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Full text of "Solubility and biocompatibility of glass"

SOLUBILITY AND BIOCOMPATIBILITY OF GLASS 



By 



ARTHUR E. CLARK, JR, 



A DISSERTATION PRESENTED TO THE GRADUATE COUNCIL OF 

THE UNIVERSITY OF FLORIDA 
IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE 
DEGREE OF DOCTOR OF PHILOSOPHY 



UNIVERSITY OF FLORIDA 
1974 




UNIVERSITY nc 



ACKNOWLEDGMENTS 

The author extends his sincere appreciation to his 
advisor, L. L. Hench , for his guidance and encouragement 
throughout the course of this study. Thanks are also extended 
to H. A. Paschall for his time and patience in helping the 
author with interpretation of the histological results of 
this study. The author will be forever indebted to his wife, 
Lisa, whose patience and encouragement made this work possible 
To C, G. Pantano, thanks are extended for the extensive use 
of his equipment and personal time. Finally, the author 
wishes to extend his appreciation to the many students, co- 
workers, and friends who have afforded assistance throughout 
the course of this work. 

This work was supported by the U.S. Army Medical Research 
and Development Command, Washington, D.C. 



TABLE OF CONTENTS 

Page 

ACKNOWLEDGMENTS ii 

LIST OF TABLES v 

LIST OF FIGURES vi 

ABSTRACT xi 

CPiAPTER 

I INTRODUCTION 1 

II THE INFLUENCE OF P"*"^, B^^ AND F"^ ON THE 
CORROSION BEHAVIOR OF AN INVERT SODA- LIME- 
SILICA GLASS 8 

Introduction 8 

Experimental Procedures 11 

Data Analysis 17 

Results '. 18 

Discussion ■ 76 

Conclusions 89 

III AUGER SPECTROSCOPIC ANALYSIS OF BIOGLASS 

CORROSION FILMS 9 3 

Introduction . - 93 

Theory 9 3 

Experimental Procedure 98 

Results 103 

Discussion 122 

Conclusions 129 

IV THE INFLUENCE OF SURFACE CHEMISTRY ON 

IMPLANT INTERFACE HISTOLOGY 130 

Introduction 130 

Experimental Procedure 130 

Results and Discussion 153 

Conclusions 159 



111 



TABLE OF CONTENTS - Continued 



CHAPTER 

V CONCLUSIONS AND SUGGESTIONS FOR FUTURE 
WORK 

BIBLIOGRAPHY * 

BIOGRAPHICAL SKETCH 



Page 

160 
166 



LIST OF TABLES 



Table Page 

1 Bioglass Compositions for Surface Chemistry 
Analyses 10 

2 d-Spacings Obtained from Corrosion Films on 
45S-6I P2O5 and 45B5S5 Glasses Corroded for 
1,500 Hrs . Corresponding d-Spacings of 

Dahllite are Included 86 

3 Bioglass Compositions Selected for Auger 
Spectroscopic Analysis 99 

4 Bioglass Compositions Implanted in Rat Tibiae. . 131 

5 Energy Dispersive X-ray Analysis of the 
Effect of Conditioning Treatment of Bioglass 
Surface 134 



LIST OF FIGURES 



Figure Page 

1 Schematic block diagram of the atomic 
emission spectrophotometer employed for 

solution analyses 14 

2 Time dependent release of Si02 from bulk 
bioglass surfaces into aqueous solution at 

37°C 20 

3 Time dependent release of Na ions from 
bulk bioglass surfaces into aqueous solution 

at 37°C 22 

+ 2 

4 Time dependent release of Ca ions from 

bulk bioglass surfaces into aqueous solution 

at 37°C 24 

5 Time dependent release of P ions from 
bulk bioglass surfaces into aqueous solution 

at 37°C 26 

6 Effect of P2O5 content of bioglasses on the 
variation of alpha with corrosion time 29 

7 Effect of P2O5 content of bioglasses on the 
variation of epsilon with corrosion time .... 32 

8 Infrared reflection spectra of freshly 
abraded Si02 and bioglass composition 

45S-6I P2O5 34 

9 Changes in infrared reflection spectra of 
four bioglasses with increasing phospJiorus 

content as a function of corrosion time .... 37 

10 Changes in infrared reflection spectrum of 
bioglass composition 45S-6'6 P2O5 as a 

function of corrosion time 40 

11 Compositional surface changes of a 45S-6o 
P2O5 bioglass exposed to a buffered aqueous 
solution 43 



LIST OF FIGURES - Continued 



Figure Page 

12 Scanning electron micrographs of corroded 

surface of bioglass compositions 46 

13 Effect of P2O5 content on the ratio of Si/Ca 
for bioglasses corroded 1 hour in an aqueous 
solution buffered at pH of 7.4 and maintained 

at 37°C 48 

14 Time dependent release of Si02 from bulk 
bioglass surfaces into aqueous solution 

at 37°C 50 

15 Time dependent release of Na ions from 
bulk bioglass surfaces into aqueous solution 

at 37°C 52 

16 Time dependent release of Ca ions from 
bulk bioglass surfaces into aqueous solution 

at 37°C 54 

17 Time dependent release of P ions from 
bulk bioglass surfaces into aqueous solution 

at 37°C '^ 56 

18 Effect of B"^^ and F' additions to the bio- 
glass composition 455-6% .P2O5 O'"^ ^he varia- 
tion of alpha with corrosion time 59 

19 Effect of b"^"^ and F' additions to the 
45S-6% P2O5 bioglass on the variation of 

epsilon with corrosion time 61 

20 Changes in infrared reflection spectrum of 
the bioglass 45B5S5 as a function of 

corrosion time "4 

21 Changes in infrared reflection spectrum of 
the bioglass 45S5F as a function of corro- 
sion time 66 

22 A comparison of the infrared reflection 
spectra of the bioglasses 45S-6''o P2O5 > 
45B5S5 and 45S5F after a corrosion treatment 
of 100 hours in an aqueous solution buffered 

at pH 7.4 and maintained at 37°C 69 



LIST OF FIGURES - Continued 



Figure P^g® 

23 A comparison of the infrared reflection 
spectra of the biogla^ 45B5S5 which had 
been corroded for 1,500 hours in an aqueous 
solution and reagent grade hydroxyapatite ... 71 

24 X-ray diffraction analysis of the crystal- 
lization of hydroxyapatite on the surface 
of a 455-6% P2O5 bioglass as a function 

of corrosion time 73 

25 X-ray diffraction spectrum of the crystalline 
hydroxyapatite film on the surface of a 

45B5S5 bioglass corroded for 1,500 hours .... 75 

26 Influence of P2O5 content on the time 
required to override the pH of a buffered 

aqueous solution 83 

27 Influence of B"^ and F" additions to the 
^SS-6-6 P2O5 bioglass on the time required 
to override the pH of a buffered aqueous 

solution 91 

28 X-ray energy level diagram depicting a 

KL,L^ Auger transition 96 

29 Schematic diagram of recording profilometer 
and the type of depth measurement plot 

generated by the profilometer 102 

30 Typical Auger spectra for three depths of 
ion milling of a 45S-6'o P2O5 bioglass 

corroded one hour at 37°C and pH = 7.4 105 

31 Corrosion film profile produced by plotting 
peak magnitudes versus ion milling time for 
a 45S-6°6 P2O5 bioglass corroded one hour 

at 37°C and pH = 7.4 107 

32 Chemical profile expressed in atomic percent 
of a 45S-6''o P2O5 bioglass corroded one hour 

at 37°C and pH = 7.4 110 

33 Chemical profile expressed in mole percent 
for a 45S-6% ^2^5 bioglass corroded one 

hour at 37°C and p?I - 7.4 112 



VI 11 



LIST OF FIGURES - Continued 



Figure Page 

34 ConTDaris on of photoelectron spectra o£ a 
freshly abraded 45S-6% P2O5 bioglass with 
the spectra of a 45S-6% Pz^S bioglass 

corroded for one hour at 37°C and pH = 7.4 . . . 115 

35 Chemical profile expressed in mole percent 
of a 45S-0I P?05 bioglass corroded one 

hour at 37°C and pH = 7.4 117 

36 Chemical profile expressed in mole percent 
of a 45S-3°o P2O5 bioglass corroded one 

hour at 37°C and pH = 7.4 119 

37 Chemical profile expressed in mole percent 
of a 45S-12''ci P2O5 bioglass corroded one 

hour at 37°C and pH = 7.4 121 

38 Changes in the Auger peak heights of 0, Ca, 
P and Si as a function of corrosion time 

for a 45S-6?d P2O5 bioglass 124 

39 Changes in infrared reflection spectrum 
of 45S-0''o P2O5 glass during conditioning 
treatment 137 

40 Changes in infrared reflection spectrum 
of 45S-6'o P2O5 glass during conditioning 
treatment 139 

41 Electron micrograph of junction between 
45S-0% glass and bone three weeks after 
implantation in rat tibia 143 

42 Light microscopy three weeks after implan- 
tation of a 45S-3% P2O5 glass 145 

43 Photomicrograph of a 455-66 P2O5 glass-bone 
interface three weeks after implantation 

in rat tibia 148 

44 Electron micrograph of the junction between 
the corrosion film of a 45S-6% ^2*^5 glass 

and mineralized bone 150 

45 Light microscopy three weeks after implan- 
tation of a 45S-12I glass 152 



LIST OF FIGURES - Continued 



Figure Page 

46 Photomicrograph of a 45S-12% P2O5 glass- 
bone interface eight weeks after implantation. . 154 

47 Electron microscopy of capillary in Figure 8 . . 156 



Abstract of Dissertation Presented to the Graduate Council 

o£ the University of Florida in Partial Fulfillment of the 

Requirements for the Degree of Doctor of Philosophy 

SOLUBILITY AND BIOCOMPATIBILITY OF GLASS 

By 
Arthur E. Clark, Jr. 
December, 19 74 

Chairman: L. L. Hench 

Major Department: Materials Science and Engineering 

The influence of phosphorus, boron and fluorine addi- 
tions on the surface chemical reactivity of a soda-lime- 
silica glass has been investigated. Several techniques, 
including infrared reflection spectroscopy, ion solution 
analysis, scanning electron microscopy, energy dispersive 
x-ray analysis, x-ray diffraction. Auger electron spectros- 
copy and ion beam milling, have been employed to develop 
insight into the morphological and chemical changes which 
occur on glass surfaces corroded in a simulated physiologic 
environment . 

The resulting corrosion layers and the influence of phos- 
phorus, boron and fluorine on their compositions and rates of 
formation are defined. Surface ion concentration profiles 
determined with Auger spectroscopy and ion beam milling 
detail the structural alterations produced by aqueous attack. 
A mechanism is postulated which explains the sequence of 
events leading to the formation of the multiple - layer corro- 
sion structures. 



Having defined the surface chemical behavior of the 
glasses in an invitro environment, an effort is made to 
relate these observations to the response elicited when iden- 
tical glasses are implanted ki laboratory animals. Stable 
interfacial fixation results when specific surface chemistry 
conditions are satisfied. Insufficient or excess surface ion 
concentrations produce negative osteogenesis and fixation 
results . 

Based upon the invivo observations, a theory is proposed 
that an ideal implant material must have a dynamic surface 
chemistry that induces histological changes at the implant 
surface wrhich would normally occur if the implant were not 
present. 



CHAPTER I 
INTRODUCTION 

Orthopedic prosthetic devices are employed for fixation, 
stabilization, and replacement of damaged or diseased bone. 
A wide variety of implant configurations are in use today. 
These include plates, nails, screws and pins for fixation, 
and weight-bearing devices such as hip, femoral, and total 
knee prostheses. 

Historically, metals have played the predominant role as 
prosthetic devices. As early as 1775 AD, evidence in the 
literature documents the use of iron wire to suture fractured 
bone segments together [1]. Since that time numerous metals 
ranging from gold, silver, aluminum, zinc, lead, copper, 
nickel, high carbon steel, low carbon steel, cobalt chromium 
molybdenum alloy, copper aluminum alloy, magnesium, iron, 
titanium, and ti tanium- aluminum- vanadium alloy have been 
investigated as candidates for prosthetic devices [2-7]. As 
might be expected, a wide range of responses is elicited by 
the various metals and alloys. These responses range from 
gross corrosion of the metal and bone necrosis adjacent to 
the implant, to situations in which the presence of the im- 
plant in a physiological environment is well tolerated and 
bone formation occurs in close proximity to the implant. As 



the investigation o£ metallic implants has progressed, a 
series of requirements for an ideal implant material has 
evolved. Included in this list are: (a) high corrosion 
resistance, (b) suitable mechanical properties for the appli- 
cation, (c) excellent wear and abrasion resistance where 
required, (d) good tissue compatibility, (e) structural homo- 
geneity and soundness, (f) non- thrombogenicity , and (g) rea- 
sonable cost [8] . 

Metal devices predominantly in use in this country fall 
into three categories: Type 316, 316L and 317 stainless 
steels (wrought); cobalt- chromium based alloys (cast and 
wrought); and titanium (unalloyed, wrought). These materials 
all exhibit superior corrosion resistance in the physiologi- 
cal environment of the body. However, it has been demonstra- 
ted that there is an absence of adherence between implants 
made from these materials and bone, because there is always 
a fibrous capsule or sheath surrounding the implant and iso- 
lating it from tissue [9,10]. 

The thickness of the fibrous capsule is an indication 
of the degree of tissue acceptability; i.e., the thinner the 
capsule the better the acceptability. The development of the 
fibrous tissue is due to either corrosion of the implant or 
mechanical irritation produced by movement of the implant 
[11,12]. 

The lack of direct attachment of living tissue to metal- 
lic implants can lead to loosening and motion. The resulting 
pain can force surgical removal. Sufficient movement can 



lead to implant failure or bone fracture. As a result of 
this situation, numerous investigations have been initiated 
to find a material which will firmly adhere to bone. 

One approach has involved the use of porous metallic 
implants. The concept involves bone ingrowth into a porous 
surface providing mechanical interlocking. The mechanical 
load is distributed over a wide area, reducing the chance of 
bone necrosis due to stress concentrations at localized sites. 

Hirschhorn e_t a_l. reported deep bone ingrowth into speci- 
mens of sintered Ti and Ti-6A-4V alloy with a pore size of 
200 ym [13]. Welsh ejt al^. documented bone ingrowtli into 
porous Co-Cr-Mo alloy (Vitallium) coatings on solid Vitallium 
rods [14]. 

Galante ej^ al_. [15] used titanium fibers which were com- 
pacted in dies and vacuum sintered. The resulting pore size 
was reported to be within an order of magnitude of the fiber 
diameter. Specimens placed in rabbit and dog femurs revealed 
bone ingrowth after 12 weeks. In another related study, hip 
prostheses were evaluated after 3 months to a year in dogs. 
Deep bone ingrowth and firm stabilization were reported [16]. 
Pore size was 230 ym. 

A process to produce porous metal implants which involves 
the use of a sacrificial metal with a low vaporization temper- 
ature has been developed at Battelle Northwest Laboratories 
[17]. A composite containing the sacrificial metal and the 
implant material is formed and machined to the desired size 
and shape. The implant is heated to vaporize the sacrificial 



metal and then sintered. Cylindrical plugs made with 304 
stainless steel, Ti , and Ti-6A1-4V powders have been implanted 
into dog femurs for time intervals up to 12 weeks. Bone in- 
growth was reported to depths of 2,500 ym [18]. 

A method for plasma spraying titanium hydride powder on 
solid titanium specimens has been developed by Hahn and 
Palich [19]. Implants with a porous surface (pore size 50-75 
ym) were implanted into femurs of sheep for 14 and 26 weeks. 
A significant increase in bond strength was noted when porous 
specimens were compared with implants with smooth surfaces. 
Although histological examination of the bone-porous surface 
was not reported, bone penetration into the pores was postu- 
lated on the basis of the differences in bond strength between 
the porous and non-porous implants. 

The use of porous metal surfaces to anchor prosthetic 
devices to bone seems promising.. One of the major points 
which remains to be shown is the effect of the increase in 
surface area associated with a porous surface and the result- 
ing corrosion which would occur over long periods of time. 

Another area of interest has centered around the use of 
inert porous ceramic materials. Due to their highly oxidized 
state, ceramics are inert materials capable of resisting 
degradation in severe environments [20]. In addition, ions 
incorporated into most ceramics (Na, K, Mg , Ca) are normally 
found in the body. Thus, release of these ions from a cer- 
amic implant would not present as serious a problem as release 
of foreign or toxic elements. 



One of the first attempts involved the use of a slip 
cast mixture of alumina, silica, calcium carbonate and mag- 
nesium carbonate. The resulting porous material (average pore 
size 17 ym) was strengthened by vacuum impregnating with an 
inert epoxy [21]. Openings at the surface were obtained by 
dissolving the epoxy to a depth of 50-70 mils with methylene 
chloride. The composite material was called Cerosium and 
exhibited mechanical properties similar to bone, Evaluation 
of this material revealed little bone ingroivth into the pores. 
This was attributed to a small pore size. In addition, a 
reduction in the strength values of Cerosium which had been 
implanted was related to epoxy degradation by body fluids [22], 

The use of porous calcium aluminate has been investigated 
by Klawitter and Hulbert [23], Calcium carbonate and alumina 
were mixed with water, pressed into pellets, dried, and fired. 
An interconnected pore structure was produced by the break- 
down of the calcium carbonate and the subsequent release of 
C0_ . Pore size was controlled by varying the particle size 
of the calcium carbonate. Invivo studies revealed that a 
minimum interconnection pore size of 100 pm was necessary for 
mineralized bone growth. In addition, there was a lack of 
inflammatory responses due to the calcium aluminate implants. 
The one unusual response was the presence of a layer of 
osteoid ('^^50 ym thick) separating mineralized bone from the 
ceramic composite. The authors speculated that a local alka- 
line pH change produced by hydration of the surface of the 
ceramic composite inhibited mineralization within 50 ym of 



the ceramic. Although there was a lack of inflammatory 
response elicited, the porous ceramic cannot be considered 
completely inert, because of the hydration and resulting 
effect on bone mineralization. 

Hulbert ej^ a_l. [24] have reviewed the invivo behavior of 
numerous porous ceramic materials and found no adverse tissue 
response and .mineralized bone ingrowth into several materials. 

Preliminary investigations have been conducted employ- 
ing dense aluminum oxide (A1_0_) as a prosthetic device [25]. 
The development of a fibrous sheath separating bone and cer- 
amic was noted as the major drawback. 

Graves et_ aJ. have recently reported on the development 
of a resorbable ceramic implant [26]. The concept of a 
resorbable ceramic material has several attractive features. 
The initial pore size can be restricted to values less than 
optimum for bone penetration. This will result in an increase 
in the initial strength of the ceramic. As resorption pro- 
ceeds, enlargement of the pore structure will stimulate bone 
ingrowth. The drop in strength associated with the increase 
in pore size will be compensated for by the presence of the 
new bone. The stress concentration at the implant-bone 
interface of permanent devices is not a problem as the mate- 
rial is completely resorbed with time. There is the poten- 
tial for influencing ossification through the release of 
specific ions incorporated into the ceramic [26]. 

Calcium aluminate ceramics with additions of phosphorus 
pentoxide were implanted into femurs of mature Rhesus monkeys. 



The results pointed to an enhancement of bone formation at 
the ceramic- tissue interface as well as within the ceramic 
as the PtO|- concentration was increased [26]. 

A completely unique approach to the problem of permanent 
fixation has been initiated by L. L. Hench et_ al_. [27-30]. 
The concept involves the use of surface reactive bioglasses 
to achieve intimate bonding between an implant and bone tis- 
sue. Invivo results, obtained at an early stage in the pro- 
gram, in the form of transmission electron micrographs, 
demonstrated glass - ceramic implants intimately bonded to bone 
at 6 weeks with no indication of an inflammatory response to 
the implant [31]. It was suggested that some chemical char- 
acteristics of the implant may have enhanced ossification at 
the glass-bone interface. 

The purpose of this text is to describe a systematic 
study of a series of glasses (referred to as bioglasses) with 
the intent of developing an understanding of their chemical 
surface behavior. New surface sensitive techniques such as 
Auger Electron Spectroscopy and Infrared Reflection Spectros- 
copy along with several other tools have been employed to 
examine the response of bioglasses to an aqueous environment 
maintained at physiologic temperature and pH. An effort is 
then made to relate the observed invitro reactions to a series 
of invivo responses. It is the author's opinion that such an 
approach has been lacking in many previous investigations of 
candidate biomaterials and, hopefully, will serve as a model 
for future studies. 



CHAP»ER II 

THE INFLUENCE OF P^^ , B^"^ AND F"^ ON THE 
CORROSION BEHAVIOR OF AN INVERT SODA-LIME-SILICA GLASS 



Introduction 

The corrosion o£ silicate based glasses can occur by 
either selective leaching or complete dissolution, but usually 
involves a combination of the two. In general, the process 
leads to the formation of a thin film or gel on the exposed 
glass surface with the composition of the gel being signifi- 
cantly different from that of the uncorroded glass. 

The composition and profile of the gel layer are usually 
a direct measure of the durability of the glass. Studies on 
binary soda-silica and lithia-silica glasses have established 
that the corrosion resistance is maximized when the reactions 
at the glass surface lead to the formation of a thin gel with 
a high surface silica concentration [32]. 

A series of invert silica glasses are under investigation 
for use as prosthetic devices [27-30], and it has been demon- 
strated that it is possible to achieve bonding between glass 
and living bone in the body [31]. The biological accepta- 
bility of a soda-lime-silica glass is significantly affected 
by the presence of small amounts of phosphorus, boron, or 
fluorine [ 33- 36 ] . 



The mechanism by which the bond is developed is essen- 
tially a controlled corrosion of the glass which produces a 
surface composition that is compatible ;vith bone. The results 
of this study have shown that the corrosion behavior of the 
bioglasses is directly related to the effects of additions of 
phosphorus, boron, and fluorine on the composition and pro- 
file of the resulting gel. 

Four nondestructive techniques, infrared reflection . 
spectroscopy (IRRS) , ion concentration analysis of the corro- 
sion solution, scanning electron microscopy coupled with 
energy dispersive x-ray analysis and x-ray diffraction are 
employed to characterize the corrosion gels. IRRS provides a 
direct measure of the surface silica concentration [37] , while 
two parameters calculated from the solution data provide a 
measure of the total amount of silica available for gel forma- 
tion [38]. The parameter a is a measure of the extent of 
selective dissolution and varies in magnitude from to 1. 
IVhen a approaches 0, selective leaching predominates. As a 
approaches 1, total dissolution is the controlling process. 
The second parameter, e, referred to as excess silica, is a 
measure of the amount of silica available for gel formation 
and is calculated from a and the concentration of SiO^ in 
solution. (For a detailed discussion see Ref. 38.) 

Six glasses were chosen for study. This series of compo- 
sitions provides information as to the influence of phosphorus 
on the corrosion behavior of the ternary soda- lime-silica 
glass (see comp. 1, Table 1) as well as the influence of 



Table 1 

Bioglass Compositions 
for Surface Chemistry Analyses 



10 



Weight 



1 


. 45S-0% P^O^ 




45% SiO^ 




2 4.5% CaO 




30.5% Na20 


2 


. 45S-3% P-0^ 








45% SiO 




2 4.5% CaO 




27.5% Na20 




3% P^O^ 


3, 


. 45S-6% P^O^ 




45% SiO^ 




2 4.5% CaO 




24.5% Na^O 




6% P^O^ 





Weight % 


4 


. 45S-12% P„0, 




45% SiO„ 




u 




2 4.5% CaO 




18.5% Na20 




12% P^O^ 


5 


. 45B^S5 




40% SiO^ 




5% B^O^ 




2 4.5% CaO 




2 4,5% Na^O 




6% P^O^ 


6, 


. 45S5F 




43% SiO 




12% CaO 




16% CaF^ 




23% Na20 




6t PjOj 



11 

boron and fluorine on the behavior of glass number 3. Glass 
number 3, which contains 6-6 P^O,- is the most compatible with 
bone. Boron and fluorine were added to facilitate flame 
spraying onto metal substrates as they both reduce the melt- 
ing temperature of the glass [39]. 

Experimental Procedures 

The glasses were prepared from reagent grade sodium car- 
bonate, reagent grade calcium carbonate, reagent grade phos- 
phorus pentoxide, reagent grade boric anhydride, and 5 ym 
silica. Premixed batches were melted in platinum crucibles 
in a temperature range of 1250 to 1550°C for 24 hours. Sam- 
ples were cast in a steel mold and annealed at 450°C for 4 to 
6 hours . 

Bulk samples of each composition were prepared by wet 
grinding with 180, 320, and 600-grit silicon carbide paper. 
After a final dry grinding with 600-grit silicon carbide paper, 
samples were immersed in 200 ml of aqueous solution buffered 
at a pH of 7.4. Buffering was accomplished with a physiologi- 
cal buffer (trishydroxymethyl aminomethane) [40]. Stock 
solutions of .2 M tris (hydroxymethyl) aminomethane and .2 M 
HCl were mixed with distilled and deionized water to produce 
a pH of 7.4. Temperature was maintained at 37°C and the 
duration of exposure was varied from .1 to 1,500 hours. All 
sample solutions were maintained in a static state. A Cole- 
man Metrion IV pH meter with ±0.05 pH accuracy was used to 
monitor change in pll. 



12 

Each sample was subjected to infrared reflection analy- 
sis immediately upon removal from the corrosion solution and 
compared with the spectrum of an uncorroded sample. The IR 
radiation reflected from the glass surface is measured over 
a spectrum of wavenumbers from 1,400 to 250 cm . The peaks 
produced are characteristic of the vibrations of specific 
ionic bonds in the glass structure [41]. By comparing the 
reflectance spectra of corroded versus uncorroded glasses and 
also the spectra of glasses of varying composition, informa- 
tion about the type of structural change as well as the rates 
of changes can be obtained [37]. All measurements were taken 
on a Perkin-Elmer 467 Grating Infrared Spectrophotometer 
equipped with a specular reflectance accessory. 

Solution analysis was performed employing atomic emis- 
sion spectroscopy and colorimetric techniques. Figure 1 is a 
schematic block diagram of the atomic emission spectrophotom- 
eter employed for these analyses. Samples of buffered aqueous 
solutions in which glass specimens had been immersed for 
specific periods of time are introduced into the flame through 
the nebulizer burner system. An atom vapor which consists of 
atoms in the ground state and thermally excited states is 
produced in the flame. As atoms in the thermally excited 
states return to the ground state, they emit radiation with a 
wavelength characteristic of the type of atom involved. This 
characteristic radiation, which is isolated in the monochro- 
mator and intensified in the photomultiplier module, can be 



s 


0) 


<D 


t/l 




X 


U 


I— 1 


H 


oj 


6 


S 


o 


nj 



X 


-p 


■P 


3 




1— 1 


<4H 


o 


o 


to 


fS 


J-l 


Oj 


o 



^ 


1— 1 


u 


a 


o 


S 


rH 


(D 


43 






r-t 


L) 


0) 


■ H 


■P 


-M 


(U 


03 


e 


e 


o 


0) 


-p 


^ 


o 


u 


nC 


CO 


Pm 



14 



E 
o 

o 

0. 




o 



O 3 

II 

o 




c 

3 



3 TJ 
^ C 

Z X 



15 

related to the concentration of the atoms in the original 
sample solution. 

The normal procedure consisted of running undiluted sam- 
ples and comparing the results with a series of premixed 
standards witli concentrations ranging from 10, 25, 50, 100, 
150 and 200 ppm of the ionic species being analyzed. Based 
on these results, the unknown samples were diluted into a 
range of 1-10 ppm. Premixed standards of 1 , 2, 4, 6, 8 and 
10 ppm were analyzed and a plot of intensity versus concen- 
tration (ppm) was obtained. The diluted samples were run 
along with the second series of standards. Plotting the 
intensities of the unknown samples on the predetermined stan- 
dard curve enabled one to obtain an accurate measurement of 
the unknown ionic concentration. This method was employed 
to determine calcium and sodium released into solution. 

The colorimetric procedure . involves the use of a Ilach 
Direct Reading Colorimeter which relates the intensity of 
light at a specific wavelength passing through a sample solu- 
tion to the concentration of a particular ion in the solution, 

The colorimetric molybdos i licate method and heteropoly 
blue method were used for silica determination [42]. In both 
of these procedures ammonium molybdate is added to the un- 
known solution, and reacts with any silica present to form 
molybdosilicate acid which has a yellow color. The intensity 
of the yellow color is proportional to the concentration of 
silica in solution. In the heteropoly blue method, the 
yellow molybdosilicate acid is reduced with aminonaptholsul - 



16 

fonic acid to heteropoly blue. The resulting blue color is 
more intense than the yellow and provides a more sensitive 
measurement of the amount of silica [43]. The molybdosili- 
cate method has a range of O-KO ppm, whereas the heteropoly 
blue method has a range of 0-3 ppm. Normal procedure in- 
volved measurement of undiluted samples with the molybdosili- 
cate method, followed by dilution and a second measurement 
with the heteropoly blue method. In both tests oxalic acid 
was used to eliminate interference from phosphate groups. 

The Phos Ver III method [42] was employed for total 
phosphate determination. This method has a range of 0-3 ppm. 
Dilutions were made until two successive dilutions yielded 
the same results. 

Several samples of each composition were examined with a 
Cambridge Scanning Electron Microscope equipped with an Ortec 
Energy Dispersive X-ray Analysis System. In this system a 
lithium drifted silicon detector is used to separate radia- 
tion according to its energy. X-rays, produced as a result 
of the primary electron beam striking the sample surface, 
excite electrons of the silicon atoms. Each of the excited 
electrons absorbs 3.8 eV of energy. Since numerous electrons 
are excited by a single x-ray, the total charge generated 
produces a current which is proportional to the energy of 
the x-ray. The current is then stored in a multichannel 
analyzer according to its amplitude, until a sufficient number 
of x-rays have been counted [44]. 



17 

X-ray diffraction patterns of selected samples were 
utilized to identify the corrosion films which formed on the 
glass surfaces. A Phillips Vertical Dif fractometer with a 
graphite diffracted beam monochromator was employed. Cu Ka 
radiation was used, with tube settings of 40 kV and 15 milli- 
amps. Pulse height selection was utilized to reduce back- 
ground noise . 

Data Analysis 

Sanders and Hench have presented the following equation 
for the calculation of a for binary silicate glasses: 

moles of SiO^ in solution /moles SiO^ in glass 
^ -' ^ " moles R2O in solution 7 moles R2O in glass 

t23 - ^^2 pp^ ^. Mw SiO^ 1-Pm 

where Pm = mole fraction R^O in glass, MW = molecular weight, 
PPM SiO^ = concentration of SiO in solution, and PPM R = 
concentration of R in solution [38]. 

Extension of the relation to a ternary soda- lime-silica 
glass leads to the following modification of equation (2). 

PPM SiO^ 
MW SiO^ ^ ' ^SiO^ 



(3) a = ^ — p 

1/2 PPM Na ^ PPM Ca ^ ' SiO 
MW Na MW Ca 

where P^-^ = mole fraction of SiO^ in glass and all other 
S1O2 2 

symbols are as presented in equation (2). All alpha values 
presented in this text were calculated from equation (3). 



The presence of small amounts of phosphorus, boron, and fluor- 
ine in the bioglasses may introduce slight inaccuracies into 
the absolute magnitudes of the individual alpha values. 
However, the significant in*Formation obtained from the a data' 
is the extent of selective leaching from the silicate network 
Avith time and its effect on the resulting corrosion layers 
which are produced. In this respect, the equation employed 
for the alpha calculations (3) becomes a sensitive indicator 
of the influence of the phosphorus, boron, and fluorine addi- 
tions on the corrosion behavior of the silicate network. 

The equation utilized for the calculation of the excess 
silica (e) was introduced by Sanders and Hench [38] and is 
presented in equation (4) . 

(4) e = PPM Si02 (— ^) 

Results • 

The time dependent behavior of ion release into solution 
is presented in Figures 2-5 for the four glasses with increas- 
ing phosphorus content. The glasses containing 0, 3 and 6 
wt.''6 PoOp exhibit an orderly decrease in the amount of Na, Ca 
and SiO„ in solution, whereas the glass containing 12°6 PoO,. 
reverses the trend \\/ith an increase in SiO„ and Ca released 
compared with the 6°b PoOj. glass. 

Figure 5 shows the phosphorus solution data for the 
three glasses with increasing phosphorus content. The 
behavior of all three compositions is similar in that a linear 



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X^ 




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-P 


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H 


C 


h 


O 


oi 


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+-> 




3 


Mh 


I— 1 


O 


o 




(/) 


o 




m 


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oj 


3 


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i-H 


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<u 


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cr 



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c 


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cu 




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(U 


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03 


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M-i 


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f-^ 


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t/) 






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CLCO 



0) 

u 

m 
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M-i 



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CD -H 



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t:) 


=S 


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ft ;3 





cr 


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Cvi 


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+-> 


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■H 



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ho 
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CO 



O 
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rH 

o u 

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c 


o 


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60 



26 




CO 



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27 

increase is followed by a drop in the phosphorus level. The 
glass containing 3"6 ^yOr exhibits an increase in phosphorus 
released for 100 hours, whereas the glasses containing 6 and 
12% ^9^^ ^^'^'^^ ^ dron after 10 hours. 

The theoretical parameters a and £ are calculated from 
the solution data. Figure 6 is a plot of a, the extent of 
selective leaching, versus time for the four glasses. The 
glass containing 0% ^7*^^ exhibits a behavior which suggests 
that selective leaching predominates throughout the entire 
process. Although the curve initially increases, indicating 
a tendency towards complete dissolution [38], the maximum a 
value attained is only 0.37 and this is followed by a levelinj 
off to an a value of 0.28. As the phosphorus content of the 
glass is increased, the maximum a value achieved increases, 
with the glass containing 12 Po^r having an a value of 0.6 
at 100 hours. 

In evaluating the influence of PoO^ content on the over- 
all corrosion process, Figure 6 can be divided into 3 time 
regimes. During the initial 20 hours of exposure the glasses 
containing 0, 3 and 6 ivt.l Po^c show a fairly consistent 
increase in their respective a values. The curve obtained 
for the glass containing 12°^ PoO^ fluctuates above and below 
the curve of the 61 PoOp glass. In region II a uniform trend 
is observed, i.e., as the P^O;- content increases the a values 
increase. At 100 hours this behavior reverses with the 
glasses containing a larger percentage of PoO;. exhibiting a 
more negative slope as the a values drop (i^egion III). 



o 

•H 
4-1 

nj 
■H 
U 
as 
> 

CD 

4-> 

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1— 1 


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■H 




^ 


s 







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• p 


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to 







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0) 





■p 


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c 




o 4:; 





+-> 




• p 


LO 


IS 







CNI/ ^ 


d, 


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'4^ 







oj 










■P 


D. 


U 


i-H 


Q) 


Oj 


M-l 




i+H t+H 


W 






H 
P 
bO 

•H 



29 



7 


r 






^ 


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••■' '' ' 




) 


^ 






1: 


A 


\\ \ / 


in in m lo 
OOOO 






'•i .' / 


C\J CvJ C\J OJ 






Q. CL D- Q. 






V-.. ; / 


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5~ 6~5~o , 






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1 1 1 1 
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1 ^\ 








1 •■ Av 








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30 



Epsilon is plotted as a function of time in Figure 7. 
As was stated earlier, epsilon is a measure of the amount of 
silica available for film formation. An increase in epsilon 
indicates that a film is forming while a decrease is a result 
of film breakdown. In order for a film to be protective it 
should have a high epsilon value. However, the magnitude of 
epsilon alone does not completely characterize the effective- 
ness of a corrosion film. The profile of the film is an 
important parameter. Thin films with a high concentration of 
silica at the surface (within 5 ym) are much more effective 
at retarding network breakdown and release of silica into 
solution than are thicker films with a more even silica dis- 
tribution. 

The data of Figure 7 illustrate that, as the P20^ content 
increases, the amount of silica available for film formation 
decreases for the glasses containing 3 and 6 wt.% P2°5" '^^^ 
curve for the 12"^ P^^^ glass deviates from this pattern. 

Infrared reflection spectra of vitreous silica and the 
glass containing 6% V^O^ are shown in Figure 8. The vitreous 
silica peak at 1,115 cm" has been attributed to a bond stretch- 
ing vibration of silicon-oxygen-silicon atoms [45,46], while 
the peak at 475 cm"-^ is produced by bending or rocking motions 
of silicon-oxygen-silicon atoms [45,46]. As alkali or 
alkaline-earth oxides are added to vitreous silica several 
events occur. The Si-O-Si (S) stretching peak experiences 
a reduction in intensity and a shift to a lower wavenumber. 
Also, the intensity of the Si-0 rocking (R) peak is suppressed, 



o 

•H 
■P 

oi 

■H 

> 

4-> 
O 



1/1 (U 

(/I S 

.-H -P 



<+-! O 

O P 

4-1 O 

c o 

CD 




CO 



-13 



t/) O 



M-i 


vT) 


O 


1 




LO 


rt 


LD 


i*-i 


•^ 


+-> 




o 


rt 


CD 


o 


P, 


• H 


to 


■P 




•H 


c 


CO 


o 


O 


•H 


P. 


■P 


e 


u 


o 


0) 


u 


1—1 




M-l 


to 





t/i 


fn 


rt 




1—1 


T3 


M 


CU 


O 


!-i 


•H 


cti 


-Q 


!-i 




U-^ Tj 


rt 


C 


i — 1 


rt 



34 




u 

DC 



39NVi331d3U 



35 

In addition, a new peak develops in the region of 950 cm 
The addition of alkali and alkaline earth oxides (i.e., Na^O, 
CaO) disrupts the continuous three-dimensional vitreous silica 
network by producing s i licon-nonbridging oxygens to satisfy 
the new cations (i.e.,- Na or Ca) . The intensity drops of 
the S and R peaks of vitreous silica are due to the decrease 
in the number of Si-O-Si bonds. The new peak at 950 cm has 
been ascribed to bond stretching of the s i 1 icon-nonb ridging 
oxygen atoms (NS) [37]. The shift of the S peak to a lower 
wave number is a result of the change in local environment 
brought about by the presence of the s ilicon-nonbridging 
oxygen-cation groups. Simon and McMahon have indicated that 
the Si-0 bond force constant is decreased by the presence of 
the cationic field of the network modifiers [47]. 

Infrared reflection spectra of corroded and uncorroded 
surfaces from the series of glasses containing PoOr ^^^ pre- 
sented in Figure 9. Comparison of the uncorroded spectra 
with the short and long corrosion times reveals several 
interesting facts. The silicon-oxygen-silicon stretching peak 
(S) at 1,000 cm" begins to sharpen and shift towards the 
location of the Si-O-Si stretching peak for pure vitreous 
SiO (1,115 cm' ) after 15 minutes for the glass containing 
0% P9O . Simultaneously there is a considerable drop in the 
intensity of the si licon-nonbridging oxygen peak (NSX) at 
950 cm' . The silicon- oxygen rocking peak (R) located at 
500 cm' also increases in intensity and sharpness after 15 
minutes' corrosion. In addition, there is a shift in 



Figure 9, Changes in infrared reflection spectra 

of four bioglasses with increasing phosphorus 
content as a function of corrosion time. 
Solutions were buffered at a pH of 7.4 and 
maintained at 37°C. 



-, s*--- 



37 



UNCORRODED 

15 MIN. 

120 MIN. 



A. 45S-0%PoO5 




"istRy 



1200 



800 600 

C. 45 S - 6%P205 



400 




1200 



1000 



800 



600 



400 



— UNCORRODED 

— 15 MIN. 
— 120 MIN. 




1200 



1000 800 600 

WAVENUMBER(CM-I) 



400 



38 



location towards the Si-O-Si rocking vibration frequency o£ 
pure silica (475 cm' ). These trends continue for a corro- 
sion exposure of 120 minutes with one exception. The inten- 
sity of the rocking peak at #475 cm is somewhat lower than 
it was at 15 minutes. 

For glasses with higher phosphorus contents, the 15- 
minute spectra show an increasing preferential attack of the 
silicon-oxygen-silicon stretching peak (S) , and a decreasing 
preferential attack of the silicon-nonbridging oxygen peak 
(NSX) . The increase in intensity, and location of the shift 
of the silicon-oxygen rocking peak (R) are also retarded for 
the higher phosphorus glasses. 

At corrosion times varying from 75 to 120 minutes, there 
is a complete reversal in behavior. For each of the three 
glasses containing PoO,. there is an increase in the intensity 
of the S peak while the intensity of the NSX peak is signifi- 
cantly reduced. The longer corrosion times for each composi- 
tion represent the maximum exposure before the glass surface 
has roughened to the point where the intensity of the spectra 
is reduced to the extent that reliable data cannot be obtained, 
The time required before surface roughening dominates is 
shortened as the PoO,. content of the glass increases. Even- 
tually the spectra of the glasses containing PoOr become flat 
curves with a very low intensity. 

However, with sufficient corrosion time a new infrared 
spectrum develops which is different from that of the glass. 
Figure 10 contains a series of IR spectra which illustrate 



Figure 10. Changes in infrared reflection spectrum 
of bioglass composition 45S-6°6 PoO- as a 
function of corrosion time. 




1400 1200 1000 800 600 400 

WAVENUMBER (CM I) 



41 



the sequence of reactions for the glass containing 6% P^O . 
This new spectrum (see Figure lOd and e) develops for all 
three glasses containing P^O^ , the only variable being the 
length of corrosion treatment required to produce it. The 
new spectrum begins to appear in as short a time as 4 hours 
for the glasses containing 12°6 P^O,-, and takes 12 hours to 
develop for the glass containing 3''o P^Oq- 

X-ray spectra taken from the glass containing 6% PoO^ 
with the energy dispersive system of the SEM are shown in 
Figure 11. The iron peak seen in each of the spectra is pro- 
duced by x-rays originating from a pole piece in the SEM 
column. The variance in the size of the iron peak indicates 
that identical conditions (i.e., specimen tilt angle and 
counting rate) were not achieved for each spectrum. A crude 
comparison of peaks from different spectra can be obtained 
by dividing the peak intensities of the various elements by 
the intensity of the iron peak in the same spectrum. Another 
way of achieving the same end is by comparing the ratio of 
two peaks in one spectrum with the same ratio from another 
spectrum. 

After two hours in solution, the Si/Ca ratio for the 
glass with 6% PoOp has increased from 0.9 to 2.2. In addi- 
tion, the sodium and phosphorus peaks have completely dis- 
appeared. The Si/Ca ratio began to drop after two hours and 
at 1,500 hours was 0.23. The phosphorus peak reappears at 
20 hours and continues to increase with corrosion time. The 
24-hour spectrum shows that the ratio of Si/Ca has dropped to 





f-l 


s 






■p 


(D 









■P 




<f) 


(U 


t/) 




m 


PuX 




CtJ 


CO 


^ 




bH 




>- 




O 




cd 




• H 


, — , 


f-i 




.0 


^ 


1 
X 




Lni~~~ 











(D 




rj 


II 


> 


• 


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<n 




d: 


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P. 


c>\° 


p.f-1 








^ — ' 








1 




Pj t/) 


CO 


c 


t/) 





LO 





•H 


fH 


^ 


•H 


Q 


(J 




+-> 




•iH 


rt 


:3 


>>2; 




I— 1 


bO 




M-4 





!h 


d 





If) 


(D 









p; 


M 


If) 


If) w 


•P 


<D 


;3 







U) 








CD 


CJ 


cu 





t-H 


aJ 


3 


P w 


^ 


cr' 


!-i 




U 


rt 


bO 








P 


CD 


n:3 


d 


■H 


U 


CD 


03 


!=l 


rt 


fH 




(=1 


M-l 


0) rc: 


03 


5-1 M-l 


P 


U 


ID 


<+^ 


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00 


U) 


13 


5 






-ID 




CD 


I— 1 




nd 


bC 


03 


Oj 


0) 


t:) 


e 




C 


•H 








•H 


fH 


•H 


■p 


Oi 


X) 


+-> 




4-> 


£ 


■ H 


TIJ 


rO 


cti 


CD 


<U 


u 





m 






P^ 


1 


(D 


cS 


e 


P^ 


, f-i 







X 


CD 


rt 


u 


(D 


12 






43 




AllSNaiNI 



44 

1.43 while the ratio of Ca/P is 2.4. At 1,500 hours the 
phosphorus peak has reached a sufficient magnitude to make 
the ratio of Si/P (.47) and Ca/P (2.04) several times smaller 
than was observed in the uncorroded glass. 

Micrographs of the corroded surfaces of the four glasses 
with variable phosphorus content are shown in Figure 12. 
Although the exposure time was only 1 hour, a thick film has 
formed on the surface of each glass, indicating a significant 
amount of corrosion has already occurred. Figure 13 is a 
plot of the change in ratio of Si/Ca as a function of PoO,. 
content for the four samples shown in the preceding figure. 
The Si/Ca ratio of each glass in the uncorroded state is also 
included. The ratio of Si/Ca drops significantly as the PoO- 
content of the glass increases. However, the ratio of Si/Ca 
is greater in the corroded glass than in the uncorroded glass 
for all four compositions. 

Figures 14-17 present the time dependent behavior of ion 
release into solution for the glasses which contain boron and 
fluorine. Since these two glasses are variations of the com- 
position containing 6-6 P-Or, its solution data are included 
for comparison. The release of SiO^ and Na into solution 
is similar for the three compositions. However, it should be 
noted that after .1 hour of exposure, the amount of silica 
released into solution is slightly higher for the boron- 
containing glass at every point on the curve. Comparison of 
the glass compositions (see Table 1) reveals that 5 wt.°6 B^O^ 



Figure 12. Scanning electron micrographs o£ corroded 
surface of bioglass compositions, 
(A) 45S-0°s P2O5, (B) 45S-3"ti P2O5 , 
(C) 45S-6I P2O5, (D) 45S-12% P2O5. 
Samples were corroded for one hour in an 
aqueous solution buffered at pH of 7.4 
and maintained at 37°C. The surfaces were 
ground with dry 600. grit SiC prior to the 
corrosion treatment. 



46 



1 






! 

i 

I 


A. 45S-0% PjOg 


B. 45S-3% P2O5 










^^^^. 














C.45S-6%P20Sj D.45S-12O>t)P205 



m U-H 


+-> 


rt o 


f-H 


r— 1 


O M 


md: 


C 


o a 


1^ -rH 


•H 


•M C 


^ -P 


•H C 


03 


:S 03 


Im 


U 


O t3 -TS (jO 


"-H O 


(U 


f-i 


C (U 


oj 0) 


■ H M 


U U-{ 


rt t3 


~-^U-i 


■P -H 


•H :3 


JD fH 


W 4D 


(D Oj 


4^ C 


o o 


f- U 


•pH 


<D 


O -M 


Is rt 


•H 3 




■P t-l 


m ^ 


rt o 


■P o 


!h in 


oi 




Q e 


0) tyi 


CD 


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■P 


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• (fl 


(DUX 


C ^ 


o c/5 


o cr 


t^ 


rt 


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ct) 


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c rt 


^ CX 


O -H 


CD (D O 


U 


C > U 


f- 


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rt cfl S 


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e-H 


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f-i u 


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w u 


f^ m p:: 



48 




..I CO 

ca|u 



M 




O 




■H 




rO 




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rH 


• 


d <-> 


^ 


o 




r-^ 


s 


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5-1 


+-> 


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rt 


(XI 


ri 


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o 


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00 


■P 




3 



(/) 


to 


03 


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(D 


o 


rH 


0) 


CD 


13 


5-1 


cr 




rt 


+-> 




C 


o 


(U 


+j 


t:) 


f:J 


d 


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d) 




p. 


LO 


cu 


(D 


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o) 


(U 


M-( 


S 


1-1 


•H 


3 


H 


U1 



50 




M 



CN4 

O 

CO 



ri4 



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C 

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f O 

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cr 




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52 




CO 



bO 
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<-H 


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r— 1 


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cr 




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c 


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to 


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^d 


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rt 


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54 




CO 



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!-io 



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c 


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Pu to 


0) 


<D 


■D 


U 




03 


o 


M-l 


E 


?H 


•H 


3 


H 


in 



CD 



56 




t0i 



57 



was substituted for SiO Thus, the glass which contains 
boron has the least amount of silica in its bulk composition. 
There is a significant difference in the behavior of 

calcium released into solution (Figure 16). At 10 hours 

+ 2 
there has been more Ca released from the glasses containing 

boron and fluorine than from the glass containing 6-0 P^O^. 
The level of calcium released remains fairly constant through- 
out the remaining 1,490 hours for the glass containing fluor- 

+ 2 
ine , while the Ca release level of the glass containing 

6% PoOr surpasses it at approximately 150 hours. The level 

of calcium released into solution for the glass containing 

boron continues to increase at a slower rate after 10 hours, 

+ 2 
but it remains above the Ca release level of the glass con- 
taining 6% PoOr ior the entire duration of the corrosion 
treatment . 

Up to 10 hours, the concentration of phosphorus in solu- 
tion is very similar for the glass containing fluorine and 
the glass containing 6% PoOq (see Figure 17). After this 
point there is a drastic drop in the P level for the glass 
containing fluorine. The glass containing boron parallels 
the glass containing 61 ^o^c but the P level is signifi- 
cantly lower at every point. 

Figures 18 and 19 show the alpha (a) and epsilon (e) 
data for the glasses containing fluorine and boron as Avell 
as the glass containing 6% ^^7^q- "^he alpha curve (Figure 18) 
for the glass with boron rapidly attains a maximum value of 
.58. After two hours there is a gradual decrease in alpha 



p: 




o 




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in 


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59 





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51 




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62 



and at l,50n hours it has dropped to a value of .25. The 
alpha curve for the glass containing fluorine remains con- 
stant at a value of .45 for two hours, and then increases to 
a maximum value of .56 at 40 hours. After 40 hours alpha 
decreases linearly to a valu^ of .4 at 1,500 hours. 

The amount of silica available for film formation (e) 
increases uniformly for all three compositions for the initial 
10 hours (see Figure 19). After 10 hours, the epsilon values 
for the glass containing boron are significantly higher than 
those of the glass containing 6 °s P9O5 » while the epsilon 
values of the glass with fluorine are lower than those of the 
glass with 6% P^'^r- 

Infrared reflection spectra of the glass containing 
boron (Figure 20) reveal the same sequence of steps as was 
seen for the glass containing 6% PoOc- Initially there is 
selective attack of the silica peak (15-minute exposure), but 
by one hour a silica-rich layer has formed on the surface. 
Surface roughening leads to a drop in intensity of the entire 
spectrum, producing a flat curve at three hours. A new spec- 
trum begins to develop within 7 hours , and is identical to 
the spectrum which was described previously for the glasses 
containing 3, 6, and 12"6 PtO^. 

A similar series of reactions was observed for the glass 
containing fluorine and the results are presented in Figure 
21. One difference between the glass containing fluorine and 
all other compositions was the shape of the peaks in the IR 
spectrum which developed after the spectrum of the glass 



Figure 20. Changes in infrared reflection spectrum 
of the bioglass 45B5S5 as a function of 
corrosion time. 



64 




D. 3 HRS. IN SOL. 




1400 1200 



1000 800 

WAVENUMBER (CM'I) 



600 400 



Figure 21. Claanges in infrared reflection spectrum 
of the bioglass 45S5F as a function of 
corrosion time. 



66 



A. 


45S5F 


FRESHLY ABRADED 


/ # 


1 


—^ 



B. 15 MIN. IN SOL. 




C. 1 HR. IN SOL. 




1400 1200 



E. 7.5 HRS. IN SOL. 




1000 800 

WAVENUMBER (CM'I) 



600 400 



67 



disappeared. Figure 22 enables one to compare the IR spectra 
of the glass containing 60 PoCi the glass containing boron, 
and the glass containing fluorine, after each had been in 
solution for 100 hours. There are three peaks in the wave- 
number region 500-650 cm and the peak at 600 cm has the 
greatest intensity for the glass containing fluorine. The 
spectra of the other two compositions have only two peaks in 
this region and the peak at 560 cm is dominant. In addi- 
tion, the main peak at 1,035 cm is sharper and more intense 
for the glass with fluorine than for either of the other two 
compos itions . 

Infrared reflection spectra of the glass containing 
boron (which had been exposed for 1,500 hours) and reagent 
grade hydroxy apatite are shown in Figure 23. The two spectra 
are very similar, the main differences being the lack of defi- 
nition of the shoulder at 1,085 cm and the broadness of the 
peak at 1,035 cm for the spectnun of the glass surface. 

Figure 24 contains x-ray diffraction curves of the glass 
containing 6''6 P^O^ which was immersed for 15, 100, and 1,500 
hours. This series illustrates the gradual development of an 
amorphous film into a crystalline product. Figure 25 illus- 
trates the diffraction curve of the glass containing boron 
which had been in solution for 1,5 00 hours. 



Figure 22. A comparison of the infrared reflection 
spectra of the bioglasses 45S-6% P2O5 > 
45B5S5 and 45S5F after a corrosion treat- 
ment of 100 hours in an aqueous solution 
buffered at pH 7.4 and maintained at 37°C, 



69 









45 S - 6% P2O5 - 100 HRS. IN SOL. 


-X 


1 1 




1^ — 


1400 


1200 


1000 


800 600 


400 




1400 1200 1000 



800 



600 400 




1400 1200 



1000 800 

WAVENUMBER (CM'1) 



600 400 



if) 0) 

(D ^ +-> 

,£^ 3 -H 

■P O +-> 

^ l# 

<+-! Ph 

O CD n5 

o >, 

rt LO X 
(h - o 

+-> t— I f-i 

O 13) 

(D !-i >^ 

Ph O ^ 
<f) '4-( 

d) 

O CD Oj 

•H Ti fH 
-POM 
O fn 
CD f-i 



O 



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M 
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P 

p 5 l-l 
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<4-i Ln (/5 
O CO 

O LD O 

t/) ^* <D 



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p. Oj 

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< JZi -H 



0) 



71 




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M ■=* 
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O U +J 

CO >4H f:: 

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rH O 

03 +-) 



fn 
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P -H 



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MH P. O 

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03 ;-^ 03 

I >N in 



a; 



73 




AilSNaiNI 



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CM 

CO 

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UJ 

C5 



AIISNBINI 



76 



Discussion 

The behavior of the glass containing 0% P20^ is easily 
interpreted since the results all point to the development 
o£ a silica-rich film through a corrosion reaction dominated 
by selective leaching. The evidence in support of this 
statement is : 

(1) The maximum value of a is .37 (see Figure 6) and 
this occurs at an early stage (10 hours). In order for com- 
plete dissolution to occur, a must approach a value of 1 [37], 

(2) After reaching its maximum value, a rapidly drops to 
.3 and remains near this value for over 1,400 hours, indicat- 
ing no tendency for the film to break down. 

(3) Epsilon (Figure 7) increases linearly with time for 
100 hours and then levels off. The rapid increase in e which 
occurs during the initial 100 hours indicates that a silica- 
rich film is developing. Any tendency for film breakdown 
would result in a drop in the £ curve. Clearly, no such ten- 
dency is observed throughout the entire 1,500 hours of expo- 
sure. 

(4) The infrared reflection spectrum in Figure 9a shows 
immediate selective attack of the silicon-nonbridging oxygen 
peak (NSX) and the development of stretching (S) and rocking 
(R) peaks associated with pure vitreous silica. After two 
hours , the intensity of the entire spectrum begins to drop 
uniformly. This drop is due to greater light scattering as 
the surface roughens. This phenomenon is unfortunate because 



77 

it does not enable one to obtain a quantitative estimate of 
the surface composition. 

Sanders and Hench have shown that infrared reflectances 
are proportional to the amount of species causing them [37]. 
This relationship assumes that the surface is sufficiently 
smooth to produce predominantly specular reflection. This is 
not the case with the glasses under investigation. However, 
qualitative interpretation can lead to information concerning 
the extent of selective leaching from the surface. It should 
be pointed out that IRRS has a maximum depth penetration of 
less than 1 pm for silicate glasses, and is therefore provid- 
ing information about changes occurring at the surface of the 
corrosion film. In this case it can be seen that a surface 
film composed almost entirely of silica forms within 2 hours. 

(5) The use of energy dispersive x-ray analysis shows 
that after 1 hour in solution the ratio of Si/Ca on the glass 
surface increased from .9 to 5.6 (see Figure 13), again demon- 
strating that the glass is being selectively leached, leaving 
behind a silica-rich film. 

The influence of P^O content on the corrosion behavior 
as seen in the data is somewhat complex. Referring to region 
I of Figure 6, the initial change in alpha suggests that the 
glass structure is more uniformly attacked as the P^O;- content 
increases. The glasses containing 6 and 12°6 PoOj. have alpha 
values slightly above 0.5, indicating that a significant part 
of the corrosion mechanism is total dissolution. This is 
substantiated by the IR spectra of Figure 9. Referring to 



the 15-minute exposures, the decrease in intensity of the S 
peak as the phosphorus content increases is a result o£ 
preferential attack of the silicon- oxygen-silicon bonds. The 
thickness of the corrosion film at very early corrosion times 
is less than 1 ym, so the IR spectra are representative of 
the entire film. Within an hour the film thickness has been 
observed with scanning electron microscopy to increase to 
values on the order of 5-10 ym [48]. Then the IR spectra are 
providing information about the surface of the corrosion film. 

The Si/Ca ratios in Figure 13 of the four glasses with 
increasing phosphorus content indicate that a silica-rich 
film has formed on each of the glasses within one hour. How- 
ever, the level of the Si/Ca ratio on the surface decreases 
as the phosphorus content increases, suggesting that the sur- 
face is more uniformly attacked as the phosphorus content of 
the glass increases. The corrosion films in Figure 12 exhibit 
less surface roughness as the phosphorus content increases , 
as would be expected if the glass structure was being uni- 
formly attacked. Examination of the corroded glass surfaces 
with a scanning electron microscope equipped with an energy 
dispersive x-ray system leads to the same conclusion derived 
from solution analysis of the ions leached from the glass 
structure . 

The glass containing 3% P20^ forms a silica-rich layer 
almost immediately, while the 6 and 12% P^^^ glasses show 
preferential silica attack within the first 15 minutes of 
exposure. This behavior is reversed within two hours for the 



79 

glasses containing 6 and 12°o P^^S ^^ ^'^® intensity of the S 
peak increases while the intensity of the NSX peak is reduced 
(see Figure 9), As was discussed earlier, light scattering 
resulting from surface roughness leads to an intensity drop 
in an IR spectrum. The fact that the intensity of the S peak 
increases after the initial drop indicates that a significant 
amount of silica is present on the surface. 

The amount of silica available for film formation (Figure 
7) increases uniformly witli time in region I for all four com- 
positions. It is during this period that the silica-rich 
film forms on the glasses. A break occurs in each of the 
curves in region II. This event corresponds to the formation 
of a calcium phosphate film for the three glasses containing 
P^O- and occurs earlier as the P^O^. content increases. 

Direct evidence for the existence of the calcium phos- 
phate film is presented in Figure 11. The series of spectra 
show the clianges which occur at the surface of the glass con- 
taining 6°6 PoO when it is exposed to an aqueous environment. 
A silica-rich film forms within 2 hours as has already been 
discussed. The phosphorus peak has reappeared in the 24-hour 
spectrimiand the ratio of Si/Ca has dropped. By 1,500 hours 
the phosphorus peak has continued to grow while the silicon 
peak has been drastically reduced. Comparison in Figure 11 of 
the respective ratios of Si/Ca, Si/P, and Ca/P clearly demon- 
strates the formation of a calcium phosphate rich layer. 

The calcium phosphate film is responsible for the infra- 
red reflection spectra which develop after surface roughening 



80 

causes the spectra of the glasses containing phosphorus to 
diminish. The new spectrum is very similar for all the 
glasses containing phosphorus and it develops more rapidly 
as the phosphorus content increases. Figure lOe illustrates 
the spectrum for the glass with 6 °6 P^O which had been immersed 
for 1,500 hours. The peaks occur in two regions, 1,045 cm 
and 560 cm . Levitt et^ al_. have identified fundamental wave- 
numbers for the phosphate ion of hydroxy apatite in these same 

regions [49]. In addition, Nakamoto [50] has predicted that 

- 3 
the infrared active fundamentals of the PO. ion in aqueous 

solution are at 1,080 cm and 500 cm . This evidence, 
along with the simultaneous buildup of calcium and phosphorus 
at the surface, identified from Figure 11, is the basis for 
specifying the origin of the new spectrum as a calcium phos- 
phate compound. . • 

The details of the calcium phosphate compound film forma- 
tion are not completely understood. It has been established 
that after 10 hours, phosphorus which has been leached into 
solution precipitates back onto the glass surface (see Figure 

5) for the compositions containing 6 and \1% P';,0 . In addi- 

+ 2 
tion, Ca release is retarded during this same time period. 

+ 2 
Figure 4 shows a leveling off in the amount of Ca released 

after 10 hours and the effect is more pronounced as the PoO- 
content of the glass increases. The data points in Figure 13 
emphasize this concept. The ratio of Si/Ca drops significantly 
with increasing PoOp content when the four glasses are cor- 
roded under identical conditions. The decrease indicates 



that proportionally less Ca is removed as the phosphorus 
content o£ the glass increases. 

The formation of the calcium phosphate film influences 
the corrosion behavior of the glasses significantly. Its 
effect is seen in region III of Figure 6. As the phosphorus 
content of the glass increases, the a curves descend with 
increasing negative slopes, indicating selective leaching is 
the controlling mechanism. The solution data (see Figures 
2-4) show that both the silicon and sodium release levels off 
during region III but that the Ca release actually increases 

after the calcium phosphate film is formed. This could be 

+ 2 
due to the excessive amount of Ca present in the glass 

compositions as compared to the P^O^ content. Once all the 
phosphorus has been used up in the film formation, the remain- 
ing Ca goes into solution. However, the film acts as a 
barrier to further attack of the bulk glass structure. 

The relative effectiveness of the films in isolating the 
bulk glass from the aqueous environment is demonstrated in 
Figure 26. It can be seen that the time required to override 
the pH of a buffered solution increases as tlie PpOp content 
increases. Since the pH increase results from a sodium- 
proton exchange between the glass and solution [51], the 
formation of the calcium phosphate film retards this reaction 
and the effect is more pronounced as the film formation is 
accelerated. 

Now let us turn our attention to tlie influence of boron 
and fluorine additions on the corrosion behavior of the glass 



xi 


c 


(U 


o 


^ 


•H 


•H 


■M 


:3 


3 


CTt-H 


OJ 


O 


>-4 


m 


(D 


t/i 


e 


^ 


'H 


o 


■M 


(D 




3 





cr 


d 


oj 





0) 


+- 


> ^ 


C ^H 


OJ 


:3 


■M 


rQ 


C 




O 


Oj 


U 






M-i 


LO O 


o 




CN 


la: 


a< 


a 


4-( 


<D 


O X 




+-> 


(D 




U 


o 


CTJ 


0) 


'H 


3 


^^ 


rH 


fH 


U-i 


CD 


p; 


> 


1— 1 


O 



CD 

3 

•H 



83 




O 

d 
o 

o 



o 
d 
o 



€/i 



§ 



,<M 



in in m lo 

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5* o o cvj 
o coco T- 



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CO cow CO 

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(O 



CM 






84 

containing 6% ^o^c- There is a pronounced difference in the 
protectiveness of the calcium phosphate film which forms on 
these glasses. Figure 27 demonstrates the effect of adding 
boron or fluorine to the bulk glass on the time required to 
override the pH of a buffered SK)lution. Obviously, the glass 
containing fluorine is much more effective than either of the 
other two glasses in preventing an increase in pH due to a 
sodium-proton exchange. In fact, the addition of boron 
actually reduces the reaction time necessary to overcome the 
buffering capacity of the solution. 

The reasons for the drastic difference in behavior are 
not intuitively obvious. Both the glass with boron and the 
composition containing fluorine exhibit a behavior similar to 
that of the glass with 6% ^y^S' ^^^^ i^ » initially there is 
selective leaching of silica which ceases after approximately 
15-30 minutes. Within the next -30 minutes a silica-rich film 
is established, and finally a calcium phosphate film is pro- 
duced at the silica-rich film-water interface (see Figures 
20 and 21) . 

The key to the variable corrosion resistance appears to 
be associated with the calcium phosphate films. Initially, 
they appear to be amorphous. Figure 24a contains an x-ray 
diffraction pattern of the surface of the composition contain- 
ing 6*^ ^2*^5 ^^hi*^^^ ^^*^ been in solution for 15 hours. Infrared 
reflection spectra of this sample showed that a calcium phos- 
phate film was present on the surface. The absence of any 
diffraction peaks indicates that the film is completely 



85 

amorphous. However, it is possible that some crystalline 
material is present but not in sufficient quantity to pro- 
duce peaks. A diffraction pattern of the same composition 
after 100 hours in solution shows peaks beginning to appear 
(Figure 24b). Figure 24c is a diffraction pattern of the 
glass containing 6 "o PtO;. which had been in solution for 1,500 
hours. The d spacings obtained from the film show reasonable 
agreement with the d spacings of carbonate hydroxyapatite 
(dahllite). The values are compared in Table 2. There is 
one discrepancy in the relative intensities and that is for 
the 3.402 d value. It is the sharpest peak and has the high- 
est intensity for the calcium phosphate film, whereas it has 
a relative intensity of 70 for dahllite. This effect could 
be accounted for if growth occurred along a preferential 
direction. Figure 25 contains a diffraction pattern of the 
calcium phosphate film on the surface of the glass containing 
boron which has been in solution for 1,500 hours. Again 
there is reasonably good agreement between its d spacings and 
those of dahllite. The relative intensities are also in good 
agreement . 

Referring to Figure 23, the similarity between the infra- 
red reflection spectrum of the reagent grade hydroxyapatite 
and the spectrum of the glass containing boron which had been 
in solution for 1,500 hours takes on added significance. 
Considering the x-ray diffraction patterns, the infrared 
reflection spectra and the energy dispersive analysis which 
shows calcium and phosphorus to be the main components on the 



86 



Table 2 

d-Spacings Obtained from Corrosion Films 

on 45S-6% P2O5 and 45B5S5 Glasses Corroded 

for 1,500 [Irs. Corresponding d-Spacings 

of Dahllite Are Included. 



Dahllite 


4 


. 120 


3 


.402 


2 


. 76 8 


2, 


.687 


2, 


,607 


2. 


,232 


1. 


931 


1. 


834 


1. 


721 



4SS-6% PoOr 
1,500 Ilrs in Sol 



3.411 
2 . 769 
2.688 
2.619 
2.268 
1.939 
1. 832 
1. 717 



45B5S5 
1,500 Hrs in Sol 



3 


.411 


2 


. 777 


2, 


.697 


2, 


,619 


2. 


,257 


1. 


,931 


1. 


839 


1. 


713 



87 

surface after 1,500 hours in solution (see Figure 11), it 
would indicate that the crystalline calcium phosphate mate- 
rial which forms contains a considerable amount of hydroxyapa- 
tite. It has been stated by Korber and Tromel [52] that in 
the system CaO-P^O^, hydroxy apatite will form at temperatures 
up to 1050°C if water is not carefully excluded. 

It should be pointed out that the most synthetic calcium 
phosphate precipitates form nons toichiometric crystal com- 
pounds with numerous possible substitutions existing, i.e., 
sodium for calcium, carbonate for phosphate, fluorine for 
hydroxyls , water for hydroxyls. McConnell [53] has stated that 
unless special precautions are taken it is practically impos- 
sible to obtain apatite crystals which do not contain carbon- 
ate groups. Furthermore, he suggests that carbonate substitu- 
tion for phosphate groups can produce distortion in the hexa- 
gonal apatite structure which can lead to line splitting in 
diffraction patterns. 

It thus seems likely that the calcium phosphate film 
which forms at the silica-rich film-water interface of the 
glasses containing phosphorus is indeed hydroxyapatite . 
However, it almost surely deviates from s toichiometry due to 
substitution of carbonate, sodium and possibly silicon. 

One explanation for the significant difference between 
the protectiveness of the calcium phosphate film of the glass 
containing fluorine and all of the other compositions is that 
the fluorine substitutes for the hydroxyl ions in the apatite 
structure. It has been reported that if water containing 



trace amounts of fluorine is brought into contact with hydroxy- 
apatite, fluorapatite will form as an insoluble product [54], 
Another source [55] has stated that in aqueous systems con- 
taining trace amounts of fluorine, fluorapatite is the most 
stable calcium phosphate compound. Referring to Figure 17, it 
can be seen that there is a drastic drop in the phosphorus 
level in solution between 10 and 100 hours for the glass con- 
taining fluorine. The level of calcium released into solution 
is also significantly lower after 100 hours for the glass con- 
taining fluorine, when compared to the data for all other 
glasses examined (see Figure 16). 

The main influence of boron is an acceleration of the 
initial attack of the glass network. Figure 14 illustrates 
that even though the glass containing boron has the least 
amount of silica in the bulk composition, more silica is 
released into solution than is released from the glass con- 
taining 6% PoOr or the glass with fluorine. This effect is 
thought to be due to a weakening of the three-dimensional 
silica network due to the presence of the boron atoms. Boron 
can exhibit either three-fold or four-fold coordination. It 
has been reported [56] that at high temperatures, boron pres- 
ent in borosilicate glasses exhibits three-fold coordinati- 
which changes to four-fold at lower temperatures. However, 
during the cooling process there is not sufficient time for 
complete reordering and some of the boron remains in three- 
fold coordination. It is the presence of the boron atoms 
with three-fold coordination which produce weak regions in 



.on 



89 

the glass network. Aqueous solutions attack these areas, 
releasing substantial amounts of boron and sodium. 

A similar type of behavior could account for the observed 
surface reactions of the glass containing boron. The pres- 
ence of three-fold coordinated boron atoms lead to an accel- 
erated release of sodium and boron atoms. This would pro- 
duce a more rapid overriding of a buffered solution which 
has been observed (see Figure 27). Release of silica would 
also be accelerated due to the increased basicity of the solu- 
tion. The data in Figure 18 substantiate this hypothesis. 
The addition of boron to the glass containing 6°6 P-^O results 
in an increase in the initial alpha values , which is a sign 
that the extent of total dissolution is increasing. It 
should be noted that this event is only temporary as a silica- 
rich film is established within 1 hour. Tlie epsilon curve of 
Figure 19 shows an increase in magnitude of e for the glass 
containing boron which is greater than the glass containing 
6% PyOr, indicating there is more silica available for film 
formation . 

Conclusions 



In summary, the following facts have been established: 
(1) The glass containing -d Pt'^c forms a silica-rich 
film which protects the glass throughout 1,500 hours of expo- 
sure . 



o 


m 




CMO 




p^ 


E 




o\° 


ft 




vD 


<D 




in 


^ 




\r> 


■p 




^ 







<o 


"d 




^ 


■H 




+J 






o 


U 




+J 


> 

o 




in 






(^ 


o 




o 


■M 




■H 






■P tJ 




■H 


H) 




■Xi 


fn 




nd 


•H 


• 


nJ 


:3 


c 




cr o 




P 


•H 


Uh 


rH 


■P 

3 


T? 


(D 


rH 


C 


e 


O 


a 


•H 


t/) 



4- (D :=! 

m ^ o 





^-> cu 


<-H 


3 


O 


C CJ' 




O nj 


d) 




u 


in "x! 


Cl 


W CD 


(D 


Oj p 


;=! 


rH (U 


I— 1 


bcqn 


tw 


O M-( 


Pi 


■H 3 


1— 1 


^ -£3 



CD 

3. 
bO 
•H 



91 




o 

g 



o 
d 

o 



6 



in 



0)00 CO 

4 • ■ 



tf> 



o 
d 



00 



(O 



C4 



92 

(2) The glasses containing phosphorus also form silica- 
rich films. However, in the case of the glasses containing 
6 and 12% phosphorus, the silica-rich film formation is pre- 
ceded by a short period (15-30 minutes) of selective silica 
attack. • 

(3) After the silica-rich film formation, the phosphorus 
containing glasses form a calcium phosphate film at the 
silica film-water interface. The rate of formation of the 
calcium phosphate film is accelerated as the amount of phos- 
phorus in the bulk glass composition is increased. 

(4) Although the calcium phosphate film appears to be 
amorphous initially, it crystallizes with time into an apa- 
tite structure. 

(5) The calcium phosphate film is more effective than 
the silica-rich film in isolating the glass from its aqueous 
environment. 

(6) The addition of fluorine to the glass containing 
6% PoOr significantly increases the resistance of the glass 
to aqueous attack. 

(J) The addition of boron to the glass containing 6% 
PpO^ accelerates the initial dissolution process in an aqueous 
solution . 



CHAPTER III 

AUGER SPECTROSCOPIC ANALYSIS OF 
BIOGLASS CORROSION FILMS 



Introduction 

Auger electron spectroscopy has been employed to further 
characterize the corrosion films which form on a series of 
bioglasses. An investigation by Clark and Hench [48] has 
established that when exposed to an aqueous environment, a 
silica-rich film forms on the glasses within two hours. A 
second film composed primarily of calcium and phosphate is 
produced at the silica film-water interface. This second 
film is produced only when phosphorus is contained in the 
glass composition and the rate of formation is related to the 
amount of phosphorus in the bulk glass. IRRS, EDXA, and X-ray 
diffraction confirmed that the film crystallized into an apa- 
tite structure with time. Auger electron spectroscopy has 
been utilized to obtain detailed chemical profiles of the 
corrosion films in hopes of elucidating the mechanism of film 
formation. 

Theory 

The technique involves bombarding the sample surface with 
a beam of monoenergetic electrons. A series of interactions 

93 



94 

leads to the release of electrons which were contained in the 
electronic structure of the surface atoms. Figure 28 illus- 
trates such a series of interactions. Impinging electrons 
from the beam create a vacancy in the K shell. An electron 
from one L shell then cascades back into the empty slot in 
the K shell. In the process, sufficient energy is available 
for the ejection of an electron from another L level. This 
process is termed an Auger transition and the electron with 
an energy characteristic of the atom from which it was 
elected is called an Auger electron. The Auger electrons 
produce peaks in the secondary electron energy spectrum and 
thus by monitoring the energy distribution due to Auger elec- 
trons, it is possible to identify the atoms producing them. 
In actual practice, the derivative of the energy spectrum is 
taken, which enhances the Auger peaks and suppresses the 
background present in the secondary electron distribution 
[57]. Due to a short mean free path, Auger electrons have a 
maximum escape depth of 50 A, making this a truly surface 
sensitive process. In addition to atom identification, it is 
possible to relate the amplitude of the Auger peaks to the 
concentration of the atoms producing them. 

A complementary process of Argon ion bombardment removes 
surface atoms a layer at a time. By simultaneously ion mil- 
ling the surface and measuring Auger spectra it is possible 
to obtain a chemical profile of the structure. 

The raw data directly observed are the changes in peak 
height with ion milling time. In order to obtain quantitative 



Figure 28. X-ray energy level diagram depicting a 
KL-.L„ Auger transition. 



96 



AUGER DE-EXCITATION 



KL^L2 ^^iV Electron 




lence/ Band 




initial 
ionization 



97 

information about the amount of atoms present at the surface, 
the differences in Auger transition probabilities for differ- 
ent atoms must be considered. Factors contributing to these 
differences are the influence of the environment on an atom's 
electronic structure as well as the distribution of atoms 
within the volume of material producing the detected Auger 
electrons. 

To overcome this problem, sensitivity factors were deter- 
mined by a recently developed process [58], These factors 
normalize the Auger peaks, enabling one to make a quantitative 
comparison of one component with respect to another. The 
sensitivity factors were obtained by analyzing Auger spectra 
of uncorroded glasses which had been ion milled for long 
periods of time to expose the bulk structure, and comparing 
these data to the known glass composition. Modifying the raw 
data with the sensitivity factors allows one to obtain a mea- 
sure of relative atomic percent versus ion milling time. 

By assuming that the cations are present as specific com- 
pounds with oxygen, i.e., SiO^ , CaO, P^^"; ' ^^® relative atomic 
percent data can be altered to provide a measure of mole per- 
cent versus ion milling time. There was usually an excess of 
oxygen near the surface which was unaccounted for. The extra 
oxygen atoms are probably associated with hydrogen atoms 
(which cannot be detected with AES) as water molecules. 
Although approximations are involved in determining the amount 
of species present, the observed changes in peak height with 
ion milling time correspond to an increase or decrease in the 



amount of species at the surface and are unaffected by the 
approximations. 

Experimental Procedure 

The four glass compositions selected for investigation 
are listed in Table 3. The glasses were prepared from reagent 
grade sodium carbonate, reagent grade calcium carbonate, 
reagent grade phosphorus pentoxide, and 5 ym silica. Pre- 
mixed batches were melted in covered Pt crucibles in a tem- 
perature range of 1250 to 1350°C for 24 hours. Samples were 
cast in a steel mold and annealed at 450°C for 4 to 6 hours. 

Bulk samples of each composition were prepared by wet 
grinding with 180, 320, and 600 grit silicon carbide paper. 
After a final dry grinding with 600 grit silicon carbide 
paper, samples were immersed in 200 ml of aqueous solution 
buffered at a pH of 7.4 (trishydroxymethyl aminomethane 
buffer). Temperature was maintained at 37°C, and all sample 
solutions were maintained in a static state. Samples of each 
of the four compositions were immersed in buffered aqueous 
solution for one hour. In addition, samples of the glass 
containing 6% ^2'^5 ^^^^ exposed to the buffered aqueous solu- 
tion for 10, 20, 30, 40, 50, and 60 minutes. 

The samples were placed in a stainless steel vacuum 
chamber maintained at a background pressure of 1 x 10" Torr. 
To prevent destruction of the corrosion films, the beam cur- 
rent was held at a low value (5-10 ya) and was slightly 



99 



• Table 3 

Bioglass Compositions Selected for 
Auger Spectroscopic Analysis 



45S-05„ 1 


-2°5 


45 


Wt.''6 


SiO, 


24. 


5 wt. 


, % CaO 


30. 


5 wt. 


, % Na20 



45S-6"6 ] 


^2^5 


45 


wt . % 


SiO^ 


24, 


5 wt 


. -6 CaO 


24. 


5 wt 


.% Na20 



6 wt."o P2O5 



45S-5I P -,0 



45 wt 



SiO 



2 4.5 wt. % CaO 

2 7.5 wt.% Na20 

3 wt. % ?^0^ 



45S-12°6 \\0, 



45 


wt. % 


SiO, 


24. 


.5 wt. 


, % CaO 


18, 


,5 wt, 


, ?6 Na20 


12 


wt. % 


P2O5 



100 

defocused. Previous attempts to obtain spectra with a beam 
current of 75-100 Ma resulted in complete degradation of the 
films. The beam energy was 3 KV for the series of samples 
corroded for one hour and 2 KV for the 10-60 minute exposures 
of the glass containing 61 ^y^c- The angle of incidence of 
the electron beam was kept at 45° to prevent unstable charg- 
ing on the surface. The energies of the emitted Auger elec- 
trons \Nrere measured with a cylindrical mirror electron analy- 
zer. 

Ion bombardment of the sample surface with 2 KV Argon 
ions was employed to remove the outermost atoms. As discussed 
in the previous section, the concurrent use of milling and AES 
produces a chemical profile of the corrosion films. 

Profiles were determined for each of the four composi- 
tions corroded for one hour. Two silicon peaks can be seen 
in the Auger spectra of Figure 30. It was observed that the 
low energy silicon peak (78 eV) changed shape as the sample 
was ion milled. The correlation between peak size and atom 
concentration does not hold if the peak shape varies. As a 
result, the high energy silicon peak (1,630 eV) was measured 
for the silicon profiles. 

A recording profilometer with a sensitivity of .02 ym was 
employed to calibrate the ion milling rate. Figure 29 con- 
tains the type of plot generated by the profilometer. Using 
the value obtained and assuming a uniform milling rate, cal- 
culations were made to convert ion milling time to depth, 
yielding an estimate of the corrosion film thickness. 
















Tj 







rt^ 




rt ■ 


+J 




fn 


>^ 




O Xi 




+J 






CD ' 


XJ 




S 







o 


+J 




rH 


Cj 




•H 


fn 




m 







o 


C 




i-< 







P- 


W) 




CxO-M 




C 


o 




•H 


1-H 




-d 


P. 




>H 






o 


+-> 




O 


p! 




•0 







f-i 







M-l 


M 




O 


;3 




e 


rt 




Oj 







f-1 


e 




M 






ct3 


^ 




•H 


■p 


!-i 


13 


Ph 







4-> 


O tD 





•H 




g 


■P 


U-i 


O 


rt 


o 


I— t 


c 




• H 


CD 





<-M 


^ 


P^ o 


U 


X ^H 


CO 


■M 


P. 







102 



■D 

O 
U 



a 
Q 



0) 

u 

3 

■o 

(A 

c 

(0 




a 
Q 




103 

Ion milling was not employed on the series of samples 
corroded at ten-minute intervals, as only Auger spectra o£ 
the surface were taken. An attempt was made to measure a 
layer as thin as possible. Since the electrons wliich produce 
the low energy silicon peak have an escape depth (^8 A) about 

O 

one-fourth that of the high energy peak (^30 A), the magni- 
tude of the low energy peak was monitored. The lower beam 
energy (2 KV] was used for these samples to minimize the 
thickness of the detected volume and to prevent radiation 
damage which can lead to splitting of the low energy silicon 
peak . 

Results 

Figure 30 shows Auger spectra obtained at three differ- 
ent ion milling times for the glass containing 6% ^i^z which 
was corroded for one hour. The location of the peaks on the 
abscissa enables one to identify the atoms producing them. As 
was discussed earlier, changes in peak height are caused by 
an increase or decrease in the amount of element in the sur- 
face layer. These changes are most pronounced for the phos- 
phorus and calcium peaks in Figure 30. Plotting the peak 
magnitudes versus ion milling time produces a chemical profile 
as is seen in Figure 31. 

Features of importance are the buildup of phosphorus and 
calcium at the surface, followed by a region in which the 
oxygen, calcium, and phosphorus levels fall off drastically, 



Figure 30. Typical Auger spectra for three depths of 
ion milling of a 45S-6% P2O5 bioglass 
corroded one hour at 37°C and pH = 7.4. 



105 



.X4 



i^ 



Ca 




t=3min 



dN(E) 
dE 




X4 




Bulk 



Glati 



1000 



2000 



Electron Energy, eV 



Figure 31. Corrosion film profile produced by 
plotting peak magnitudes versus ion 
milling time for a 45S-6I P2O5 bioglass 
corroded one hour at 37°C and pH = 7.4. 



107 



Corrosion-Film l^rofile 




Ion Milling Time,min. 



108 



and finally a buildup in the oxygen, calcium and phosphorus 
levels to values characteristic of the uncorroded glass. 
Modifying the raw data with the sensitivity factors and con- 
verting ion milling time to depth of milling produces a semi- 
quantitative chemical profile of the corrosion film. Figure 
32 illustrates the results of this process for the glass con- 
taining 6% PoO;^ which was corroded for one hour. When com- 
paring Figures 31 and 32 it is important to note that, 
although the magnitudes of the elements have been altered 
with respect to each other, the changes observed with milling 
time or depth of milling have been maintained. Ion milling 
through the corrosion films into the bulk glass was achieved 
only for the glass containing 6% ^o'^c. (Figure 32). The thick- 
ness of the silica-rich film is on tlae order of 2.0-2.5 ym, 
while the outermost film rich in calcium and phosphorus is 
only 0,5 ym thick. 

Figure 33 is the result of converting atomic percent of 
surface species to mole percent. This final adjustment of 
the data can only be applied for the corrosion films, because 
the sodium has been leached out. Since the bulk glass con- 
tains a significant amount of sodium which is not detected 
with AES, it would be very difficult to accurately compute 
mole percentages in the region of uncorroded glass. 

The absence of sodium which will be seen in all of the 
chemical profiles is not unexpected. It has been reported by 
several investigators that leaching of alkali is one of the 
initial steps in the corrosion of silicate glasses in aqueous 






O 4-1 

u a 

u o 

•H ^ 

e 

O dJ 

oj o 

• H (D 

Id 

Td o 

<U 't-< 

I/) u 

m o 

u 

P. t/) 

X (/) 

CD nj 

t— i 

O bJj 

.-H O 



CNl . 

t— I CU '^ 
CD 

•H \o 
(U oo 

X Lo a: 



110 




03 
CO 
< 

-J 
3 
CQ 
4k 



-;r 



LU 

H 



LU 
CO 

LL. 



X 

o 
cr 

< 

o 

-J 

CO 

2'x 

1 O 
CO cc 
O 



K- C3 





t3 




c 


ct! 


rt 


fH U 


O ! 


D 


Uh 


t~- 




ro 


4-1 




rt 


+-> 


OJ 


rt 


o 




!-i 


!-i 


(U 


;3 


p. o 




x; 


o 




rH 


o 


O 


s 


g 


o 


(^^ 


•H 


0) 




nd 


t3 


O 


OJ 


^ 


cyi 


r-i 


t/) 


o 


dj 


u 


!-i 




P- 


to 


X 


V) 


(U 


rt 




rH 


QJ 


bC 


r— 1 


o 


• H 


•H 


m 


X) 


o 




^ 


LO 


&0 




CNl . 


I— 1 


Ph 'd- 


rt 




o 


o\o I^ 


• H 


\0 


e 


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OJ 


00 


^ 


LTl "X 


o 


■* p. 



112 



o 

I 

(/} 
IT) 




0) 
U 
(0 

3 
C/) 



a 

0) 

Q 



113 

solution [59]. In spite of these findings, one factor which 
had to be considered is the difficulty in detecting the pres- 
ence of sodium with AES. Previous work [60] has suggested 
that electrostatic conditions produced by electron bombard- 
ment cause the extremely mobile sodium atoms to migrate out 
of the area of analysis. Another possibility is that the 
Argon ion milling process preferentially removes the sodium. 
For these reasons two samples of the glass containing 6% PoOr 
were examined with Electron Spectroscopy for Chemical Analysis 
(ESCA) . This technique involves bombarding the surface with 
a beam of x-rays and detecting the ejected photoelectrons . 
Information on composition and chemical binding can be ob- 
tained from this process. By examining a sample which had 
been corroded for one hour along with an uncorroded sample, 
the absence of sodium in the corrosion films was shown to be 
real and not an artifact of AES.' Figure 34 compares the 
sodium, phosphorus, and silicon peaks for the uncorroded and 
corroded samples using ESCA or photoelectron spectroscopy. 
Chemical profiles of the glasses containing 0, 3, and 
121 PoOq are shown in Figures 35, 36, and 37. They were 
determined by the same technique previously described for the 
glass containing 6% ^2*^5 " Note in Figure 36 that the F2*^5 
level is intensified near the surface but the CaO level remains 
relatively constant and even drops within .05 ym of the sur- 
face. Immediately underlying the phosphorus -enriched region 
is a silica-rich film. The profiles of the glasses contain- 
ing 6 and 121 P„0^ (Figures 33 and 37) both contain areas of 



o 

r-l O ro 



^1 aJ 



CO 





"+^ U ^H 

nJ PL, o 
, (/) X 

O cu 

-P ^ !-i 
U 4-J O 
0) •M <4-l 
P. 12 
CO 13 
to <u 
c! (/) TJ 
O 03 o 
>-i t— I 5h 

+-> w; ^ 
U O o 

CJ 'H (J 

O CO 

o Lo to 
■M O 03 

O (NJt— I 

^ P-, to 

D. O 

o\a .H 

^ VO ^ . 
O I -^ 

C/D LO . 

C Lo O t--- 

O ^ (NO 

to iDh II 

■H TJ 

03 "Tj \0 p^ 
C-i 03 I 

o jD LO fi; 
U 03 >* 03 






115 



u 
a> 

a 

(A 



c 
o 

■*■> 

u 
_aj 

0) 

o 

■«-> 
o 

Q. 






> 



c 

UJ 



c 

■5 

_c 



Figure 35. Chemical profile expressed in mole percent 
of a 45S-O1; P2O5 bioglass corroded one 
hour at 37°C and pH = 7. 4. 



Relative 

Mole 

Percent 



117 



45S-0% P2O5 



SiO. 



'2V 



CaO 




1.0 1.2 



Depth From Surface [** mj 



Figure 36. Chemical profile expressed in mole percent 
of a 45S-3"o P2O5 bioglass corroded one 
hour at 37°C and pH = 7.4. 



45S-3% P2O5 



119 




SiO 



2^L 



60 



Relative 
Mole 40 

Percent 




J I L 



III III I L 



Depth From Surface L^^nJ 



r-i • 

CO • 

it 

05 

o 

Pi rt 


U U 

(U 4-> 

o 

:3 

p; o 

Tj 0) 

in o 
(/I 

f-< 0) 

X O 

0) f-i 

CD O 

rH O 

O CO 

U rt 

tH o 

Oj .H 
O^ 
■ H 

o o 






121 




E 

5. 



Q) 
O 
CO 

1- 

CO 

E 
o 



a 

0) 

Q 



Q) 


-^-i 


> 


c 


•*-' ^ 





03 0) 


O 


— -1 ■"" 


i_ 


O 


0) 


DC^ 


CL 



122 



P and CaO enrichment near the surface with silica-rich 
regions below them. The calcium-phosphorus - rich film of the 
glass containing 12% ^7*^1; ^^ larger than that of the glass 
containing 6°s Po^S' ^ 

Figure 38 presents the raw data from the Auger spectra 
of the sample corroded at 10-minute intervals. The silicon 
peak was not detected after 20 minutes of corrosion, whereas 
the Ca and P levels remained above their uncorroded values 
for the entire 60 minutes. 

Discussion 

The profiles of Figures 33 and 35-37 clearly show the 
existence of silica-rich films for all four glasses. Further- 
more, as the phosphorus content of the glass increases, a 
calcium phosphate film of increasing thickness overlaps the 
silica-rich film. 

The profile of Figure 36 indicates that there is a mini- 
mum phosphorus level which must be reached near the surface 
before the calcium begins to buildup. This level should 
depend on the phosphorus content of the uncorroded glass as 
well as the length of the corrosion treatment. In the case 
of the glass containing 3% PoOr there is not a sufficient 
amount of P^Op to initiate the calcium buildup within one 
hour. Previous work [48] has shown that the calcium phosphate 
film will form at the surface of the glass containing 31 PoOq 
with time. 



00 X! 

p; o 



rt 


1 


U CO 




ut 


" 


^ 


o 






ri 


Mh 




o 


H 




O 


w 


>+H 


+-> 




^ 


Q) 


bog 


•H 


•H 


(U 


+J 


^ 






P5 


r!^ 


O 


OJ 


■H 


<D 


(/) 


ft O 




H 


M 


U 


(D 


O 


bO O 


;3 




<;4H 




o 







^ 


Pi 


■M 


o 




■H 


PI 


•P 


•H 


u 




d 


(fl 


S 





Mh 


bO 




Pi 


rt 


rt 




^ 


tfl 


u 


OS 



124 




c^ 5 " <o ^ S 

(s;!un-qje)}qB!8H >je8d -o;- )|ead JeBnv 



125 



The results shown in Figure 38 point to the formation 
of a thin surface layer (10-15 X) rich in calcium and phos- 
phorus. This layer is established within 20 minutes of cor- 
rosion time during which silicon is preferentially removed. 
This thin calcium phosphorus film is present on the surface 
during the time when the silica-rich layer is forming beneath 
it. In fact, the change from selective silica leaching to 
the formation of the silica-rich film coincides with the time 
when the thin calcium phosphorus layer has formed. The evi- 
dence indicates that the thin calcium phosphorus film prevents 
further preferential silica removal, but allows the other com- 
ponents of the bulk glass composition to be continually 
leached. Once a sufficient amount of calcium and phosphate 
has been leached into solution the thin calcium phosphate 
film serves as a nucleation site for the formation of the 
calcium phosphate layer which eventually crystallizes into 
an apatite structure. One point which is not clear is whether 
the silica-rich film formation which is produced only after 
the thin calcium phosphate layer has formed, plays a role in 
the growth and crystallization of the calcium phosphate film. 

These results are in complete agreement with those pre- 
sented in the previous chapter, and add some additional in- 
sight into the sequence of steps involved in the corrosion 
process. The following series of reactions are now known to 
occur when the glass containing 6% P^Or is placed in an 
aqueous environment buffered at a pH of 7.4 and maintained at 
37°C: 



126 

(1) Within the first 15-30 minutes silica is preferen- 
tially leached. 

(2) During this same time a thin layer rich in calcium 
and phosphorus is established at the surface (10-15 A thick). 

(3) Once the thin calcium-phosphorus layer has formed, 
the preferential silica attack ceases and a silica-rich 
layer, 2-3 ym thick, is formed within one hour. 

(4) After the silica-rich layer has formed and there is 
sufficient calcium and phosphate in solution the thin calcium 
phosphate layer begins to grow. It was reported in the pre- 
vious chapter that the calcium phosphate film formed at the 
silica-rich film-water interface. The techniques which were 
used to characterize the corrosion process were not suffi- 
ciently sensitive to detect the presence of the thin calcium 
phosphate film which forms initially. Only through the use 
of Auger Electron Spectroscopy w.as the detection of this thin 
film possible. 

(5) The calcium phosphate film crystallizes into an apa- 
tite structure with time. 

This sequence of steps can be explained through the 
following mechanism. Phosphorus is a network former which 
exists in four-fold coordination. Due to the +5 charge of 
the phosphorus atom one of the phosphorus oxygen bonds must 
exist as a double bond. McMillan has stated that the exis- 
tence of the double bond in the phosphorus tetrahedra leads 
to conditions which promote separation of the phosphate groups 
from the silica network. Furthermore, he states that it would 



127 

be probable for the P^O- to be associated with alkali or 
alkaline earth oxides present in the glass composition [61]. 
Tomozawa has reported that ^n^r additions to sodium silicate 
and lithium silicate glasses promote phase separation by 
widening the immis cibility boundary and accelerating the 
kinetics [62]. The influence on the immiscibility boundary 

is related to the relative magnitude of the cationic field 

+ 4 +5 
strength with respect to that of Si . P , which has a 

2 2 

larger cationic field strength [Z/a (P) = 1.91, Z/a (Si) = 

1.58] than Si, was shown to promote phase separation while 

+ 4 +4 
Ti and Zr , which have smaller field strengths than Si, 

were both found to suppress phase separation in the soda 
silica system [62]. Although this effect was only substanti- 
ated for simple binary systems, Tomozawa felt that the chances 
for this relation to hold in more complex silicate glasses 
were quite possible. 

Based on these findings, it seems likely that the PoOr 
additions to the soda- lime -sili ca glass promote a tendency 
towards phase separation and, in the process, disrupt the 
silicate phase by tying up some of the calcium from the 
ternary phase. This would have the effect of reducing the 
corrosion resistance of the silicate phase as calcium addi- 
tions have been shown to increase the durability of soda 
silicate glasses [63]. Evidence for phase separation of the 
glass containing 61 P^Or was presented by Hench et_ a_l. [27]. 
A scanning electron micrograph showed a second phase which 
existed as droplets, and was thought to be tlie phosphorus- 
rich phase. 



128 

The net result of this situation would be that the soda 
silica phase would be preferentially attacked by the alkaline 
aqueous solution. This effect would be enhanced as additional 
phosphorus tied up an increasing amount of calcium. As the 
silicate phase is attacked, a surface layer rich in calcium 
and phosphate would be produced which would then shield the 
remaining silicate phase from further network breakdown. 
Diffusion of Ca and Na into solution would still be pos- 
sible, thus leading to the formation of a silica-rich layer 
under the calcium phosphate layer, Wien sufficient phosphate 
and calcium have been released into solution, a reaction 
between these two components and water would cause the calcium 
phosphate layer to grow and eventually crystallize into the 
apatite structure. 

Reactions of this type have been cited in the literature. 
Weyl has postulated that phosphate opacification in soda-lime 
silica glasses is produced by the formation of apatite crys- 
tals [64]. The crystal formation occurs when calcium and 
phosphorus react with water in the glass melt. It was also 
reported that the reaction of calcium and phosphorus with 
moisture in the atmosphere can lead to apatite formation at 
the glass surface, producing surface roughness and brittleness 
of the phosphate opacified glass [64]. 



129 

Conclusions 

1. Chemical profiles have been measured with Auger 
Electron Spectroscopy and ion beam milling which define the 
silica-rich and calcium phosphate corrosion layers. 

2. IVhen the bioglasses are corroded under identical 
conditions, the thickness of the calcium phosphate layer 
increases as the phosphorus content of the bullc glass compo- 
sition increases. 

3. There is a minimum phosphorus level which must be 
reached near the surface before the calcium begins to build up 

o 

4. A thin surface layer ("^10-15 A) rich in calcium and 
phosphate forms during the initial 15 minutes of corrosion of 
the 45S-6°6 PoO,. bioglass. The data indicate that the thin 
calcium phosphate layer initiates the formation of the silica- 
rich layer and serves as the nucleation site for growth of 
the calcium phosphate layer once sufficient calcium and phos- 
phorus have been leached into solution. 



CHAPTER IV 

THE INFLUENCE OF SURFACE CHEMISTRY 
ON IMPLANT INTERFACE HISTOLOGY 



Introduction 

A series of bioglasses with variable phosphorus content 
have been implanted in rat femurs and their response has been 
related to the previously defined invitro chemical behavior. 
In previous invivo studies bioglass implants were treated in 
a conditioning solution prior to implantation. The influence 
of this process on the structure of the bioglass surface has 
been investigated. Infrared reflection spectroscopy and 
scanning electron microscopy with energy dispersive x-ray 
analysis have been utilized to characterize the surface 
changes produced by the conditioning solution. Light micros- 
copy and transmission electron microscopy were employed to 
examine histological sections of the glass-bone tissue inter- 
face . . . 

Experimental Procedure 

Bioglass compositions 1-4 (see Table 4) were selected to 
study the influence of phosphorus additions on the behavior 
of bioglass implants. Samples were prepared under identical 
conditions employed for the invitro studies (see page 11). 

130 



Table 4 

Bioglass Compositions Implanted 
in Rat Tibiae 



131 



1. 



45S-0?6 F\ 0^ 



45 wt. % SiO 
2 4.5 wt 
30.5 wt 



2 
CaO 



Na20 



45S-6°5 P -,0^ 



45 wt. % SiO 
24. 5 wt 
2 4.5 wt 
6 wt 



2 
CaO 



Na20 



^2^5 



45S-3I P,0, 



45 


wt.% 


Si02 


24 


.5 


wt , 


.^0 CaO 


27 


.5 


wt, 


.% Na 


3 wt. 


.i P^Oj 



4. 



5S-12^o P ^O^ 



45 wt.% 
24.5 wt 
18. 5 wt 



SiO, 



CaO 
Na20 



12 wt.% P2OP 



132 

One series containing the glasses with ^o and 61 V^O^ was gas 
sterilized and soaked in conditioning solution for 72 hours. 
Samples of each of these two ^compositions were subjected to 
IRRS and SEM analysis after gas sterilization, 24, 48 and 72 
hours in the conditioning solution. 

A second series was gas sterilized and soaked in condi- 
tioning solution for 72 hours before implantation. The con- 
ditioning solution contains Eagles MEM (Minimum Essential 
Medium) and Earle's balanced salt solution, 10°o fetal calf 
serum, and 10''o newborn calf serum [65]. 

Samples of bioglass 5 mm by 5 mm by 1 mm were placed in 
defects produced in the metaphysis of the tibia just distal 
to the epyphyseal plate of Sprague Dawley male rats. The 
limbs were not immobilized and the animals were s acrif iced at 
3 and 8 weeks. 

The tibiae were dissected clean of all soft tissues and 
the area of bone surrounding the bioglass was cut into 1 mm 
thick sections with bone on either side of the glass. The 
slices of bone and glass were immediately placed in cold 
cacodylate buffered gluteraldehyde , fixed for two hours and 
then washed with fresh cold buffer. The tissue sections were 
then placed in 2% osmium tetraoxide collidine buffered at a 
dH of 7.4 and fixed for an additional hour. After a final 
wash with additional buffer, the blocks were dehydrated in 
■graded alcohols and embedded in Epon 812. Sections were pre- 
pared on a Porter-Blum MT-2 ultra microtome. Thick sections 
(1 ym) were cut with glass knives, stained with Richardson's 



133 

methylene blue azure II stain and examined with a light 
microscope. A diamond knife was used to cut thin sections 
(600 A thick) . Prior to TEM analysis the thin sections were 
stained with saturated fresh alcoholic uranyl acetate and 
lead citrate [66]. All TEM sections were examined with a 
Hitachi HU IIC electron microscope. 

Results and Discussion 

Table 5 illustrates the time dependent change in the 
surface ratios of Si/Ca and Ca/P for the glasses containing 
and 6% P7O1. during the conditioning treatment. These ratios 
were obtained with a scanning electron microscope equipped 
with an energy dispersive x-ray analysis system. X-rays pro- 
duced as a result of the electron beam striking the sample 
surface are detected and identified according to their energy. 
As different atoms have their own discrete energies, the 
resulting spectrum can be used to determine the atoms present 
on the surface. For a more detailed discussion refer to 
page 16. The gas sterilization treatment produces little or 
no change for either composition. After 24 hours in the solu- 
tion there is a significant increase in the ratio of Si/Ca 
for both glasses. In addition, the Ca/P ratio for the glass 
containing 6-6 PoO^ drops drastically. These trends continue 
through 48 hours. Between 48 and 72 hours of exposure the 
ratio of Si/Ca remains constant for the glass containing 0% 
P^O^.. During the same period, the ratio of Si/Ca has dropped 



134 



Table 5 

Energy Dispersive X-ray Analysis of the Effect 
of Conditioning Treatment on Bioglass Surfaces 



Condition of 
Sample 


45S-0^o P^ 
Si/Ca 


°5 


45S-6I 
Si/Ca 


-i^205 
Ca/P 


Freshly abraded 


.910 




.912 


6.2 


Gas sterilized 


.912 




,.912 


6.1 


Gas sterilized + 










24 hrs in cond. sol. 


2,03 




1.43 


2.38 


Gas sterilized + 










48 hrs in cond. sol. 


2.41 




1.75 


1.97 



Gas sterilized + 

72 hrs in cond. sol, 



2.40 



0. 80 



1. 



135 

from 1.75 to 0.80 for the glass containing 6% PpOr , while the 
ratio of Ca/P continued to drop to a value of 1.89. 

Figures 39 and 40 show infrared reflection spectra of 
the glasses containing and 6% PoOr ^'^ selected intervals 
during the conditioning treatment. The spectra of the glass 
with 01 PoOp (Figure 39) reveal the formation of a silica-rich 
surface layer which is present at the conclusion of the 72- 
hour conditioning treatment. Little change is noted between 
the freshly abraded spectrum and the spectrum of the gas 
sterilized sample. After 24 hours in solution, there is 
selective attack of the silicon-nonbridging oxygen peak at 
840 cm . The silicon- oxygen- si licon stretching (S) and 
rocking (R) peaks, located at 955 and 500 cm respectively, 
begin to sharpen, increase in intensity and shift towards the 
location of the S and R peaks of vitreous silica. These 
changes continue to occur through 48 hours of exposure. The 
curve after 72 hours exhibits no additional changes indicating 
a stable condition has been achieved. The data obtained with 
infrared reflection spectroscopy and the x-ray system of the 
scanning electron microscope both point to the formation of 
a silica-rich surface layer on the glass with "o Pt'^'s* This 
glass exhibited the same type of behavior in the invitro 
studies presented in Chapters II and III. 

The IR spectra of the glass containing 6-6 P2O2 (see 
Figure 40) are similar to the spectra of the glass with 01. 
P„Oj^ through 24 hours of exposure. That is, little change 
can be noted between the freshly abraded and gas sterilized 



Figure 39. Changes in infrared reflection spectrum 
of 45S-0% ^z'-'s gl^ss during conditioning 
treatment. 



137 




1200 1000 800 600 

WAVENUMBER (CM-1) 



400 



Figure 40. Changes in infrared reflection spectrum 
of 45S-6I P2O5 glass during conditioning 
treatment. 



133 




1200 



1000 800 600 

WAVENUMBER (CM-1) 



400 



140 

spectra. After 24 hours in solution, selective attack of the 
silicon-nonbridging oxygen peak occurs, and the peaks associ- 
ated with the silicon-oxygen-silicon bonds exhibit changes in 
shape and location which indicate the concentration of silica 
is increasing on the surface. The 48-hour spectrum of Figure 
40 contains the S and R peaks of silica but their intensities 
have dropped to values below their level at 24 hours. This 
trend continues with the 72-hour spectrum. Behavior of this 
type was also observed in the invitro studies on the glass 
containing 6% ^y^c- After the silica-rich layer is formed, 
the calcium phosphate layer begins to grow. Apparently the 
rate of these reactions is slower in the conditioning solution 
and there is not a sufficient amount of calcium phosphate on 
the surface at 72 hours to produce the infrared reflection 
spectrum seen invitro. However, the data obtained with the 
x-ray analysis shows the ratio of Ca/P is becoming smaller 
with time, while the ratio of Si/Ca drops significantly from 
its 48-hour level, indicating an increase in the calcium and 
phosphorus concentration on the surface. 

These observations clearly show that the surface struc- 
ture of a bioglass implant is drastically influenced by the 
conditioning treatment and interpretation of the histological 
results of conditioned samples should take these changes into 
consideration. 

Small pieces of glass implant were attached to bone in 
almost every case, but a distinct variation ivas observed in 
the tissue responses evoked by the different compositions 



141 



which had been conditioned prior to implantation. 

Figure 41 is a transmission electron micrograph of a 
45S-0% P^Or glass-bone interface at three weeks. The mate- 
rial which exhibits the regular fracture pattern appears to 
be the silica-rich corrosion film (CF) which forms on the 
surface of the glass implant. The relative softness of the 
corrosion layer compared to the glass produces the uniform 
fracture pattern, with long non-branching fracture lines. 
The corrosion film contains a tear which was probably produced 
during the sectioning process. Close examination reveals that 
a thin layer of the corrosion film (CF) remains attached to 
bone (B) along the interface (I) , indicating the corrosion 
film-bone interface has considerable strength. The elongated 
cell (EC) in close proximity with bone has the appearance of 
a normal endosteal cell on a resting bone surface and does 
not appear to be actively engaged in laying down new bone. 
Examination of thick sections containing the glass with 0% 
P^O;- revealed a small number of viable osteocytes present in 
newly formed bone and bone surfaces characterized by a lack 
of active bone formation and very few active osteoblasts. 

A 45S-3-6 P-pOr glass-bone interface at three weeks is 
shown in Figure 42. Small pieces of implant are attached 
along the surface. It should be pointed out that before sec- 
tions are cut, the glass is chipped out of the block. If 
this was not done it would be very difficult to cut sections 
as glass knives are used and they would constantly break. 
The presence of small pieces of glass attached to bone 



a 




Oj <D 


bO 


^ 


d 


t/) H 


o 


(/) 


t— 1 


rt 


03 


rH • 


^ — , 


M rt 


^x 


•H 


W o 


o\° ^ 


"— 'o 


O -H 


LTi 


1 -M 


(D " 


CO 


PJ 00 


LO +-1 


O V— ' 


^ nJ 


-P 


^ 




C 


o • 


0) C 


+-> /— > 


(U -H 


-^ 


|3 


-a ^ 


■P S 


s 


<D O 


Hi (/) 


XI -H 


o ;s 


+-> 


rTi o 


sn rt 


^ 


O 4-" 


>^!h 


■H fH rH rt 


+J oj 


-p 


U rH 


^ X 


d P^ bO^ 


=5 S 


•H 


•^■H 


4-1 t3 




CD 


M-i ^ 


I/) +-> 


O O 


•H OS 


+J 


d 


^'+H 


,-^ bO 


p.; Oj pL, .H 


rt 


U t/1 


Jh t/) 


^— ' OJ 


M^ 


■1:3 


O (D 


e 


U CD 


I— ( (D 


O 5 


•H O 


• H 


Mh rt 


6 CD 


Mh 





C ^^ 


a !- 


O (D 


O X 


• H -P 


fH 4-> 


tfl C 


•p 


O -H 


O <U 


fH 


0) C 


f-1 


rH O 


O ^ 


w ^ 


U 4-> 



U ■ 

■;3 

bO: 



143 




WO 0) 

00 u 

'^ -p 
■p 

Mh 

O f-H 
O ,— ^ 

•H U 

+-> ^-^ 
+-> +-> 

I— I 1—1 
Oh P- 

e s 

•H -H 



Sh CO 



(D -P 


O 


10 
i^ P 

^ c 

P 03 

Ph X 
O Cii o 

u o 



to 



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u 01 

■ H Oj . 

e i-i 

P o 

tXO 

■ H CO O 
-J &. P 



145 







/> /^ 



146 



indicates that there is considerable strength associated with 
the glass-bone interface because fracture occurs within the 
glass implant rather than at the interface. 

The mineralized bone adjacent to the implant interface 
of Figure 42 contains several osteocytes and an area of 
unmineralized osteoid. There is a layer of plump osteoblasts 
which appear t'o be laying down new bone. 

Figure 43 is a photomicrograph of a 45S-6% ^y^c. glass- 
bone interface at three weeks. Large pieces of bioglass (G) 
are intimately attached to bone (B) and several normal osteo- 
cytes (0) are present in the mineralized area. There is a 
well-defined layer of osteoblasts actively engaged in laying 
down new bone (OF) and this front is separated from the 
mineralized area by a transition zone of partially mineralized 
osteoid. These features indicate that induction of normal 
osteogenesis has been achieved. An electron micrograph of 
the same section (Figure 44) shows the corrosion layer directly 
attached to mineralized bone along the wavy interface I. 

A 45S-12°6 P-jO^ glass-bone interface at three weeks is 
shown in Figure 45. There is an absence of activity along 
the ossification front with no evidence of osteoid and only 
one osteoblast in the area. Figure 46 is a photomicrograph of 
a 45S-12% P^Oq glass-bone interface at eight weeks. An impor- 
tant feature to note is that the implant G has been separated 
from the bone B by an interval containing a capillary C. 
Electron microscopy of this section (Figure 47) reveals inter- 
cellular crystallization (X) has been induced along the edges 



0) 

u o a 

'^ u o c 

(U cd +-> oi 

•H (U O I 

QJ 03 L) 

C -H nJ t/) 

42 +-> (U 
^ .H 4-> -P 

1 -P 03 >. 
t/1 U , 
(/) P >-, O 
o3 03 1—1 0) 

T-H Jh 0) P 

W) P CO 

C ^ O 

O -H rH 

CNl pi P Oj 

ciH o p; 5h 

•H -H 0) 

o\= p > 

^O 03 0) (D 

I P !-i LO 

00 c; 03 

Ln 03 {/) 

'^ t-l ^> p; 

03 S^ 






03 
P 

o 

CD Cti O 

^ P rH 

o3 oS o a 

fn -HO 

tiO in ^ ^ 
o ^ 

U (D O 

O (D <D 
p (D O . 
O f-i ^— , 

^ ,^ -H pq 
CL, p cx^-^ 



OX 



CD 00 



tiO 



148 




o m 

• H " — ' 

to 

o <u 

>H o 

o ^ 
o 

<U 0) 

^ N 



(D 0) 
O 03 



U P-. 

;3^ 



^ 


(Nl 




P^Cn 




ri 






f-< 


cJP 




CxOvO 




O 


1 




S-H 


uo 




u 


LO 




•H 


■^ 




s 


Oj 


^ ^ 


rt 




X 


O Mh 


o 


f-l 


o 


o 



w m 



150 






";"' ?^'*>("'' # •■■■■"■■■ 



■'^^;:^-\..^'''-'/^^V^ 



» . 'iS V ^■^'■ 









Figure 45. Light microscopy three weeks after implan- 
tation of a 45S-12% glass. Glass (G) is 
attached to bone (B) . There is an absence 
of activity along the new bone surface 
(OF). (1,800X) 



152 




Figure 46. Photomicrograph of a 45S-12I P2O5 glass- 
bone interface eight weeks after implan- 
tation. Glass implant (G) has been 
separated from bone (B) by an interval 
containing a capillary (C) . (1,800X) 



154 




Figure 47. Electron microscopy of capillary in 

Figure 8. Note intercellular crystal- 
lization (X) along edges of capillary, 
(44,200X) 



156 




157 

of the capillary. It can also be observed that part of the 
corrosion film (CF) remained attached to the bone when the 
interval containing the capillary separated the implant from 
the bone. 

Referring to Figure 46, note the unhealthy appearance 
of the osteocytes (0), They have withdrawn from their lacunar 
walls and the nuclei are pyknotic. There is also an absence 
of new bone formation at the bone surface. 

The invivo results of this study show that direct attach- 
ment of glass to bone is achieved within three weeks for the 
four compositions studied. 

The invitro studies in Chapters I and II establish that 
silica-rich corrosion films form on the surface of the bio- 
glasses in a simulated physiologic environment. Furthermore, 
the invitro results of this chapter show that the conditioning 
treatment produces the same response. 

Carlisle has reported that silicon-rich regions are 
associated with active mineralization sites in young mice and 
rats and, once mineralization has gone to completion, the 
silicon content drops [67]. Recent invitro investigations by 
Hench and Paschall [36] have shown that 45S-6'd PoOp glass 
implants are bonded to bone by an amorphous cement- like layer, 
probably comprised of SiO^ , CaO, and PoOi- , which serves as 
the active site for collagen attachment followed by mineral- 
ization. 

In view of the findings of this study as well as those 
in the literature, it seems likely that the silica-rich layer 



158 

serves as an induction site for osteoblasts to lay down the 
organic intercellular substance of bone. This substance 
contains collagen and mucopolysaccharides. Normally, miner- 
alization would begin to occur §s soon as the organic inter- 
cellular substance was secreted by the osteoblasts. The 
exact mechanism of mineralization is not completely defined; 
however, the concentration of Ca and PO^ ions in the area is 
thought to play an important role [68]. 

The phosphorus content of the bioglasses may be the 
important parameter which influences mineralization. The 
buildup of calcium and phosphorus which occurs on the surface 
of the silica-rich films could provide a source of ions for 
mineralization. The results obtained indicate that, as the 
phosphorus content of the glass increases from through 6% 
P the appearance of the total ossification process becomes 
increasingly healthy. In the cs^se of the glass containing 6% 
P„0 the resulting situation is one of normal ossification. 

The results obtained with the glass containing 121 P20^ 
suggest that there is an optimum phosphorus content which 
should not be exceeded. The ectopic crystallization seen in 
Figure 9 might well have been induced by an excessive amount 
of phosphorus. Matthews et al . have reported that the addi- 
tion of phosphates to a fixative, followed by incubation, 
will result in apatite crystal formation [69]. Furthermore, 
they reported that release of phosphate from cells which led 
to the formation of an amorphous calcium phosphate was 
prompted as a response to administered doses of thyrocalcitonin, 



159- 

In the case o£ a bioglass, a specific enzyme would not 
be necessary to release large amounts of calcium and phos- 
phorus as the response of the bioglass surface to body fluids 
would accomplish the same end. If tlie calcium and phosphorus 
released from the glass when combined with calcium and phos- 
phorus present in the body fluids resulted in a critical 
supersaturation , apatite crystal formation would result. 

Conclus ions 

Based upon the evidence obtained, the following theory 
is proposed for implant materials design and selection: 

An ideal implant material must have a dynamic surface 
chemistry that induces histological changes at the implant 
interface which would normally occur if tlie implant were not 
present . 

In the case of the bioglasses the optimal response is 
elicited by a composition which has the ability to form a 
silica-rich corrosion film and provide an adequate but not 
excessive supply of ions to be incorporated in the minerali- 
zation process. The glass containing 6-0 PoOr appears to be 
the best candidate based upon the relatively short implanta- 
tion times of this study. 



CHAPTER V 
CONCLUSIONS AND SUGGESTIONS FOR FUTURE WORK 

The objectives of this study fall into two categories. 
The first has been an effort to understand the influence of 
compositional variations on the surface chemical behavior of 
a series of bioglasses in a simulated physiologic environment, 
and the relation of this behavior to that exhibited when 
identical glasses are implanted in animals. The second objec- 
tive has been an attempt by the author to bridge the gap 
between the fields of materials science and the biological 
sciences so that an intelligent and practical approach may 
be developed for the selection of a material for potential 
use as a prosthetic device. This has involved developing an 
awareness of problems associated with the body's response to 
prosthetic devices and some of the procedures which are em- 
ployed to examine normal and abnormal responses to foreign 
devices . 

The results of Chapter II have shown that the glasses 
investigated develop a corrosion layer or layers in response 
to attack by an aqueous solution buffered at a pH of 7.4 and 
maintained at 37°C. Sodium and calcium are preferentially 
leached from the soda- lime-silica glass (45S-0I P^Op) , pro- 
ducing a silica-rich film which serves as a buffer zone 



160 



• 161 

protecting the remaining bulk glass from aqueous attack. As 
phosphorus is added to the glass composition, a second film 
is generated at the silica-rich film-water interface. The 
second film is an amorphous calcium phosphate compound which 
crystallizes to an apatite structure with time. Increasing 
the phosphorus content of the glass reduces the time required 
for the calcium phosphate film to form. Partial substitution 
of B„0_ for SiO„ leads to weakening of the silicate network 
and acceleration of the initial dissolution process. Fluorine 
additions significantly enhance the resistance of the glass to 
aqueous attack, probably by substituting for hydroxyl ions in 
the apatite structure of the corrosion film. 

The results of Chapter III confirm the observations of 
Chapter II by providing chemical profiles of the corrosion 
films which define the silica-rich layer and the calcium phos- 
phate layer. The thickness of the calcium phosphate layer 
was found to increase as the phosphorus content of the bulk 
composition increased when glasses were corroded under iden- 
tical conditions. The application of Auger spectroscopy and 
ion beam milling to obtain detailed maps of compositional 
changes over a depth of several micrometers has turned out 
to be a valuable technique in characterizing the corrosion 
behavior of the bioglasses. It should also be noted that 
the results obtained with Auger spectroscopy have substan- 
tiated the usefulness of the techniques employed in Chapter 
II such as infrared reflection spectroscopy and ion solution 
analysis, which infer rather than directly measure information 



162 

about the corrosion films and which are somewhat easier to 
apply to a large number of samples. 

Additional results in Chapter III point to the existence 
of a thin surface layer (10-15 *) rich in calcium and phos- 
phorus which forms during the initial 15 minutes of corrosion. 
The observed sequence of events indicate that the thin cal- 
cium phosphate layer initiates the formation of the silica- 
rich layer and serves as the nucleation site for growth of 
the calcium phosphate layer once sufficient calcium and phos- 
phorus have been leached into solution. 

Based upon the results of Chapters II and III a mechanism 
which explains the formation of multiple corrosion layers 
has been proposed (see page 126) which includes phase 
separation induced by phosphorus. Future work should include 
an investigation of the influence of phosphorus additions on 
the micros tructure of the bioglasses. A new instrument 
ideally suited for such a study is the scanning transmission 
electron microscope [70] with supplemental attachments which 
enable one to obtain elemental analysis and crystallographic 
identification via electron diffraction on a very fine scale. 

The invivo results of Chapter IV have demonstrated that 
the four compositions (see Table 4) implanted all exhibited 
direct attachment to bone. There was a wide variation in the 
appearance of the tissue near the implant. Only the inter- 
face of the glass containing 6 wt.% ^?^c, exhibited a healthy 
zone of ossification characterized by numerous osteocytes in 
close proximity to the glass, a layer of unmineralized 



163 

osteoid, and a layer of osteoblasts actively engaged in lay- 
ing down new osteoid. The other three glasses exhibited a 
lou density of viable osteocytes and an absence of an active 
osteoid front. At 8 weeks this situation had degenerated 
further for the 12 wt.% P^O glass. The osteocytes that were 
present appeared to be dying and osteoblasts were not actively 
producing new osteoid. In addition, the glass had been split 
near the glass-bone interface. This area was filled by a 
capillary containing several types of cells and electron 
microscopy revealed an intercellular crystallization that 
was apparently induced by the excess phosphorus. 

The induction of normal bone growth was related to the 
ability of a bioglass to form a silica-rich corrosion film 
and provide an adequate but not excessive supply of ions to 
be incorporated in the mineralization process. It has not 
been established whether the undesirable results attributed 
to an excess phosphorus concentration are related to the 
amount of phosphorus present or an unbalance produced in the 
ratio of Ca to P. This question could be answered by implant- 
ing a series of bioglasses in which the phosphorus content 
would be held constant while varying the Ca content. It 
would be desirable to analyze the invitro corrosion behavior 
of the same series employing the techniques discussed in 
Chapters II and III. 

The invivo results presented in this study have been 
limited to some type of visual observation of the glass-bone 
interface. The positive results obtained in the invitro 



164 

studies employing Auger spectroscopy to define the corrosion 
profiles (see Chapter III) have opened up the possibility of 
a similar analysis on glass-bone samples. If successful, 
the results would provide a maj:^ of the change in atomic 
composition from the glass through the attachment zone into 
bone. 

Re -examination of the EM grids containing glass-bone 
sections with the scanning transmission electron microscope 
described previously would allow one to achieve interfacial 
compositional and crystallographic identification of the 
interfacial zone of bonding. 

The invitro results presented in Chapter IV describe the 
effect of the conditioning treatment on the surface structure 
of the bioglass implants. Corrosion layers similar to the 
layers produced in the invitro studies of Chapters II and III 
form on the implant surface. 

It is important to know whether the conditioning treat- 
ment is necessary to produce the observed invivo responses. 
Possibly the body would produce the same structural changes 
on the glass surface if unconditioned glasses were implanted. 
In other words, how is the time sequence of events of the 
interfacial reactions influenced by the conditioning treat- 
ment? To answer this question it would be necessary to sub- 
ject the four bioglass compositions employed in Chapter IV to 
an identical implantation experiment eliminating the condi- 
tioning treatment. The results might indicate that a critical 
mixture of Si, Ca, and P ions on the surface is necessary for 



165 

the induction of bone growth. If this were the case, it 
would produce new possibilities for materials for prosthetic 
devices, such as ion impregnation of metals or ceramics with 
the desired amounts of calcium, phosphorus, and silicon. 



BIBLIOGRAPHY 

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BIOGRAPHICAL SKETCH 

Artliur E. Clark, Jr., was born in Savannah, Georgia, in 
1947. He attended hifih school at the American School, Makati , 
Rizal, Philippines. He received a Bachelor of Science degree 
in Metallurgical Engineering in June of 1969. Since obtain- 
ing his bachelor's degree, the author has been pursuing his 
doctorate at the University of Florida. 



171 



I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



L. L. Hench, Chairman 
Professor of Materials Science 
and Engineering 



I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



yv- /• 



J- 




R. T. DeHoff 

Professor of Materials Science 

and Engineering 



I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 




E. D. Verink, Jr. Tf xV 

Professor of Materials'' Science 
and Engineering 



I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 




H. A. Paschall 
Associate Professor of 
Orthopedic Surgery 



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