Electronic Packaging, 
Landshut, Germany and Fero, Cleveland, OH).  The 9 glases investigated 
comprised glas-to-metal sealants and electronic and pharmaceutical packaging 
glases, listed in Tables 2.2 and 2.3.  Their compositions varied from alumino-
borosilicates (high working temperature) and bismuth-based (low working 
temperatures, T
w
). Unfortunately, the industrial glases proved to be inadequate for a 
 
 
 
 
13 
number of reasons that wil be discussed later. The need for developing laboratory-
made glases tailored to our needs thus became imperative.  
First, the glas properties critical to our dental sealant application were 
identified. As mentioned previously, our aim was to produce a glasy chemical bond 
for solid freform fabricated, or robocast separate veneer and core layers.  
Acordingly, our approach was to fabricate tailored in-house glas 
compositions to provide CTE matching to our selected alumina- and zirconia-based 
bilayer systems.  We used base glas compositions for which CTE properties have 
been determined from the glas seal literature [7], along with rule of mixture formulae 
to adjust individual glas components.  These rules involve a empirical linear 
additivity relations in terms of concentrations and mol% ?factors?, from which CTE 
can be calculated. 
The results of such calculations for specific glas compositions, are shown in 
Table 2.4 for factors from Winkelmann and Schott [24-26], and in Table 2.5 using 
factors from Appen [27-29]. In addition, empirical tabulations from Fluegel [35,36] 
were used to estimate viscosities at working temperature (800
 
C), along with glas 
transition, softening and melting temperatures (T
g
, T
s
, and T
m
), and Young?s modulus. 
Table 2.6 shows the results for the compositions in Table 2.5.  
It wil be noted that the laboratory glases compositions listed in Tables 2.4 to 
2.6 contain no toxic elements like lead. While the present study is meant simply to 
establish feasibility of the methodology, toxicity is stil of interest.  
Firing time cycle (min)
T
emperature (C)
0
0
200
400
600
800
1000
105
15
T
emperature (C)
Vacuum
20
0
0
200
400
600
800
1000
105
15
Vacuum
20
a)
b)
Figure 2.3 Manufacturer recomended firing cycles for a) Nobel Rondo 
alumina-veneering porcelain, and b) Sakura Interaction zirconia veneering 
porcelain.
14
A
l
u
m
i
n
a
 
s
y
s
t
e
m
 
m
a
t
er
i
a
l
s
 
Na
m
e
 
F
unction
 
Supp
l
ie
r
 
CTE 20
-
500 
?
C,
 
x
 10 
-
6
/C
 
F
l
e
x
u
r
a
l
 
St
r
ength,
 
M
P
a
 
E
l
astic
 
M
odu
l
us
 
E, G
P
a
 
F
r
actu
r
e
 
T
oughness, 
M
P
a.
m
-
1/2
 
T
g
 
?C
 
T
s
/T
m
 
?C
 
P
r
oce
r
a
 
 
C
o
re
C
o
re
 
N
o
b
e
l
 
B
i
o
care
 
7.0
-
7.3
 
600
-
687
 
370
 
4.48
 
 
>2,000
 
Rondo
 
 
V
e
n
eer
 
N
o
b
e
l
 
B
i
o
care
 
7.20
 
120
 
~ 65
 
 
620
 
>890
 
 Z
i
rc
o
n
i
a
 
s
y
s
t
e
m
 
m
a
t
er
i
a
l
s
 
Na
m
e
 
F
unction
 
Supp
l
ie
r
 
CTE 20
-
5
00 
?
C,
 
x
 10 
-
6
/C
 
F
l
e
x
u
r
a
l
 
St
r
ength,
 
M
P
a
 
E
l
astic
 
M
odu
l
us
 
E, G
P
a
 
F
r
actu
r
e
 
T
oughness, 
M
P
a.
m
-
1/2
 
T
g
 
?C
 
T
s
/T
m
 
?C
 
B
io
z
y
r
a
m
 
C
o
re
 
C
y
r
t
i
n
a
 
De
n
t
a
l
 
10.5
 
1
,
000
 
205
 
8.5
 
 
>2,600
 
Sa
k
u
r
a
 
Inte
r
action
 
V
e
n
eer
 
E
l
ep
h
a
n
t
 
De
n
t
a
l
 
9.9
 
>50
 
~ 65
 
 
485
 
>890
 
 
 
T
able 2.1 
 Properties of veneer and core materials
NA NA
NA
NA
1.0
AD999
 
 
 
 
8.0
 
500
 
406
 
 
 
 
3.9 1.0
Coors 
15
Properties and compositions of Schott industrial lead-free aluminoborosilicate and 
bismuth glasses 
Glass 
No. 
Function Composition  
CTE  
x10
-6
 C
-1
 
T
g  C  C
 
T
w
log =4) 
Modulus  
E, GPa 
Poisson?s 
ratio 
 
8415 Illax?, amber tubing 
glass for 
pharmaceutical 
packaging 
75% SiO
2
, 
 8?12% B
2
O
3
, 
 up to 5% alkaline 
earths and Al
2
O
3
 , 
 
7.8 535 1050 74 0.21 
8421 Sealing glass for 
seals with NiFe45 
(DIN 17745) and 
compression seals 
 
No Information 
available 
9.7 525 1000 74 0.22 
8422 Sealing glass for 
seals with NiFe47 or 
49 (DIN 17745) and 
compression seals 
 
No Information 
available 
8.7 540 1010 76 0.21 
G018-
193 
Sinter glass ceramic 
for LTCC 
application 
10-50% SiO
2
 
10?50% B
2
O
3
  
1-10% Al
2
O
3
 ,  
10-50% CaF
2
 
1-10% ZnO 
1-10% ZrO
2
 
1-70% alkaline earths 
 
10 580 830 70-80 ~ 0.22 
G018-
249 
Low melting Lead-
free Solder Glass 
1-10% B
2
O
3
 
0.1-2% Al
2
O
3
 , 
1-10% SiO
2
 
10-50% ZnO 
>50% Bi
2
O
3 
 
10.11 365 380-500 70-80 ~ 0.22 
G018-
250 
 
Low melting Lead-
free Solder Glass for 
CTE range 7-8 
1-10% Li
2
O 
1-10% B
2
O
3
 
1-10% Al
2
O
3
 , 
1-10% SiO
2
 
1-10% ZnO 
>50% Bi
2
O
3 
 
7 380 350-540 70-80 ~ 0.22 
 
(
Table 2.2  
16
Properties and compositions of Ferro industrial lead-free bismuth glasses 
EG2914 Lead-free sealing 
glass 
Bi
2
O
3
  
ZnO 
B
2
O
3
 
SiO
2
 
R
2
O *  
 
10.25 430 465-525 70-80 ~ 0.22 
EG2964 Lead-free bonding 
and sealing glass 
Bi
2
O
3
  
ZnO 
B
2
O
3
 
 
8.45 480 520-560 70-80 ~ 0.22 
EG2998 
 
Lead-free lamp 
sealing glass 
Bi
2
O
3
  
ZnO 
B
2
O
3
 
 
9.80 405 440-500 70-80 ~ 0.22 
Table 2.3  
Glass 
No. 
Function Composition  
CTE  
x10
-6
 C
-1
 
T
g  C  C
 
T
w
log =4) 
Modulus  
E, GPa 
Poisson?s 
ratio 
(
* R = alkali metal ion
17
 
 
 
 
18 
 2.4 Preliminary Tests Using Industrial Glasses 
 
As indicated above, we originaly set out to investigate the possibility of 
efecting glas joins using the commercialy available glas powders listed in Tables 
2.2 and 2.3.  In these cases, we mixed the powders into a thick aqueous slurry and 
painted the mixture as layers up to 200 ?m thick onto the alumina and zirconia plates. 
The plates were then fired at manufacturer-recommended times and working 
temperatures (T
w
 in Tables 2.2 and 2.3). The fired plates were then examined in an 
optical microscope. 
In short, no combinations of firing temperature could produce quality coatings 
on for the industrial glases on any of the core materials.  Several problems were 
encountered, as depicted in Fig. 2.4. These included, (a) porosity, (b) poor weting, 
(c) discoloration and poor aesthetics and (d) crazing.  Consequently, we turned our 
atention to laboratory fabricated glases, where we had beter control of the 
properties. 
 
 2.5 Glass Preparation by Sol-Gel Procesing 
 
 2.5.1 Sol-gel wet mixing 
In order to prepare the large number of experimental compositions, smal 
batches of glas were made using sol-gel procesing.  We adopted the wet mixing 
methodology proposed by Yoldas et al. [37,38].  Silicon tetraethoxide (TEOS), 
 
 
 
 
19 
Si(OC
2
H
5
)
4
, an alkoxysilane, was combined with aqueous solutions of metal nitrates 
in acordance with composition calculations. As TEOS is imiscible with water, it 
was first disolved in 200 proof ethanol, and a smal amount of water added.  
Soluble nitrates were disolved in de-ionized water then in ethanol, which was 
maintained as the solvent medium for the first few hours of reaction.  The nitrate 
solution was then added to the TEOS solution, and the mixture stired on a hot-plate 
at 60-70 C, where the following reactions took place: 
 
Si-OR + H
2
O ? Si-OH + ROH  (hydrolysis)   (1) 
Si-OR + HO-Si ? Si-O-Si + ROH  (re-esterification)  (2) 
Si-OH + HO-Si ? Si-O-Si + H
2
O  (alcoholysis)   (3) 
 
Al thre reactions are concurrent throughout the sol, forming a silicate polymeric 
network.  The sol is then alowed to cool, seting to a translucent gel as water and 
exces alcohol evaporate. Modifier metal ions remain within the network, and as the 
gel dries, cause structural variations that modify properties of resultant polymers 
[21]. The efect of these variations extends to the glases obtained from the gels, and 
can even modify the glas?s high temperature properties, including sintering, 
crystalization, and viscosity [22,39].  The resulting gels were then applied in their 
wet state onto alumina disks, dried, and fired at diferent temperatures in a dental 
furnace (Vita Vacuumat 2500, Vita Zahnfabrik, Bad Sackingen, Germany).  This 
resulted in foaming of the gel due to water and organics trapped in the glas precursor 
at higher temperatures. 
 
Glas s Notat ion  
(CT E ? 10
-6
/? C)
 
SiO
2
 
mol % 
Al
2
O
3
 
mol % 
CaO  
mol % 
Na
2
O 
mol % 
BaO  
mol % 
B
2
O
3
 
mol % 
 
7.60 60.00 3.00 5.00 10.00 13.00 9.00 
8.00 60.00 3.00 8.00 10.00 12.00 7.00 
8.30 60.00 5.00 8.00 10.00 12.00 5.00 
8.56 60.00 2.73 8.18 11.82 11.82 5.45 
8.70 60.00 2.67 8.00 12.44 11.56 5.33 
8.71 60.00 3.36 10.08 12.28 7.56 6.72 
8.81 60.00 3.47 10.40 12.64 6.56 6.93 
8.90 60.00 3.57 10.71 13.02 5.56 7.14 
8.92 60.00 4.29 12.87 12.86 1.40 8.58 
9.02 60.00 3.24 9.70 12.66 10.38 4.04 
9.14 60.00 3.28 9.83 14.14 6.20 6.55 
9.24 60.00 3.13 9.40 14.64 6.56 6.27 
 
Compositions of Winkelmann and Schott Factor-formulated lab glasses for bonding 
electronic alumina substrates
Table 2.4
20
 
Glass Notation  
(CTE ?10
-6
/? C)
 
SiO
2
 
mol % 
Al
2
O
3
 
mol % 
CaO 
mol % 
Na
2
O 
mol % 
BaO 
mol % 
B
2
O
3
 
mol % 
 
6.00 60 3.32 9.20 5.54 2.22 19.72 
6.50 60 3.29 8.22 5.48 5.48 17.53 
7.00 60 3.03 8.37 6.41 6.05 16.14 
7.50 60 3.07 8.39 7.16 7.06 14.32 
8.00 60 2.87 8.61 8.61 6.51 13.40 
Glass Notation 
(CTE ?10
-6
/? C)
 
SiO
2
 
mol % 
Al
2
O
3
 
mol % 
CaO 
mol % 
Na
2
O 
mol % 
BaO 
mol % 
B
2
O
3
 
mol % 
 
9.80 60 3.36 10.09 12.27 7.56 6.72 
10.20 60 3.28 9.83 14.14 6.20 6.55 
10.40 60 3.13 9.40 14.64 6.56 6.27 
 
Compositions of Appen Factor-formulated lab-tailored alumina-compatible and 
zirconia-compatible glasses
Table 2.5 
Alumina system materials 
Zirconia system materials
21
Glass
 
N
ot
a
t
i
on
 
C
TE (b
y
 
Appen)
 
?
 1
0
-
6
 
o
C
-
1
 
log
1
0
 (
v
is
c
osi
t
y 
P
a
.
s) a
t 
8
0
0
 ?C
 
*
T
g
 , ?C
 
*
T
s
 , ?C
 
 
*
T
m
 , ?C
 
Y
ou
n
g
?
s 
Modu
l
us 
GP
a,
 
 
AP
0
6
0
0
 
6
.
00
 
3.9
 
5
8
6
 
607 613622
1
,
2
8
1
 
73
 
AP
0
6
5
0
 
6
.
50
 
4.6
 
5
9
5
 
1
,
2
6
5
 
74
 
AP
0
7
0
0
 
7
.
00
 
4.9
 
5
8
6
 
 
615
 
608 600 597
 
1
,
2
0
1
 
74
 
AP
0
7
5
0
 
7
.
50
 
4.8
 
5
8
5
 
1
,
1
7
1
 
77
 
Glass
 
N
ot
a
t
i
on
 
C
TE (b
y
 
Appen)
 
?
 1
0
-
6
 
o
C
-
1
 
log
1
0
 (
v
is
c
osi
t
y 
P
a
.
s) a
t 
7
8
0
 ?C
 
*
T
g
 , ?C
 
*
T
s
 , ?C
 
 
*
T
m
 , ?C
 
Y
ou
n
g
?
s 
Modu
l
us 
GP
a,
 
 
AP
0
9
8
0
 
9
.
80
 
4.8
 
5
7
3
 
1
,
1
7
1
 
78
 
AP
1
0
2
0
 
1
0
.
2
0
 
4.6
 
5
6
7
 
1
,
1
6
9
 
77
 
AL1
0
40
 
1
0
.
4
0
 
4.6
 
5
6
3
 
1
,
1
6
2
 
77
 
 
Calculated properties of selected join 
Appen glass compositions
T
g
 ?C (log 
h
 in Pa.s =12): glass transition 
T
; 
T
s
 ?C (log 
h
 in Pa.s =9): glass flow 
T
;  
T
m
 ?C (log 
h
 in Pa.s =1.5): glass 
melting 
T
.
A
l
u
m
i
n
a
 
s
y
s
t
e
m
 
m
a
t
er
i
a
l
s
 
Z
i
rc
o
n
i
a
 
s
y
s
t
e
m
 
m
a
t
er
i
a
l
s
T
able 2.6 
 
*
22
(a)
(c)
(b)
(d)
Figure 2.4   Lead-free commercially available glasses from Schott (Bad Sackingen, 
Germany and Ferro, USA).  (a) Schott glass No. 8415, CTE 7.8x10
-6
C
-1
 on electronic 
alumina (EA) sintered at 800 C, showing pitting; (b) Ferro EG2964 CTE 8.5x10
-6
C
-1 
on 
alumina, showing poor wetting; (c) Schott G01-8249 CTE 10.1x10
-6
C
-1 
on zirconia, 
showing discoloration;  (d) Schott G01-8421 CTE 9.7x10
-6
C
-1 
fired at 1000C on 
alumina, showing improved melting but crazing.  
23
 
 
 
 
24 
We subsequently decided that it would be best to heat the dried gel at a 
temperature high enough to remove these organics (but low enough to avoid melting, 
i.e. calcine them), then to grind the resultant powders and reapply them to the alumina 
disks. This alternative route is indicated in the flow diagram in Fig. 2.5. Melting tests 
were then conducted to determine appropriate firing temperatures.  Two problems 
nevertheles persisted with this alternative route.  First, in some of the mixtures, 
precipitation occurred when the TEOS and modifier salt solutions were combined. 
This meant that the glas was inhomogeneous on a molecular level, which could 
result in phase transformations during melting, and a highly unpredictable join 
behavior. The second problem was that some foaming of the powder stil occurred 
betwen 600 and 1000 C, leading to some porosity. This was taken as evidence that 
some residual organics must have remained even after calcining. 
Both these later problems were tackled concurrently.  Precipitation is 
common in alkoxysilane gels and has been documented elsewhere [37,38]. We first 
tackled this problem by using a more soluble barium salt, in the acetate form. The 
resultant gels were inconsistent in terms of precipitation.  We then mixed the 
solutions with a higher ethanol component. The slow evaporation of the ensuing sol 
resulted in a gel that was completely homogenous to the eye, and therefore aceptable 
for our needs.  
As to the second problem, calcining of the gels at 600 C for up to 15 hours 
resulted in powders that stil foamed during firing on alumina disks. We ran a series 
of tests in which the specimens were heated in intervals of 50 C betwen 650 C and 
1200 C. For this, the calcined powders were placed onto an alumina substrate before 
 
 
 
 
25 
heat treatment.  Specimens were examined after each temperature interval.  We 
observed some foaming in our melts up to about 850 C, and complete absence at 1200 
C. Separate thermogravimetric analysis (TGA) was caried out in a Shimadzu TGA 
50 at a heating rate of 10 C/ min to 1000 C, to check removal of organic volatiles. 
TGA profiles are shown in Fig. 2.6, for several gel compositions (Table 2.5). Weight 
loss is substantial up to 150 C, with subsequent gradual faloff, to a plateau at about 
850 C, i.e. at the point where foaming ceases.  
To reduce porosity (cf. Fig. 2.4a), we subsequently decided to run a second 
heating cycle up to 800 C. This succeded in removing most of the pores. However, 
some minor crazing persisted in some specimens (cf. Fig. 2.4d), and some specimens 
showed evidence of poor weting (cf. Fig. 2.4b).  Another problem that arose, 
especialy in those systems that used a higher content of ethanol to reduce 
precipitation was discoloration?in these cases, the calcined gels tended to turn to a 
dark grey or black powder (cf. Fig. 2.4c), presumably due to the entrapment of larger 
amounts of carbonized organics.  
 
 2.5.2 Melt and quench proces 
In an atempt to eliminate the problems outlined at the end of Sect. 2.5.1, we 
decided to augment the procesing route with an addition melt and quench stage. The 
calcined gel powder was melted and fined at 1200 C for 3 hr, before quenching into 
an ice/water bath.  This caused the glas melts to fragment into dense shards.  The 
fragments were then broken down by grinding into powder using a corundum mortar 
and pestle, and sieved to a maximum 38 ?m frit. This route is ilustrated in Fig. 2.7.
Figure 2.5
 
Sol-gel/wet mixing glass precursor preparation process diagram
 
Hot  plate, 60?C 
Stirrer  spee d 
4.5, 15 min. 
 
Sol-gel 
glass 
precursor 
granules 
Stirrer  4.5 speed, 48 hrs  
Air-dry , hood 
 
Hot  plate 60 ?C, Stirrer  4.5 
speed , 24 hrs  
Dry Gel  
 
Stirrer  spee d 
4.5, 15 min. 
NaNO
3
 
Ca(NO
3
)
2
*4H
2
O 
Al(NO
3
)
3
*9H
2
O 
Ba(NO
3
)
2
 
H
3
BO
3
 
H
2
O ETOH  
H
2
O 
TEOS 
Organic  Precursor Solution  
 
CAL CINING 
Furnace , ramp 
20 ?C/min. 
Hold at 600 ?C 
for 30 mi n. 
 
26
Figure 2.6   
TGA
 of gels to assess minimum temperature for complete removal of volatiles
0.002.004.006.008.00
10.0012.0014.0016.0018.00
0
100
200
300
400
500
600
700
800
900
1000
T
emp C
 
6.0 
TGA
 mg
6.5 
TGA
 mg
7.0 
TGA
 mg
7.5 
TGA
 mg
8.0 
TGA
 mg
27
 
 
 
 
28 
The frits were then mixed into aqueous slurries. At first, 0.5 wt% polyvinyl 
alcohol was tried as a binding agent, in order to enable smooth application of the 
slurry mixture onto the core ceramic substrates with a paint brush. However, even at 
the slowest drying rates available with our furnace, some discoloration occurred, from 
encapsulation of carbon particulates [40].  We then tried a solution of 50 vol% 
ethanol and water mixture. This was a litle more inconvenient, because it required 
painting on of several layers of the slurry. However, the end result was a significant 
improvement, with no detectable porosity, good weting, good aesthetics and no 
crazing. A typical example is shown in Fig. 2.8, for an Appen-formulated glas on a 
dental alumina substrate.  
 
 2.6 Material Characterization 
 
Characterization tests were run on the core, veneer and final glas 
compositions, (AP700 for alumina systems, and AP 1040 for zirconia systems) 
material components to measure pertinent properties.  These included 
thermomechanical analysis (TMA) primarily to confirm calculated CTE values. 
Additional dilatometry tests were run on just the glases, to obtain additional 
information on the glas working temperatures. 
After the melting and fining proces, some of the glas melts were poured into 
preheated graphite bar molds and then annealed 580 ?C for 1 hour. The glases were 
then alowed to cool to room temperature, emptied from the molds, and cut into plates 
7 m on each side and 1 m thick. Specimens of similar size were also prepared for 
 
 
 
 
29 
the dental alumina and zirconia core materials, as wel as for the corresponding 
porcelain veneers.  These were placed in the TMA furnace (IPC TM-650, 
Northbrook, IL) and heated to 500 C at a rate 25 C/min. Measured relative expansion 
is plotted as a function of temperature for each material in the alumina and zirconia 
systems in Figs. 2.9 and 2.10, respectively. Data are shown only within the range 100 
C to 500 C, because the specimens take some time to ?setle in? on initial heating 
(hence the relative displacement of the curves along the expansion axis).  The key 
observation is a common curve-fit slope (inclined lines) within 1x10
?6
 C
?1
, 
representing the CTE for the 3 materials in each plot.  This indicates good thermal 
matching in the critical temperature range with respect to mismatch streses.  
Additional tests were run using a dilatometer (Orton dilatometer DIL 2016 
STD, Westervile, OH) on just the glas materials. These were conducted on larger 
bar specimens, 25 long and 3 m square cross section.  These tests enabled us to 
expand the temperature range up to 625 C, thus alowing for determination of 
working temperatures.  For these tests, the heating rate was much slower, 3 C/min, 
eliminating much of the setling-in problem observed in the TMA. Figures 2.11 and 
2.12 contain results from these tests, again for the selected glases used in the alumina 
and zirconia systems.  Results are shown as the curves, with the lines representing 
calculated CTE values from Table 2.5.  These values compare favorably (within 
1x10
?6
 C
?1
) with the fited lines from the TMA data in Figs. 2.9 and 2.10.  
Also indicated in Figs. 2.11 and 2.12 are the glas transition temperatures T
g
 
and dilatometry softening temperature T
s
. These values also compare favorably with 
those calculated in Table 2.6.  
Fig. 2.7  Diagram of glass preparation from sol-gel precursors  
 
MELTING 
Furnace: 
Heat at 20 
?C/ min to 
1200 ?C. 
 
QUENCHING 
Ice/water 
equilib. mix, 
4?C: Pour 
liquid glass 
into mix. 
 
FINING 
Furnace: 
Hold at 
1200 ?C 
for 3 hrs 
 
Viscous
glass  
Liquid 
Lab-
tailored 
join 
glass 
 
GRINDING 
Mortar & 
Pestle, Ball 
mill: Particle 
size  36 m 
Glass 
shards 
Sol-gel 
glass 
precursor 
granules 
 
30
Figure 2.8 Crazing in Appen-formulated glass, AP650 (CTE = 
6.5 x 10
-6
 C
-1
) fired at 800 ?C for 10 mins on Procera alumina 
substrate.  No crazing, porosity is visible, and wetting is good.  
Matt finish is due to surface roughness.   
31
y = 7.33E-06x + 3.97E-03
y = 7.00E-06x - 1.43E-04
y = 6.54E-06x + 8.65E-04
0
0.0010.0020.0030.0040.0050.0060.0070.0080.009
0
100
200
300
400
500
600
T
emperature C
Procera AP0700 Rhondo Linear (Procera) Linear (AP0700) Linear (Rhondo)
32
Figure 2.9 Thermomechanical Analysis (TMA) of alumina system with  data of tailored glasses overlaid. The matching CTEs over the sintering temperature range are represented by the slopes of the graphs. Procera is the alumina core material, AP0700 is the glass join, and Rondo is the veneering porcelain.
 
y = 1.10E-05x + 1.50E-03
y = 1.19E-05x - 7.23E-04
y = 9.98E-06x - 1.34E-02
-0.015
-0.01
-0.005
0
0.005
0.01
0
100
200
300
400
500
600
T
emperature C
Cyrtina AP1040 Sakura Linear (Cyrtina) Linear (AP1040) Linear (Sakura)
33
Figure 2.10 Thermomechanical Analysis (TMA) of zirconia systems with data of tailored glasses overlaid. The matching CTEs over the sintering temperature range are represented by the slopes of the graphs. Cyrtina  is the zirconia core material, AP1040 is the glass join, and Sakura is the veneering porcelain.
Figure 2.1
1 
Dilatometry 
of 
AP700 
alumina 
systems 
joining 
glass. 
The 
red 
line 
represents 
a 
slope 
of 
y 
= 
mx, 
where 
m 
represents 
the 
coef
ficient 
of 
thermal 
expansion 
at 
the 
denoted 
temperature 
range 
 
and 
is 
calculated 
as 
7.00 
x 
10
-6
 
C
-1
 
for 
this 
glass. 
T
g
 
and 
T
s
 
(transition 
and 
softening 
temper
atures 
respectively) 
areare 
denoted 
 
by 
the 
intersection
 
of 
the 
green 
dotted 
lines 
representing 
change 
in 
slope, 
or 
transition 
of 
expansion 
behavior
. 
The 
blue 
shaded 
ellipse 
identifies 
the 
area 
where calculated and measured thermal expansion behaviors match.
34
Figure 2.12 
Dilatometry 
of 
AP1040 
zirconia 
systems 
joining 
glass. 
The 
red 
line 
represents 
a 
slope 
of 
y 
= 
mx, 
where 
m 
represents 
the 
coef
ficient 
of 
thermal 
expansion 
at 
the 
denoted 
temperature 
range 
 
and 
is 
calculated 
as 
10.40 
x 
1
0
-
6
 
C
-
1
 
f
o
r 
t
h
i
s 
g
l
a
s
s
. 
T
g
 
a
n
d 
T
s
 
(
t
r
a
n
s
i
t
i
o
n 
a
n
d 
s
o
f
t
e
n
i
n
g 
t
e
m
p
e
r
a
t
u
r
e
s 
r
e
s
p
e
c
t
i
v
e
l
y
) 
a
r
e 
d
e
n
o
t
e
d 
 
b
y 
t
h
e 
intersection 
of 
the 
green 
dotted 
lines 
representing 
change 
in 
slope, 
or 
transition 
of 
expansion 
behavior
. 
The 
blue 
shaded ellipse identifies the area where calculated and measured thermal expansion behaviors match.
35
 
 
 
 
36 
 2.7 Discusion 
 
We have set out to select glases for the subsequent bonding of porcelain 
veneers to dental alumina and zirconia substrates.  Detailed melting experiments of 
selected glas frits onto alumina and zirconia substrates have been caried out, to 
determine which compositions wil provide good weting and bonding without 
porosity or crazing, without discoloration, and with a sufficiently low melting 
temperature.  Various procesing routes were investigated, from sol gel to wet 
chemistry to melting and quenching. We discovered that obtaining suitable glases to 
match both veneer and core ceramics, yet provide good weting and bonding, was not 
straightforward. Several early failures, followed by experimental design modification 
finaly led us to glases that appeared to met the ned.  
Apart from the need to choose glases that wet wel and yet show no porosity 
or crazing, other criteria had to be met. These included biocompatibility, with zero 
toxicity.  Of particular concern was the need to avoid any lead (or bismuth) in the 
compositions.  Only a smal range of glases met al these conditions.  Our first 
atempts to use commercialy available glases failed because of our lack of control 
over their critical properties such as CTE matching.  This forced us to expand our 
scope to include glas compositions fabricated in-house. Many early atempts at such 
fabrication also failed, and much efort was expended before we were able to 
manufacture fully compatible glas joining compositions.  
The next stage was then to use these glases to join the veneers to the 
corresponding core ceramics. This forms the subject of the next chapter. 
 
 
 
 
37 
 
 Chapter 3: Fused Joins: Laminar Ceramic Structures 
and Glass-Ceramic Interface Chemistry 
  
 
 3.1 Introduction 
 
Once glas compositions that match specific porcelain/core combinations have 
been obtained, it is necesary to join the two layers into a bilayer.  It has been 
indicated that dental crowns are laminar structures, with generaly complex 
geometries.  Since we are concerned only with the joining proces in this thesis, 
experiments wil be conducted here on flat layers, always mindful that the techniques 
wil need to be applicable to more general geometrical shapes.  Traditionaly, 
porcelains are painted on to a core material layer by layer, a time consuming and 
expensive proces.  One alternative fabrication route is fre-form fabrication, in 
which each layer is shaped individualy (e.g. by robocasting) and then integrated with 
a glas join, as in Fig. 3.1.  Such techniques are currently being investigated by 
materials procesors, and hold the promise of new-generation crown fabrication [41-
43].  
The integrity of the glas-ceramic interface is crucial to succes of our 
approach. First and foremost, a strong bond is required, so that delamination failure 
does not occur before any other failure modes. In order to maintain this integrity, it is 
also required to avoid thermal expansion mismatch streses and to ensure high 
 
 
 
 
38 
resistance to degradation in an aqueous environment, particularly over time.  
This demands characterization of the interface chemistry. It wil be necesary 
to determine interdifusion characteristics either side of the nominal interface, with 
particular atention to ionic difusion and elucidation of new phases.  Such changes 
wil be evident as gradients in chemical composition.  
This chapter wil describe ways to efect glas joins betwen CTE-matched 
porcelains and alumina or zirconia cores. For this, it is necesary to heat the bilayer 
combination to an elevated temperature, sufficiently high to make the glas flow and 
wet each component, but not so high to cause deformation of the porcelain (typicaly 
950
o
C). At this stage, the test that adequate CTE matching has been achieved through 
the heating and cooling cycle wil simply be that the bilayer remains intact during 
preparation and subsequent handling. On the other hand, evaluations of the integrity 
of the resulting interfaces, using electron microscopy and microprobe methods, wil 
form an important part of the characterization. It wil be recaled that we have chosen 
materials to minimize residual streses from thermal expansion mismatch (Chapter 2). 
Confirmation of the absence of any such residual streses in our structures wil be 
defered to the next chapter.  
Methods employed to cary out characterization of the interface chemistry wil 
be outlined in this chapter.  This wil be done by first using conventional optical 
microscopy to determine geometrical elements, e.g. interfacial roughnes. Then SEM 
(scanning electron microscopy) and SEM microprobe analysis wil be used to map 
out the grain structures in the core and veneer layers, and the width and composition 
of the resulting interdifusion layer (IDL). Difusion kinetics wil be determined by 
 
 
 
 
39 
observing changes in the compositional maps over time. The coeficient of thermal 
expansion is also a crucial factor in these applications, but we wil proced on the 
asumption that we have achieved good matching betwen components (Chapter 2), 
and return later to confirm that the level of any residual streses from such mismatch 
is negligibly smal (Chapter 4). 
 
 3.2 Background 
 
 3.2.1  Glass as a seal 
Historicaly, glas has been used as a ceramic/glas seal in clay-based ceramic 
pottery, to negate the weaknes and porosity of the fired ceramic.  Clearly, this 
ancient technology was developed by trial and eror. Functionaly, glazing promoted 
waterproof capability and prevented bacterial infestation in clay vesels, while also 
alowing decorative artistic expresion to be preserved on the surface.  For this, we 
owe our ancestors a debt of gratitude, in the name of culture as wel as technology.  
More recently, a similar method has been developed to form bioactive glas 
coatings on medical ceramic devices, such as alumina hip replacements.  Further 
functional grading is required, however, to prevent leaching of alumina into the 
bioactive glas at the high temperature of glazing, as wel as to prevent thermal 
mismatch betwen the bioactive glas-ceramic and the alumina component.  Such 
mismatch has been mitigated by the addition of an intermediate glas layer to 
improve CTE match and further protect the bioactive glas from alumina leaching 
[44]. 
Figure 3.1  All-ceramic crown with join interlayer translated  to a flat model crown
40
 
 
 
 
41 
In other medical applications, such as dental crowns, the notion of a glasy 
interface has special appeal, because of its bioinertnes. Unlike metalic components, 
glas is not susceptible to electrochemical interaction with the oral environment. The 
isue of toxicity is also avoided, provided certain elements (especialy lead, and 
possibly also bismuth) are not present.  
 
 3.2.2  Traditional applications in glass/ceramic joining 
Although the use of glas to join layers is a new approach in dentistry, it is a 
method widely used in lighting [7], fuel cels [9] as wel as electronic packaging [17], 
which require hermetic seals. When glas is used to join materials together, it is often 
applied at viscosities betwen 10
3
?10
5
 Pa.s (kg?m
?1
?s
?1
), in order to promote flow and 
weting of the separate layers [36]. In silicate glases, these viscosities are atained at 
temperatures in the range 600?1200 ?C.  During the joining proces at these 
temperatures, interdifusion or atomic mixing ocurs betwen the layers.  While, 
some interdifusion is critical to weting and good bonding betwen layers, intuition 
and ceramic procesing science indicates that too much interdifusion is likely to be 
detrimental.  The amount of interdifusion that occurs depends on the relative 
component compositions, the temperature and time required for joining. 
Glas has long been used in industry to provide seals for metals [7].  As 
always, the secret to succes has ben matching of CTE betwen the glas and metal. 
This rich history in materials technology gives credence to our proposition that 
glases could be used as veneer/core joining agents in dental crown systems.  In 
modern technology, glas sealants are in use for their low electrical conductivity in 
 
 
 
 
42 
electronic packaging [17]. Glas joins are in common usage in solid oxide fuel cels 
(SOFC) [9,18,45,46], where resistance to thermal cycling is an isue.  
Glas joining is also a highly developed field in electronics packaging 
applications.  Some groups [17] have caried out similar research involving sealing 
glas for silicon carbide electronic packages.  Glases have been selected for low 
dielectric loss, high service temperature and high insulation, and then asesed for 
reactivity with silicon carbide.  Weting behavior, mas transport across interfaces, 
and formation of new phases have been investigated to create and maintain the best 
hermetic seal. Thermal expansion matching was found to be important, since a 5% 
diference in CTE could result in cohesive or adhesive failure, compromising the 
component. Sealed cross sections have been examined for chemical reaction and 
structural integrity using electron microscopy.  Reaction kinetics were linked to 
reactant transport across the interface, leading to an optimum glas and  interdifusion 
layer.  The end result has been a succesful means of interlayer joining in these 
systems, seting a precedent for our study. 
 3.2.3 Dental all-ceramic restorations 
As indicated, al-ceramic crowns are currently fabricated by painstaking 
sequential procesing, in which thin layers of porcelain are painted on step by step. 
This stepwise procedure is needed to avoid buildup of residual streses, as wel as to 
control aesthetic appearance.  We have repeatedly mentioned that such streses can 
lead to catastrophic failure of the component, most likely during the preparation itself 
but also subsequently in service, from buildup of damage within one or other layer. 
Most al-ceramic crowns use alumina or zirconia as a core material, for strong support 
 
 
 
 
43 
of an otherwise weak porcelain veneer. The tendency is more toward zirconia these 
days, because of its superior strength and lifetime, but alumina remains a useful 
material for baseline comparison [47-49]. 
To avoid the problem of residual stres, two methods of fabrication have been 
considered. One simply uses a single ceramic for the entire crown, thus avoiding the 
mismatch isue altogether.  Glas-ceramics have been used in the field for this 
purpose [50]. The problem is that those ceramics which provide pleasing aesthetics 
are invariably weak, and do not last aceptably long in the mouth. This is because the 
same enamel-like particulates within the glasy matrix that provide the tooth-like 
aesthetics also weaken the microstructure. The second method of fabrication of this 
kind is to join a porcelain veneer to a strong ceramic core, as advocated in this study, 
but using a polymeric adhesive that can be cured at room temperature. This approach 
has in fact ben investigated by others in our laboratory [2,3]. While avoiding CTE 
streses, adhesive joining is limited by a relatively weak veneer/core interface, which 
is susceptible to degradation over time in harsh oral conditions. The use of glas as a 
sealing agent as proposed here represents a compromise approach, with procesing 
temperatures low enough to keep residual streses low but high enough to produce a 
strong chemical bond.  The higher modulus of the glas also diminishes flexural 
modes in the veneer distortion during mastication, thereby reducing the possibility of 
porcelain fracture (Chapter 4) [3].  Of course, other isues now present themselves, 
such as the role of interdifusion gradients at the interface, and these wil be part of 
our investigation.  
 
 
 
 
 
44 
 3.3  Experimental Methods 
 
In al the joining applications described in 3.2 above, including the dental 
systems of interest here, the integrity of the bond is influenced by the properties of the 
interface difusion layer (IDL). These layers tend to be smal, of the order ?m, with 
large gradients in chemical composition. A feature that plays an important role in the 
integrity of the join is the existence of residual streses, so a great deal of atention is 
paid to minimization of thermal expansion mismatch, typicaly to les than 1x10
-6
 C
-1
. 
In the case of dental ceramics of interest here, there are virtualy no comparative data 
on the chemical properties of IDLs. This deficiency provides a principal motivation 
for the present chapter. In what follows, we wil outline how the joins are made, and 
then describe how the interfaces are characterized. 
 
 3.3.1  Preparation of bilayer specimens 
Flat alumina and zirconia layers were prepared as outlined in the previous 
chapter, along with their matching porcelains.  A thin layer ~ 50?100 ?m of the 
appropriate glas slurry was placed onto the surfaces of each of the core ceramics and 
their matching porcelains, ensuring coverage over the entire surfaces to be bonded. 
Biscuit firing was conducted for each specimen layer at a rate of 30 C/min to 700?
750 C for 5 min, as shown for alumina (Fig. 3.2a) and zirconia (Fig. 3.3a), to bind the 
glas particles onto the surface and begin weting. The specimens were then slowly 
cooled within the furnace and removed for examination.  
Firing time cycle (min)
0
0
200
400
600
800
1000
10 20 30 40 50 60 70 80 90
T
emperature (C)
0
0
200
400
600
800
1000
10 20 30 40 50 60 70 80 90
T
emperature (C)
Vacuum
Low Med High
(a) Biscuit cycle
(b) Sintering cycle
Figure 3.2 Firing cycles for alumina systems. a) Glass slurry was 
applied to both veneer and core layers, then biscuit-fired at 750 C 
(to promote adhesion of the glass to the substrates). b) Veneer and 
core blocks were then sintered at 830 C (to complete the join).  
45
Firing time cycle (min)
0
0
200
400
600
800
1000
10 20 30 40 50 60 70 80 90
T
emperature (C)
0
0
200
400
600
800
1000
10 20 30 40 50 60 70 80 90
T
emperature (C)
Vacuum
Low Med High
Biscuit cycle
Sintering cycle
Figure 3.3 Firing cycles for zirconia systems. a) Glass slurry was 
applied to both veneer and core layers, then biscuit-fired at 700 C (to 
promote adhesion of the glass to the substrates). b) Veneer and core 
blocks were then sintered at 800 C (to complete the join).
46
 
 
 
 
47 
The fired porcelain biscuits were then placed on top of their respective cores, 
alumina (Fig. 3.2b) and zirconia (Fig. 3.3b), and heated at a rate of 40 C/min under 
dead weight (4.4 g refractory ceramic) to 800?830 C. This was wel below the firing 
temperature of the veneer recommended by most manufacturers, and was chosen in 
order to avoid ?slumping? or loss of veneer shape.  The specimens were held for 
prescribed hold times of 15 min (low), 30 min (medium) and 45 min (high) to fuse 
the surfaces together.  These times were chosen to investigate appropriate levels of 
interdifusion.  
After cool-down, the sandwich specimens were sectioned perpendicular to the 
interface using a low-speed diamond saw blade, bonded with resin onto aluminum 
specimen holders. The surfaces were polished down to 1 ?m for optical microscopy, 
SEM and microprobe analysis. 
 
 3.3.2 Characterization of layer interfaces 
The fired bilayer sections were subjected to several tests, to establish the 
integrity of the joins.  The first set of these were simple tests of specimen survival 
from preparation.  To begin, the newly fabricated bilayers were subjected to 
interfacial shearing by hand, the so-caled ?finger test?.  Although crude, this test 
usefully eliminates the weakest interfaces in the population tail. The second such test 
of survival was a ?cutting test?, in which specimens were sectioned with a diamond 
saw, followed by surface grinding and polishing down to 1 ?m finish (Sect. 3.3.1). 
Those that broke during the cutting and polishing operation were also eliminated?a 
kind of proof test. 
 
 
 
 
48 
Those specimens that survived intact were then examined by various forms of 
microscopy. First, reflection optical microscopy was used. The interfaces could be 
readily resolved by this method.  The quality of the observed joins, uniformity of 
thicknes and lack of porosity, were taken as first indicators of a ?good? interface. 
Some specimens were also examined by scanning electron microscopy.  
Wavelength dispersive spectroscopy (JEOL JXA-8900 SuperProbe, Japan 
Electron Optics, Tokyo, Japan) was used to examine any difusion of ions betwen 
the bonding glas and the ceramic sandwich layers. Across the ceramic layers, a 1 
?m-wide electron beam was scanned across the join, with readings taken every 1.5?
2.8 ?m. Results were plotted as intensity versus location across the interface. 
 
 3.4 Results 
 
 3.4.1 Screning tests 
Tables 3.1 and 3.2 shows results of the finger and cutting screning tests for 
specimen integrity, for glases of diferent CTEs and for diferent firing temperatures. 
Table 3.1 contains data from the initial finger tests. This table lists the pas (P) and 
fail (F) results for electronic alumina (EA) and dental alumina (DA), and for dental 
zirconia (DZ). Some specimens failed during preparation before even geting to the 
finger test (FP). The pas rate for EA specimens was 5/13, for DA specimens was 
5/6, and for DZ specimens 3/3.  The pas rate using Winkelmann?Schott (WS) 
glases was only 5/12, and for Appen (A) glases was 7/9.  The survivors were then 
pased on to the sectional cutting and polishing phase. Results for pas and fail rates 
 
 
 
 
49 
during this phase are shown in Table 3.2. The pas rates for surviving EA specimens 
was 3/5, for DA was 5/5, and for DZ was 3/3. The pas rate using Winkelmann?
Schott (WS) glases was only 3/5, and for Appen (A) glases was 7/7. 
From these results, it was concluded that good bonds could be formed for DA 
and DZ interfaces only.  By contrast, EA specimens responded poorly, and were 
excluded from further consideration.  A glases gave beter pas rates in both the 
finger and cutting tests. With these results we chose as conditions for further study 
those bilayers with the lowest firing temperature and yet survived the screning tests. 
For the DA bilayers, this meant a glas CTE of 7.0x10
-6
 C
-1
 and a firing temperature 
of 830 C; for the DZ bilayers, a CTE of 10.4x10
-6
 C
-1
 and a firing temperature of 800 
C. 
 3.4.2 Optical microscopy 
The selected sectioned and polished specimens were then examined by 
microscopy for structural integrity. Interfaces for porcelain/glas/alumina structures 
in Fig. 3.4 and porcelain/glas/zirconia structures are shown in Fig. 3.5.  Typical 
width of the glas layer in these specimens was 20?70 ?m. In Fig. 3.4, the porcelain 
was Rondo and the alumina was Procera (Table 2.1). The glases used were Appen 
(a) AP600 with frit size < 100 ?m and (b) AP700 with frit size < 38 ?m (Table 2.6), 
applied in slurry of water/ethanol slip suspension in both cases. The porosity evident 
in the glas in Fig. 3.4a is due to the large grit size (Chapter 2).  In Fig. 3.5, the 
porcelain was Sakura and the zirconia was Cyrtina (Table 2.1). 
T
a
bl
e
3.
1
F
i
nge
r
t
e
s
t
,
t
o
de
t
e
r
m
i
ne
?
pa
s
s
?
(
P
)
or
?
f
a
i
l
?
(
F
)
f
o
r
bonde
d
a
l
um
i
na
-
a
nd
z
i
r
c
oni
a
-
ba
s
e
d
s
pe
c
i
m
e
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52
(a)
(b)
50 ?m
Figure 3.4  Showing joins in glass-bonded Rondo porcelain and 
Procera alumina bi-layer.  Glass in (a) AP600, in (b) AP700. 
White horizontal lines are used to highlight the veneer/glass junc-
tion.  Some porosity is evident in (a).    
   
52
(a)
(b)
Figure 3.5  Showing joins in glass-bonded Sakura porcelain and 
Cyrtina zirconia bi-layer.  Glass in (a) AP600, in (b) AP700. 
White horizontal lines are used to highlight the veneer/glass junc-
tion.  Some porosity is evident in (a).      
    
                                               
50 ?m
53
 
 
 
 
54 
The glases used were (a) AP1020 with frit size < 38 ? and (b) AP1040 with 
frit size < 38 ?m) (Table 2.6), applied with PVA binder in water solution in first case 
and water/ethanol slip suspension in second. In this case, the porosity evident in Fig. 
3.5a was due to the PVA binder (Chapter 2). Hence we chose smaler grit sizes and 
avoided PVA binders in our glas slurry preparations. 
Note that the specimens in Figs. 3.4 and 3.5 al pased the screning tests for 
integrity indicated above.  Clearly, such tests are insufficient to determine ultimate 
interface integrity. We leave this to Ch. 4.  
 
 3.4.3 Electron microprobe analysis 
Electron microprobe analyses with backscatered electron microscopy were 
conducted on the samples to determine compositional gradients across the glas-
bonded interfaces. Results are shown in Figs. 3.6 to 3.7. 
In these figures the veneer, glas and core zones are delineated, and the 
various ionic species indicated in the legend.  Specificaly, these plots enable the 
quantification of interdifusion layers (IDLs) at the veneer/glas (V/G) and glas/core 
(G/C) interfaces.  Widths of IDLs are shown as the vertical colored bands, as the 
distance over which ion concentrations have undergone significant change, although 
this distance is dificult to quantify.  Of special interest here is the difusion of the 
network-modifier ions, i.e. Na
+
, K
+
, Ca
+
 and Ba
+
, across the boundaries. 
The first set of plots shows gradients across interfaces in alumina-core 
systems fired at 830 C (Fig. 3.6) and zirconia-core systems fired at 800 C (Fig. 3.7), 
for (a) low (15 min), (b) intermediate (30 min) and (c) high (45 min) anneal cycles, as 
 
 
 
 
55 
depicted earlier in Figs. 3.2 and 3.3. Note the considerable scater in intensity for 
some ions within the glas bond, indicating some fluctuation in composition. For the 
alumina-based system in Fig. 3.6, the IDL extends over 10?15 ?m, increasing with 
anneal time. The levels of Na
+
, K
+
 and Ca
+
 and Ba
+
 show signs of difusion across 
the V/G interface, Ca
+
 and Ba
+
 from the glas into the veneer and K
+
 and Na
+
 in the 
opposite direction.  The curve for K
+
 is most interesting. As al the glas joins were 
formulated with no K
+
 in the glas, the presence of this ion indicates interdifusion 
from the porcelain. For this particular ion, the difusion has extended wel beyond the 
shaded V/G band in Fig. 3.6, and is shown more clearly in Figs 3.8a.  There is not 
much evidence of a large difusion of ions across the G/C boundary into the core, 
consistent with the higher density and crystalinity of the core material. On the other 
hand, there is indication of some difusion of Al
++
 into the glas from the core in the 
more highly annealed system in Fig. 3.6c.  Ba
+
 and Ca
+
 ions vary throughout the 
IDL width, with anneal time, as can be sen in Figs 3.8b and 3.8c for alumina. A plot 
of Si at the V/G interface and Al at the G/C interface mapping of major components 
at the veneer and core interfaces, highlights core interdifusion in Fig 3.8d  
The behavior of the zirconia-based system in Fig. 3.7 is similar in the Na
+
, K
+
 
and Ca
+
 and Ba
+
 difusion, but there are some minor diferences too. The IDL at 
the V/G interface appears to extend over a similar distance, i.e. 10?15 ?m, again 
increasing with anneal time.  However, the behavior at the G/C interface shows 
higher difusivity of the Zr
++
 (also detailed in Fig 3.9d) and Y
++
 ions than their 
Al
++
 counterpoints, commensurate with normal difusion behavior.  
As indicated above, the difusion of network modifier ions, Na
+
, K
+
 and Ca
+
 
 
 
 
 
56 
and Ba
+
, is of special interest.  Detailed intercomparison of difusion profiles for 
these ions is dificult to ascertain from the plots in Figs. 3.6 and 3.7, because of the 
compresed intensity scales, but plots of  K
+
, (Fig. 3.9a), Ca
+
 (Fig. 3.9b) and (Ba
+
 
Fig. 3.9c) show similar behavior for as for alumina systems, particularly in the case 
of K
+
. 
It is particularly dificult to compare numbers for the diferent anneal cycles in 
(a), (b) and (c) of those figures.  Therefore, individual plots for each of these ionic 
species is given with expanded scales in Figs. 3.8 and 3.9, comparing low, 
intermediate and high anneal cycles on each plot.  Composition fluctuations in the 
glas are magnified in these plots and the IDL zone is generaly clearer. Systematic 
variation betwen the diferent anneal cycles is not strongly apparent, especialy 
given the data scater, suggesting that longer anneals above some optimal level are 
only marginaly beneficial. Basicaly, an anneal that thoroughly melts the glas frits 
is enough to produce a strong bond.  
 
 3.5 Discusion and Sumary 
 
In this chapter the aim has been to bond porcelain to either alumina or zirconia 
using glas seals. The methodology required to do this has been laid out. We have 
described an optimal procedure for selecting the right materials and fusing conditions 
to produce joins of apparent structural integrity. A principal stipulation was that the 
various components, veneer, glas and core, should have close CTE matching. 
0
102030405060708090
100
380
400
420
440
460
480
500
520
      B        Ba       N
a 
      Ca       Al       K        Si 
57
Figure 
3.6a 
 
Microprobe 
scans 
of 
ion 
concentration
 
across 
glass-bond 
interface 
for 
alumina
-
core 
system, 
for 
low 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers. 
 
V
eneer 
(V), 
glass (G) and core (C) zones indicated.   
Alumina system - low interfacial diffusio
n
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
0
102030405060708090
100
380
400
420
440
460
480
500
520
      B        Ba       N
a 
      Ca       Al       K
  
      Si 
58
Figure 
3.6b 
 
Microprobe 
scans 
of 
ion 
concentration
 
across 
glass-bond 
interface 
for 
alumina
-
core 
system, 
for 
intermediate 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers.  
V
eneer (V), glass (G) and core (C) zones indicated.
   
Alumina system - intermediate interfacial diffusio
n
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
0
102030405060708090
100
340
360
380
400
420
440
460
480
      B        Ba       Na       Ca       Al       K
  
      Si 
Figure 
3.6c 
 
Microprobe 
scans 
of 
ion 
concentration
 
across 
glass-bond 
interface 
for 
alumina
-
core 
system, 
for 
high 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers. 
 
V
eneer 
(V), 
glass (G) and core (C) zones indicated.   
Alumina system - high int
erfacial diffusio
n
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
59
60
Figure 
3.7a 
 
Microprobe 
scans 
of 
ion 
concentration
 
across 
glass-bond 
interface 
for 
zirconia
-
core 
system, 
for 
low 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers. 
 
V
eneer 
(V), 
glass (G) and core (C) zones indicated.   
Zirconia system - low interfacial diffusion
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
N
a
Figure 
3.7b 
 
Microprobe 
scans 
of 
ion 
concentration
 
across 
glass-bond 
interface 
for 
zirconia
-
core 
system, 
for 
intermediate 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers.  
V
eneer (V), glass (G) and core (C) zones indicated.
   
Zirconia system - intermediate interfacial diffusio
n
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
61
Figure 
3.7c 
 
Microprobe 
scans 
of 
ion 
concentratio
n 
across 
glass-bond 
interface 
for 
zirconia 
core 
system, 
for 
high 
anneal 
cycle. 
 
Colored 
bands 
indicate 
interdif
fusion 
layers. 
 
V
eneer 
(V), 
glass (G) and core (C) zones indicated.   
Zirconia system - high interfacial diffusio
n
Di
s
t
ance ac
r
oss join (?
m)
Ion intensity (%)
V
G
C
62
V G
CG
(a)
(b)
Potassium
63
Figure 3.8a  Potassium diffusion in alumina systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue    )  anneal times. The vertical axis is vertical intensity, 
and the horizontal axis is distance ( ?m).
 
 
V G
C
G
(a)
(b)
64
Calcium
Figure 3.8b  Calcium diffusion in alumina systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue   )  anneal times. The vertical axis is relative intensity, 
and the horizontal axis is distance (?m).
G
G
C
65
V G
C
G
(a)
(b)
Barium
Figure 3.8c  Barium diffusion in alumina systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue    )  anneal times. The vertical axis is relative intensity, 
and the horizontal axis is distance (?m).
66 
V G
CG
(a)
(b)
Aluminum
Silicon
Figure 3.8d  Diffusion in alumina systems, a) silicon at 
veneer/glass (V/G) join interface, b) aluminum at 
glass/core (G/C) interface. For low (yellow    ), 
i n t e r- m e d i a t e ( r e d    ) a n d h i g h ( b l u e   ) a n n e a l 
times.The vertical axis is intensity, and the horizontal 
access is distance (mm).
V G
CG
(a)
(b)
Potassium
67
Figure 3.9a  Potassium diffusion in zirconia systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue    )  anneal times. The vertical axis is relative intensity, 
and the horizontal axis is distance (?m).
V G
C
G
(a)
(b)
68
Calcium
Figure 3.9b  Calcium diffusion in zirconia systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue    )  anneal times. The vertical axis is relative intensity, 
and the horizontal axis is distance (?m).
69
V G
CG
(a)
(b)
Barium
Figure 3.9c  Barium diffusion in zirconia systems, a) 
veneer/glass (V/G) join interface, b) glass/core (G/C) 
interface  For low (yellow    ), intermediate (red   )  and high  
(blue    )  anneal times. The vertical axis is relative intensity, 
and the horizontal axis is distance (?m).
G
70
V G
CG
(a)
(b)
Zirconium
Silicon
Figure 3.9d  Diffusion in zirconia systems, a) silicon at 
veneer/glass (V/G) join interface, b) zirconium at glass/core 
(G/C) interface  For low (yellow    ), intermediate (red   )  
and high  (blue    )  anneal times. The vertical axis is relative 
intensity, and the horizontal axis is distance (?m).
 
 
 
 
71 
The higher pasrate for A glases indicated that Appen-formulated 
compositions were a beter fit with respect to CTE matching. This is due to the 
improved capacity of the Appen approach to beter predict the coeficient of 
expansion of our glases. Detailed atention to glas frit preparation and subsequent 
join firing conditions was then found to be necesary. As indicated in Chapter 2, it is 
crucial to choose the right conditions, to avoid non-weting, crazing and porosity in 
the glas join itself. Finer frits produced the best joins. To determine optimum fusion 
conditions, specimens were produced over a range of firing temperatures. Those that 
survived the firing were then pased on to simple, preliminary screning for 
mechanical integrity, by means of a ?finger? bend test.  The survivors were then 
subjected in turn to a ?cutting? scren test, in which specimens were cut into shape 
using a diamond saw. 
Optical microscopy was used to examine section of the intact fused 
interlayers, so as to help select the conditions for the best joins, i.e. uniformity of 
interlayer thicknes and fre of porosity or cracking. Those joins that showed clean 
interfaces and also survived the mechanical screning test were then selected for 
further examination. 
Electron microprobe analysis was then used to examine difusion of network 
modifier species in the glas join, quantified by measurement of an interdifusion 
layer (IDL). Extensive testing was caried out on this aspect of the work. For this 
part of the work, 3 firing (anneal) times at the optimal firing temperatures were 
examined. Basicaly, we found that Na
+
, K
+
 and Ca
+
 and Ba
+
 ions moved across the 
veneer/glas interface, at diferent rates, with K+, the most mobile, moving furthest 
 
 
 
 
72 
from veneer to glas. This was expected as monovalent cations have a greater 
difusivity than divalent modifiers.  Some Al
++
 and Zr
++
 ions difused from the 
core ceramics into the glas joins.  At low anneal times, the IDLs of both V/G and 
G/C were narow, 10?15 ?m for both alumina and zirconia systems, extending 
somewhat in range for larger anneal times. These were al indicators of wel-bonded 
interfaces, and demonstrate the feasibility of applying the technology to the formation 
of dental crown systems.   
 Having succesfully demonstrated capacity to make these joins, it remains 
now to evaluate mechanical integrity more closely. This forms the subject of the next 
chapter. 
 
 
 
 
 
 
73 
 Chapter 4:  Mechanical Evaluation 
 
 
 4.1 Introduction 
 
Mechanical properties are important in our application. As indicated earlier, 
our goal is to join two britle layers by fusion with a chemical bond.  In such a 
system, both veneer and core layers are vulnerable to fracture, most generaly to 
cracks that traverse the layer thicknes??transverse? cracks.  More importantly, in 
the context of the present work in which the bond is acomplished by glas fusion, the 
interlayer betwen the two layers is itself vulnerable. Acordingly, we need to ensure 
that the interface has sufficient strength or toughnes to survive the stringent 
conditions under which dental crowns are expected to operate.  Can the glas bond 
arest cracks in the veneer layer from penetrating into the core, or vice-versa? Above 
al, can the glas interlayer prevent cracks from delaminating the interface? The last 
of these isues is most key to aceptable performance.  
Acordingly, our aim is to conduct simple tests on fabricated layer structures 
to evaluate the strength of the interfacial bond that joins the veneer and core ceramic 
layers. For this, we use the indentation microprobe method with a Vickers diamond.  
This approach is simple but powerful, and is used extensively in materials research to 
characterize fracture and deformation properties, measure residual streses, and, most 
important, evaluate the tendency for cracks to penetrate interfaces or delaminate them 
[51-54].  
  
 
 
 
 
74 
 4.2 Background 
 
As indicated above, layer systems used in al-ceramic crowns are susceptible 
to fracture, particularly from cracks traversing the layer thickneses.  Many studies 
have been made of such crack systems, particularly by the NIST group [51-55]. At 
isue here is how a fused interface betwen the veneer and core layers interacts with 
such cracks. In particular, does the fusion layer impede the progres of these cracks, 
or deflect them along the interface? If the second of these scenarios is true, then the 
strength of the interface is suspect. 
A schematic indicating some of the more deleterious fracture modes is shown 
in Fig. 4.1. This system consists of an aesthetic veneer (porcelain) layer joined to a 
core (alumina or zirconia) support layer, with the whole cemented onto a soft and 
compliant substrate representing tooth dentin.  The veneer/core layer system is 
subject to concentrated loading at the top surface, simulating an occlusal contact. The 
concentrated ?Hertzian? streses can initiate a variety of cracks within the veneer, 
depending on the specific loading conditions [56,57].  These include two kinds of 
cone cracks initiated within the Hertzian elastic contact zone: elasticaly generated 
outer (O) cracks [58] and hydraulicaly pumped inner (I) cracks [59,60]. Such cracks 
are generaly axisymmetrical with the geometry of truncated cones, and penetrate 
downward and outward into the veneer subsurface.  
Substrate
Veneer ceramic
Occlusal contact
Core ceramic
O
IM
R
Join
Figure 4.1 Schematic showing various cracks that can form in ceramic 
veneer/core layers joined by glass, and cemented onto a compliant support 
layer.  At top surface, elastic contact can generate outer (O) and inner (I) 
cone cracks, and plastic contact can generate median (M) cracks.  At 
bottom surface, flexural stresses can generate lateral cracks.  All these 
cracks grow transversely though the layer thickness, and ultimately inter-
sect the join interface.                                                                                                     
75
 
 
 
 
76 
In addition, in the case where some plasticity is generated beneath the contact, 
median (M) cracks can initiate and propagate directly downward on median planes 
containing the load axis [61]. These cracks are particularly evident in cyclic loading 
in moist environments, from fatigue efects [62-64].  
Even though the supporting core layer is generaly much stifer, harder, 
tougher and stronger than the vener, it too is susceptible to fracture.  This comes 
about because the combined veneer/core system experiences some flexure beneath the 
concentrated surface load, placing the core undersurface in tension [52,65].  Cracks 
initiate from flaws at the bottom surface, and pop in to form so-caled radial (R) 
cracks which spread upward and radialy outward, again on median planes.  Thus 
atention has to be paid to the possibility of fractures in both veneer and core layers.  
The important point about al these crack systems is that they traverse the 
layers, from the outer surfaces toward the interface.  Idealy, we would like the 
interface to be infinitely strong, to prevent the cracks penetrating from one layer to 
the next, or, more importantly, from delaminating along the interface itself.  Of 
course, no interface is infinitely strong?in fact, most (not al) interfaces tend to be 
weaker than the layers they bond. Generaly, from studies of interface mechanics, if 
the toughnes of the interface is greater than one half that of the bulk material in the 
adjacent britle layer, the crack wil arest and/or penetrate rather than delaminate 
[51]. We wil explore this isue in more detail below.  
As had been made clear in many places in this thesis, it is crucial also to avoid 
high CTE mismatch streses in the composite layer system, because that wil 
introduce some tension in one or both layers. The existence of such streses in any 
 
 
 
 
77 
bilayer that survives the joining proces wil inevitably enhance one or other of the 
fracture modes depicted in Fig. 4.1.  
In what follows, we wil use a simple indentation test developed by 
researchers at NIST [53] for joined bilayers to determine the integrity of our fused 
interlayers and to confirm the absence of significant CTE streses.  We wil follow 
that work closely, since that group has already outlined the basic theory behind the 
methodology.  The major diference in the present study is the presence of the 
additional, intervening glas bond layer betwen the veneer and core layers.  
 
 4.3  Experimental Methods 
 
Vickers indentation tests were caried out on sections of glas-joined 
porcelain/alumina and porcelain/zirconia layers. The sections were cut normal to the 
top surface with a diamond saw, ground and polished with diamond paste to 1 ?m 
diamond finish. A final polish was made with a 0.1 ?m colloidal silica suspension on 
a felt cloth.  Figure 4.2 shows the section geometry.  The thicknes d of the join 
varied betwen 10 to 30 ?m in these specimens. 
Conventional microindentation testing was then caried out on the polished 
sections using a Vickers indenter in a Zwick tester (Zwick, Riverview, Michigan 
USA).  Preliminary tests were made on each of the porcelain veneer and core 
materials before joining, to ensure wel-formed indentations. The sizes of cracks c
1, 
c
2
, c
3
 and c
4
 emanating from the indentation corners were then measured in a high 
power microscope, from which the fracture toughnes was calculated for each 
material using  the following relation [66] 
c
1
c
2
c
4
c
3
h
d
c
1
c
2
c
4
c
3
x
x
a
h
(a)
(b)
E
1
,
E
2
, 
T
1
 
T
2
 
E
1
,
E
2
, 
T
1
 
T
2
 
a
d
Join
Join
Figure 4.2  Schematic showing Vickers indentations in ceramic 
layers joined by glass.  Coordinate system shown for measuring 
lengths c of crack arms at different distances h from interface 
(taken at boundary between core and glass), for crack in (a) veneer 
and (b) core layers.         
 
78
 
 
 
 
79 
  T = ?(E/H)
1/2
P/c
3/2
 
 
where P is the indentation load, c is the size of the crack traveling towards the 
interface measured from indent center to crack tip (c
1
 in the core, and c
2
 in the 
veneer), E is Young?s modulus, H is hardnes, and ? = 0.016 is a dimensionles 
coeficient. The calculated modulus of the glas is 70-78 GPa, while the measured 
hardnes was 4-5 GPa both of which are similar to values for porcelain. 
Indentations were then placed in the specimen sections at prescribed distances 
h from the interface betwen core and fused glas interface, with cracks paralel and 
perpendicular to interface, as depicted in Fig. 4.2. Diferent loads were used in each 
layer material, so as to maintain indentations with wel-defined cracks in each case: 
in the porcelains, P = 10 N (higher loads tended to produce excesive chipping, 
disrupting the crack paterns); in the alumina, P = 10 N, whereas in the zirconia, P = 
35 ? 40 N (needed to initiate corner cracks).  Indents far from the interface were 
compared with those in the dummy, unjoined specimens, to examine for the presence 
of any residual compresion or tension CTE streses?any such streses would reveal 
themselves by shortening or lengthening the crack arms paralel or perpendicular to 
the interface [67]. Typicaly, cracks that are les than 30% longer than the average 
crack length in unjoined material at the same load may be considered to be indicative 
of the presence of an insignificant stres level, i.e. < 30 MPa [54].  
The behavior of the cracks with diminishing distance h from the interface was 
then observed for the layered structures. In particular, observations were made as to 
whether perpendicular cracks arest, deflect, bend or penetrate at the interface. From 
 
 
 
 
80 
such observations, one can evaluate minimum interface toughnes values [53]. 
Bending of paralel cracks toward or away from the interface indicates either presence 
of significant residual stres or a stres field modifying efect of a lower modulus 
interface layer.  
 
 4.4 Results 
 4.4.1 Qualitative observations 
Figures 4.3 and 4.4 show indentations in the veneer/glas/core-ceramic 
systems remote from the glas interface. In these figures, cracks are equi-sized in al 
directions, indicating wel-behaved materials without significant CTE streses from 
the join proces.  The crack paterns are in fact almost identical to those sen in 
control, separate porcelain, alumina and zirconia specimens. As an example, a CTE 
mismatch of 1x10
6
 C
?1
 corresponds to a residual stres of 30 MPa [53], which would 
in turn show as a diference of about 30% in crack lengths in the paralel and 
perpendicular [54]. We observe no measurable diferences at al, certainly not more 
than 10%, within the scater in data. From these observations we can conclude that 
the CTE streses in either of the layers in Figs. 4.3 and 4.4 are < 10 MPa, and can 
therefore be neglected.  
Figures 4.5 and 4.6 show indentations much closer to the glas/core interface, 
such that the nearest, lead corner crack intersects the interface itself.  
(a)
(b)
Figure 4.3  Vickers indentations in porcelain/glass/alumina layer 
structures, remote from bonding glass interface in (a) porcelain 
and (b) alumina.  Note essential symmetry of corner crack patterns 
in both materials, indicating absence of significant residual 
stresses.                         
 
20?m
20?m
81
50 ?m
(a)
(b)
Figure 4.4  Vickers indentations in porcelain/glass/zirconia layer 
structures, remote from bonding glass interface in (a) porcelain 
and (b) zirconia.  Note essential symmetry of corner crack patterns 
in both materials, indicating absence of significant residual 
stresses.                         
 
50?m
50?m
82
 
(a)
(b)
Figure 4.5  Cracks from Vickers indentation in glass-bonded 
porcelain/ alumina bi-layer. Indentations in a) porcelain and b) 
alumina. Note crack arrests at alumina traverses glass bonding 
layer and arrests at alumina interface in a), and penetration across 
glass into porcelain in b). White horizontal lines are used to high-
light the veneer/glass junction.             
50 ?m
20 ?m
83
(a)
(b)
Figure 4.6  Cracks from Vickers indentation in glass-bonded 
porcelain/ zirconia bi-layer. Indentations in a) porcelain and b) 
zirconia. Note crack traverses glass bonding layer and arrests at 
zirconia interface in a), and penetration across glass into porcelain 
in b). Grey horizontal lines are used to highlight the veneer/glass 
junction.                
50 mm
50 mm
84
 
 
 
 
85 
For the cracks in the porcelain veneer (Figs. 4.5a and 4.6a), the lead crack 
arests at the interface. (In the examples shown, the crack tip is dificult to discern, 
but can be sen at the interface in very high magnification.) It was dificult to place 
these cracks much closer to the interface without causing chipping, as reported by 
Kim et al. [53], then only for cracks approaching the core from the porcelain side. In 
no case did the cracks penetrate into the core.  
For indentations in the core, on the other hand, the lead cracks appeared to be 
atracted to the interface, before penetrating into the veneer (Figs. 4.5b and 4.6b). 
This later suggests that the interface does litle to arest upward extending fractures; 
but, most importantly, nor does it indicate any breakdown of the interface itself. 
These are robust joins. Note the bending of lateraly extending corner cracks toward 
the interface in the case of alumina cores in Fig. 4.5. Since we have established that 
there are no significant residual streses in any of the layers, this ?atraction? of the 
crack arms can be atributed to a modification of the local indentation stres field by 
an adjacent low modulus interface layer [68], especialy manifest in the case of the 
porcelain/alumina system where the modulus mismatch is particularly large (note 
tables 2.1 and 2.6 for properties of ceramics and glases respectively). 
 
 4.4.2 Quantitative analysis 
 
To quantify these observations, we plot measured crack lengths c as a function 
of location h relative to the glas/core interface in Figs. 4.7 and 4.8. The subscript on 
the c quantities in the insets define the crack orientation (compare Fig. 4.2). 
c
3
(a)
(b)
Crack length, 
c
 (?m)
 
150
100
0
Interface distance, h (?m)
0 20010050 150 250
50
200
150
0
100
50
Crack length, 
c
 (?m)
 
c
1
c
2
c
4
c
1
c
2
c
3
c
4
Porcelain
Alumina
Porcelain
Alumina
Figure 4.7 Crack sizes c
1
, c
2
, c
3
 and c
4
 versus distance h of indentation 
center to interface in glass-bonded porcelain veneer/alumina core systems. 
Indentations at P = 10 N in (a) porcelain and (b) alumina, orientation 
relative to interface indicated by inset.
P = 10 N
P = 10 N
86
(b)
0
Interface distance, h (?m)
0 200 400 600 800
100
150
200
0
50
c
1
c
2
c
3
c
4
Porcelain
Y-TZP
(a)
Crack lemgth, 
c
 (?m)
 
200
150
100
50
Crack lemgth, 
c
 (?m)
 
c
1
c
2
c
3
c
4
Porcelain
Y-TZP
Figure 4.8 Crack sizes c
1
, c
2
, c
3
 and c
4
 versus distance h of indentation 
center to interface in glass-bonded porcelain veneer/ zirconia (Y-TZP) core 
systems. Indentations in (a) porcelain and (b) zirconia, orientation relative 
to interface indicated by inset. In zirconia graph, the red points represent 
glass AP1040, indents made at P = 40 N, while the blue  points represent 
glass AP1020 indents madeat P = 35 N.
P = 10 N
87
 
 
 
 
88 
In the figures for zirconia system, Fig. 4.8, the data are for two bonding 
glases, and diferent indentation loads (P = 35 N for AP1020 glas bonded specimen, 
and P = 40 N for AP1040 glas bonded specimen). Horizontal dashed lines on these 
plots indicate asymptotic limits c
1
 = c
2
 = c
3
 = c
4
 at large h, i.e. for indents away from 
the interface for remote indentations, although at smal h (i.e. closer to the interface) 
there is wide scater exacerbated by the diferent loads used.  These crack sizes are 
similar to those measured in individual, unbonded ceramic specimens, in which no 
macroscopic CTE streses exist. 
Most interesting are the data for the lead indentations, c
1
 in Figs. 4.7a and 4.8a 
and c
2
 in Figs. 4.7b and 4.8b (indicated by the dark symbol at the respective insets). 
Solid curves through these data are empirical fits.  As the indents approach the 
interface (diminishing h), there are tendencies for the lead cracks in the porcelains 
(Figs. 4.7a and 4.8a) to become smaler (?repulsion?). At very smal I, the lead cracks 
arest at the glas/core interface.  In our case, we saw no delaminations at the 
interface, indicating very good bonding. This contrasts with observations by Kim et 
al. [53], in which delamination occurred at similar fused interfaces but without a glas 
bond. Conversely, lead cracks in the alumina and zirconia cores (Figs. 4.7b and 4.8b) 
become larger (?atraction?).  This indicates that the cracks are sensing a lower 
modulus in the adjacent glas/veneer material. Cracks tend to acelerate when they 
approach a les stif adjacent layer (and especialy when approaching a fre surface).  
 These interactive efects of crack repulsion and atraction has ben 
interpreted by Kim et al. [52,53] in terms of an ?efective toughnes?. Those authors 
begin by defining a stres intensity factor for a Vickers corner crack  
 
 
 
 
89 
K
0
 = ?P/c
3/2
     
 
where ? = ?(E/H)
1/2
 is an elastic?plastic coeficient, E is Youngs modulus and 
H is hardnes, and is ? a dimensionles constant. Figures 4.9 and 4.10 are plots of K
0
 
versus the distance x = c ? h of the lead crack tip from the glas/core interface (Fig. 
4.2).  The limit x = 0 corresponds to intersection of the crack tip with the interface 
origin. The vertical dashed lines represent asymptotic toughnes values T measured 
for the individual porcelain and core materials. Note that the data deviate away from 
these limits as the interface is approached, suggesting a change in the toughnes 
values.  However, the toughnes does not change?instead, the deviations are a 
measure of the interfacial influence on the crack, a factor not built in to the K
0 
equation above.  This influence is analyzed more closely in the work of Kim et al. 
[53], and is not of direct interest here. Of more importance is that the toughnes of 
the porcelains is considerably les than that of the core materials, so that a crack in 
the porcelain wil arest, and that in the core wil penetrate, the interface. This is in 
acord with the observations in Figs. 4.5 and 4.6. 
We may take the analysis of Kim et al. one stage further, using a wel 
documented scheme by He and Hutchinson [51], to determine a lower bound to the 
toughnes of our interfaces. Basicaly, delamination wil not occur if the interfacial 
toughnes is les than about one half that of the adjacent layer material. For indents 
in the porcelain, the cracks arest at the interface. They do not penetrate into the core 
because of the relatively high toughnes of the later material.  
Stress intensity factor, K
0
 (MPa m
1/2
)
0 1 2 3 4 5 6
100
0
-100
-200
100
50
0
-50
Crack coordinate, x (?m)
Crack coordinate, x (?m)
Porcelain
Alumina
Porcelain
Alumina
Figure 4.9  Stress intensity factor as function at different location x = c? h from 
interface, for glass-bonded alumina veneer/core system.  Indentation corner 
cracks in (a) porcelain (x = h ? c
2
) and (b) alumina (x = c
1
 ? h), indicated by 
inset.  Solid lines are empirical fits to data, dashed lines are asymptotic tough-
ness bounds.  
P = 10 N
P = 10 N
90
Stress intensity factor, K
0
 (MPa m
1/2
)
0
1 2 3 4
5
6 7 8
200
0
-600
-200
-400
-800
-1000
600
200
400
0
-200
Crack coordinate, 
x
 (?m)
Crack coordinate, 
x
 (?m)
Figure 4.10  Stress intensity factor as function at different location x = c
1,2
 - h 
from interface, for glass-bonded zirconia veneer/core system.  Indentation 
corner cracks in (a) porcelain (x = h - c
2
) and (b) zirconia (x = c
1
 - h), indicated 
by inset.  Solid lines are empirical fits to data, dashed lines are asymptotic 
toughness bounds. In b) for zirconia, the red points represent glass AP1040, 
indents made at P = 40 N, while the blue points represent glass AP1020 indents 
made at P = 35 N. Note the toughness values converging towards the same 
values. 
Porcelain
Y-TZP
Porcelain
Y-TZP
91
 
 
 
 
92 
Presumably, if we could get the indents closer to the interface, with very high 
contact loads, we might be able to get delamination to occur.  However, we were 
never able to achieve that in any of our experiments.  For indents in the core 
materials, the cracks always penetrated into the veneer, meaning that the toughnes of 
the interfaces was at least one half that of the bulk veneer (Table 2.1), i.e. around 0.5 
MPa m
1/2
.  This is a very respectable toughnes for any interface, and confirms the 
mechanical integrity of our glas bonding.  
 
 4.5 Discusion and Sumary 
 
In this chapter, Vickers indentations were used to probe the efect of a glasy 
join interlayer on the mechanical behavior of layered dental ceramics. The interfaces 
of porcelains bonded by glas to alumina and zirconia were indented, and crack 
lengths were measured as a function of the distance to the glas/core interface. Lead 
cracks originating in the porcelain layer crossed the glas join, aresting at the core 
interface.  Lead cracks originating in the stifer, tougher core were increasingly 
atracted to the weaker porcelain, and penetrated the glas join, extending abruptly 
into the porcelain.  No delamination was observed, confirming good bonding at the 
interface. Measured toughnes values tended to asymptotic limits of 3.5 MPa in the 
alumina, and 5 MPa in zirconia, with values of both veneers tending to 1 MPa, 
corresponding to measured values of toughnes for these two material. Propagation 
of the crack from the core ceramic into the veneer across the interface indicated an 
interface toughnes greater than one half the toughnes of the veneer, i.e. ? 0.5 MPa. 
 
 
 
 
93 
Most importantly, the interfaces never delaminated in our joined specimens, atesting 
to the mechanical integrity of the joins.  
A comparison with previous work on resin-based adhesives, specificaly 
polymer matrix particle-filed joins, is of interest. Such adhesives have the advantage 
of simple procesing, without the need for high temperature fusion. In an exhaustive 
study of such adhesives, Wang et al used a BisGMA-TEGDMA matrix filed with 
Al
2
O
3
, SiO
2
 and diamond particles to increase the elastic modulus of the join, and at 
the same time efect good bonding.  However, a problem faced by resin-based 
adhesives is the degre of integrity within the join material?they tend to be much 
weaker that fused interfaces. They are also susceptible to chemical degradation, and 
to loss of strength during cyclic loading. In our work, we do not have this integrity 
isue as the glas is wel-homogenized by the stage of application as a join, although 
the need for high fusion temperatures is an offset. Another advantage of glas joins is 
that the properties of the glas can be matched more closely to that of the typical 
dental porcelain, eliminating the prospect of catastrophic radial cracking from flexure 
of the porcelain on a soft interface support [69].  
In high temperature joining, residual streses due to CTE mismatch betwen 
adjacent materials is a frequent problem, one which we have avoided by 
compositional CTE matching of the glas join. Residual streses arising from the 
efect of CTE mismatch in dental systems can be calculated from finite element 
calculations, as has been demonstrated by Hermann et al. [54].  As we have sen, 
CTE mismatches below 1x10
?6
 C
?1
 are unlikely to result in residual streses above 30 
MPa, and our crack size measurements indicate wel below that level. This is very 
 
 
 
 
94 
important, because large streses can lead to catastrophic fracture of the layer system, 
either spontaneously or in long-term service, especialy in cyclic loading in reactive 
environments.  
Finaly, some comments on clinical relevance are in order.  This isue has 
been discussed by Kim et al., but only for fused layers without a glas bonding phase 
[53].  Recal from Fig. 4.1 the various modes of fracture that compete in top-loaded 
layer structures. These include inner and outer cone cracks and median cracks that 
initiate in the occlusal contact region and propagate downwards into the top layer (i.e. 
the porcelain veneer), driven by Hertzian streses. There are also radial cracks that 
initiate at the bottom surface of the lower layer (the core) and propagate upwards, 
driven by flexure on the compliant dentin underlayer. Our experiments with Vickers 
corner cracks simulate the likely response of such cracks as they approach an 
intervening interface, in our case the glas bond interlayer. Cracks in the veneer layer 
may then be expected to intersect the glas join layer, continue growing, and 
ultimately arest at the core interface. Only at exceptionaly high loads would such 
cracks deflect along the interface to cause delamination. It is worth repeating that we 
saw no such delaminations in our tests, suggesting an entirely adequate bond. 
Coupled with this is the chemical intermixing evident in the veneer/glas join/core 
difusion profiles considered in Chapter 3, which adds an element of strength to the 
veneer/core bond. On the other hand, cracks that initiate in the core layer wil be 
expected to penetrate directly into the veneer, and potentialy proced through the 
veneer layer to the occlusal surface, exposing the tooth interior to the oral 
environment [70]. However, flexural streses in the tooth are expected to be 
 
 
 
 
95 
compresive at the top surface, thereby inhibiting full fracture.  
We can conclude that glas bonding is an efective way of joining veneer and 
core layers, with inbuilt provisions to resist catastrophic fracture in dental crown 
applications.  
 
 
 
 
 
96 
 Chapter 5: Summary and Future Work 
 
 5.1 Sumary 
 
This thesis has outlined how to fabricate and test veneer porcelain and ceramic 
core joined by glas fusion, for potential applications in dental crowns.  The 
methodology has involved procesing of new glas types, optimizing fusion 
proceses, and mechanical evaluation of the finished layer system.  
Glas preparation and selection has been discused in Chapter 2.  This part 
proved to be a daunting task, because great care had to be taken to match properties 
of the glas to both veneer and core materials. Commercial glases were found to be 
inadequate for this purpose, so new glas compositions had to be designed and 
fabricated.  Sol gel procesing formed the basis of the procesing route.  Diferent 
glas compositions led to various degres of weting, crazing and porosity on the 
ceramic substrates, so the composition and procesing route itself had to be 
optimized. 
The actual joining proces has been considered in Chapter 3.  Again, the 
proces of joining, firing times and temperatures, along with compositional 
considerations, have been discussed.  Some of the combinations produced intact 
interfaces with good characteristics, as determined by optical microscopy.  Simple 
mechanical screning tests were used to eliminate unsatisfactory layer systems. 
Detailed microprobe analysis has been applied to determine interdifusional layers at 
the veneer/glas and glas/core interfaces. 
 
 
 
 
97 
In Chapter 4, the mechanical integrity of the glas joins investigated.  A 
Vickers indentation probe method has been used to introduce controlled cracks in 
veneer/glas/core sections, to investigate the integrity of the join interfaces.  No 
delamination was observed in any of our experiments, confirming strong bonding. 
The indentation method also enabled confirmation of insignificant residual streses, 
as wel as providing an estimate of the join interface toughnes. 
 
 5.2 Future Work 
 
This work has provided the first steps in demonstrating the feasibility of glas 
joining in the preparation of dental crowns. There is room for further study of this 
methodology. Further refinement of the sol gel procesing route is in order, e.g. role 
of finer glas frits in the joining proces.  Our work has examined only porcelain 
fused to alumina and zirconia cores, but other material systems could be explored. 
The possibility of developing a graded structure, by fusing together several layers 
with gradualy changing modulus from veneer top surface to core bottom surface is a 
potential future development. 
From the mechanical evaluation standpoint, there is a need to look more 
closely at the role of water and other chemical species on the aging properties of glas 
joins. Al ceramics are notoriously susceptible to chemical degradation. The possible 
role of cyclic loading is another isue that should be investigated.  
Finaly, there needs to be some atempt to demonstrate transferability of the 
glas fusion technology to the dental laboratory.  This would involve dental 
 
 
 
 
98 
clinicians, to se if the new fabrication methods can withstand oral environments. 
 This is work for the next generation of researchers.  
 
 
 
 
 
 
99 
 Apendix 
CTE Calculations 
 
 
a) Winkelmann Schott and Appen factors table. This table was hyperlinked to the 
next, the Calculation sheet for compositions, to provide a linear correlation betwen 
composition and CTE. In the table we se the efect of each oxide on the CTE, by wt or 
mol % as denoted. Network formers SiO
2
, B
2
O
3
 have the smalest efects for both sets 
of factors (W-S & A). Al2O3?s behavior as an intermediate oxide, has a deleterious 
efect on CTE above certain mol %s. 
 
Oxide Winkelmann 
&  Schott 
Factors ( wt 
% )
M olec Wt. W &  S Factors 
per mol
Appen Factor 
per mol
Appen Factor per mol 
*(10^-7)
B2O3 0.1 69.62 2.32E-07 1.00E-07 1.00
MgO 0.1
SiO2 0.8 60.09 1.60E-06 3.80E-06 38.00
ZnO 1.8
As2O3 2
Li2O 2
P2O5 2
BaO 3 153.33 1.53E-05 2.00E-05 200.00
PbO 3
Al2O3 5 101.78 1.70E-05 -3.00E-06 -30.00
CaO 5 56.09 9.35E-06 1.30E-05 130.00
K2O 8.5
Na2O 10 61.98 2.07E-05 3.95E-05 395.00
 
 
 
 
100 
b)
 
 
E
xa
m
pl
e
 
c
a
l
c
ul
a
t
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on 
s
he
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t
 
f
or
 
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.
 
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y 
ke
e
pi
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t
he
 
vol
um
e
 
of
 
T
E
O
S
 
c
ons
t
a
nt
 
(
20 
m
l
)
,
 
a
nd 
t
r
a
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l
a
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t
ha
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t
o 
a
 
c
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m
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%
 
(
60 
or
 
65 
m
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%
)
,
 
t
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of
 
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m
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(
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,
 
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a
 
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)
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,
 
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a
 
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f
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ppr
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nge
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a
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gl
a
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gl
a
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* 
A
 
F
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t
or
 
X
 
m
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%
 
c
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t
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*
 
 
 
 
 
101 
c) Some of our compositions, as calculated by the Appen factor method. 
 
 
 
 
102 
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110 
 Curiculum Vitae 
 
MEY SAIED 
4319 Rowalt Dr, Apt. 302, College Park, MD 207 
Cel: 617.821.3724, meysaied@umd.edu 
 
 
EDUCATION           
University of Maryland -  Ph.D. Mater. Sci. & Eng. 12 - 2008 College Park, MD 
Royal College of Art -  M.Phil Ceramics Research, 01 - 1997 London, UK 
Imperial College -   B.Eng. Mater. Sci. & Eng, 06 - 1994 London, UK 
 
EXPERIENCE                    
University of Maryland,     College Park, MD (Aug 2003- present) 
? Graduate Research Asistant & PhD Candidate 
? PROMISE Per Mentor (Aug 2004-May 2005) 
? Eleanor Roosevelt High School student research mentor (Aug 2006 ? May 2007) 
Harvard University     Cambridge, MA (Jul 2002-Aug 2003) 
Lab Administrator 
? Grant management, purchasing, and administration. 
Nubian Studies Procedings   Boston, MA (Jan 2002- Oct 2002, pt/time) 
Editing & Publishing Asistant, 
? Layout & desktop publishing of 1998 Nubian Studies Conference Procedings 
Education Development Center   Newton, MA (Sept 2001-Dec 2001) 
Consultant 
? NSF-funded project DigNubia: ancient ceramics & archaeological materials characterization. 
Junetenth Productions     North State, Sudan (Dec 2002- Feb 2001) 
Production Coordinator, ?Nubia & the Mysteries of Kush? Documentary 
National Council for Antiquities & Museums  Bayuda Desert, Northern State, Sudan 
Interpreter and Asisting Archaeologist 
United Nations World Food Program  Khartoum, Sudan (Apr 1999-Dec 2000) 
Reports & Information Asistant, Gender Task Force Member 
Rupert Spira Ceramics      Shropshire, UK (Oct 1997- May 1998) 
Ceramics Apprentice 
 
 
 
 
 
111 
AWARDS             
? Ruth L. Kirschstein NRSA/NIH Individual Predoctoral Minority Felowship Award 
 (May 2006 ?May 2010) 
? University of Maryland Centre for Women & Minorities Graduate School Felowship  
(Oct 2003- Aug 2005) 
CONFERENCES            
? Saied M, Lloyd I, Rekow E,  ?Investigating the use of glas joins for dental restoration 
systems?, # 0658, ADR Dalas, TX, 2008, 
? Saied M, Slepitza J, Lloyd I, OC Wilson J, Janal M, ?Characteristics of glas-joined veneer-
core layers for dental applications?, #2016, IADR New Orleans, LA, 2007 
? Saied M, Wang Y, Lloyd I, OC Wilson J, Rosenblum M, Thompson V, ?Progres on glasy 
joins for laminar dental al-ceramic restorations?, #1880, ADR Orlando, FL, 2006 
? Saied M, Wang Y, Dreyer E, Lloyd I, OC Wilson J, Rosenblum M, ?Glasy ceramic-ceramic 
joins for laminar dental restorations?, #1770, IADR Baltimore, D, 2005 
 
JOURNAL PUBLICATIONS          
? MA Saied, IK Lloyd, OC Wilson Jr. ?Design of join glas composition for layered al 
ceramic restorations?, Journal of Non-Crystaline Solids, in progres. 
? MA Saied, IK Lloyd, BR Lawn, WK Haler, ?Joining Dental Ceramic Layers by Use Of 
Lead-Fre Glas", Journal of Dental Research, in progres. 
? MA Saied, BR Lawn, IK Lloyd, ?Interfacial Characterization of joined laminar ceramic 
restorations?, Dental Materials, in progres. 
 
THESES             
? M.Phil. Thesis Title: An Investigation of Sudanese Raw Materials for Ceramic Manufacture 
? B.Eng. Senior Disertation: Ceramic Matrices and Composites Fabricated by Procesing of 
Silicate Sols and Gels 
  
LANGUAGES            
Fluent in English and Arabic, Proficient in French  
 
REFERENCES            
Available upon request.