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dissertation entitled

THE EFFECT OF GRAIN SIZE, MICROCRACKING AND
GRAIN BOUNDARY GROOVING ON OSTEOBLAST
ATTACHMENT IN HYDROXYAPATITE

presented by

IAN ORLAND SMITH

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THE EFFECT OF GRAIN SIZE, MICROCRACKING AND GRAIN BOUNDARY
GROOVING ON OSTEOBLAST ATTACHMENT IN HYDROXYAPATITE

By

Ian Orland Smith

A DISSERTATION

Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of

DOCTOR OF PHILOSOPHY
Department of Chemical Engineering and Materials Science

2007

ABSTRACT

THE EFFECT OF GRAIN SIZE, MICROCRACKING AND GRAIN BOUNDARY
GROOVING ON OSTEOBLAST ATTACHMENT IN HYDROXYAPATITE

By

Ian Orland Smith

This research examined the effect of particle size, microcracking and
grain-boundary grooving in hydroxyapatite (HA) ceramics on osteoblast (OB)
attachment, with the overall goal of understanding the role of physical
characteristics in optimized scaffolds for bone tissue engineering.

Bimodally porous HA scaffolds were fabricated by foaming and sintering
either micron-scale or nano-scale HA powder, yielding two sets with average
grain diameters of 8.6 i 1.9 pm and 588 i 55 nm, respectively. 035 were seeded
onto these scaffolds and counted at 0.5, 1, 2 and 4 hours for attachment and 1, 3
and 5 days for proliferation using a hemacytometer. Results showed that OB
attachment and proliferation was not significantly affected by the change in grain
size and may depend more on the bimodal porosity of the implant. However, as
our attempt to reduce the error in the hemacytometer counts was not fully
successful, a more accurate method of counting the 088, such as a quantifiable

dye, must be used to verify this trend.

While microcracks occur as a result of thermal processing of HA, these
TEA-induced cracks are not easily controlled. For our studies we used Vickers-
induced microcracks to quantify the effect of microcracking on OB attachment in
HA. OB attachment was not significantly affected at one hour, but increased at
four hours to 61% higher than on non-microcracked control specimens. This
increase indicates that microcracking does have an effect on OB attachment and
should be studied further, to assess its effect on OB proliferation and
differentiation. It is not surprising that microcracks have a positive effect on 08
attachment, as this mimics the natural process of bone remodeling. However,
they are not likely to occur in nano-grained HA as a result of processing, as its
small grain size falls below the known values of critical grain size for
microcracking (GCR) in HA.

grain boundary grooving in dense HA is also investigated in this dissertation.
085 were seeded onto grain boundary grooved and control dense HA discs and
counted at one and four hours. At both times, the presence of grain-boundary
grooves did not have a significant effect on OB attachment.

These findings eliminate microcracking and grain boundary grooving as
contributing factors in increased OB attachment on HA substrates with nano-
scale grains. Having ruled out these two phenomena, the observed increase in
08 attachment on nano-grained HA is likely due to either the associated
increase in surface roughness or the presence of nano-scale charge variations at

grain boundaries. The latter is the subject of a current study by our group.

Copyright by
IAN ORLAND SMITH

2007

 

ACKNOWLEDGEMENTS

I would like to thank my parents, Keith and Karen Smith. Their undying
support during my time at Michigan State University made this all possible.

I would also like to thank my academic advisor, Professor Melissa
Baumann, for her guidance and assistance during the pursuit of my studies at
Michigan State University. Her oversight helped me to focus and accomplish
everything that I set out to do. She has helped me to realize my potential in the
realm of engineering research.

Additionally, I would like to thank Professor Eldon Case for all of his help.
By sharing his wealth of knowledge, he has truly enriched my learning

expenence.

TABLE OF CONTENTS

LIST OF TABLES ................................................................................. 11
LIST OF FIGURES ............................................................................... 12
CHAPTER 1 ......................................................................................... 1
INTRODUCTION TO CALCIUM PHOSPHATE BONE TISSUE ENGINEERING
Bone structure .............................................................................. 1
Need for bone replacement ............................................................. 4
Calcium phosphate biomaterials ....................................................... 6
Osteoblast response to hydroxyapatite .............................................. 9
Microcracking ............................................................................. 1 1
Porous hydroxyapatite for bone tissue engineering ............................. 12
Overall aims ............................................................................... 13
References ................................................................................ 1 4
CHAPTER 2 ....................................................................................... 22
EXPERIMENTAL PROTOCOLS
I. Overview .......................................................................................... 22
ll. MC3T3-E1 osteoblast culture .............................................................. 22
A. Freezing cell stock ................................................................... 23
B. Thawing from stock .................................................................. 24
C. Splitting/feeding ...................................................................... 25
Ill. Specimen fabrication ........................................................................ 26
A. Porous hydroxyapatite scaffolds ................................................. 26
B. Dense hydroxyapatite ............................................................... 31
IV. Cell culture/counting ......................................................................... 35
A. Porous hydroxyapatite .............................................................. 36
B. Dense hydroxyapatite ............................................................... 39
References ......................................................................................... 42
CHAPTER 3 ....................................................................................... 43

MC3T3—E1 OSTEOBLAST ATTACHMENT AND PROLIFERATION ON
POROUS HYDROXYAPATITE SCAFFOLDS FABRICATED WITH
NANOPHASE POWDER

Abstract .................................................................................... 43
Introduction ................................................................................ 43
Materials and Methods ................................................................. 47
Results ..................................................................................... 49
Discussion ................................................................................. 53
Acknowledgements ..................................................................... 57
References ................................................................................ 58

vi

CHAPTER4 ....................................................................................... 62
A BRIEF COMMUNICATION REGARDING THE INACCURACIES OF
HEMACYTOMETER MEASUREMENTS WHEN ASSESSING OSTEOBLAST
ATTACHMENT ON POROUS NANO- AND MICRON-SCALE
HYDROXYAPATITE SCAFFOLDS

Introduction ................................................................................ 62
Materials and Methods ................................................................. 63
Results and Discussion ................................................................ 67
Conclusions ............................................................................... 69
References ................................................................................ 71
CHAPTER 5 ........................................................................................ 73

MICROCRACKING AND POROSITY IN CALCIUM PHOSPHATES AND THE
IMPLICATIONS FOR BONE TISSUE ENGINEERING

Abstract .................................................................................... 73
Introduction ................................................................................ 74
Background ................................................................................ 75
Microcracking and critical grain size ....................................... 75
Microcracking in bone ......................................................... 77

HA in bone tissue engineering .............................................. 79
Experimental procedure ................................................................ 79
Results and discussion ................................................................. 81
Microstructure of calcium phosphate specimens ....................... 81
Microcracking in HA specimens ............................................. 83
A re-evaluation of key results from the literature on microcracking in

HA ..................................................................................................... 87
Conclusions and summary ............................................................ 94
Acknowledgments ....................................................................... 96
References ................................................................................ 97
CHAPTER 6 ...................................................................................... 101

INDUCED MICROCRACKING AFFECTS OSTEOBLAST ATTACHMENT ON
HYDROXYAPATITE SCAFFOLDS FOR TISSUE ENGINEERING

Abstract ................................................................................... 101
Introduction .............................................................................. 103
Microcracking in bone ........................................................ 103
Microcracking in hyd roxyapatite ........................................... 104
Materials and methods ................................................................ 106
Powder processing ........................................................... 106
Microcrack induction .......................................................... 107
MC3T3-E1 culture ............................................................. 108
OB attachment study ......................................................... 109
Results and discussion ............................................................... 110
Conclusion ............................................................................... 1 14
References ............................................................................... 1 15

vii

CHAPTER 7 ...................................................................................... 1 19
EARLY STAGE MC3T3-E1 OSTEOBLAST ATTACHMENT ON
MICROCRACKED HYDROXYAPATITE SCAFFOLDS FOR TISSUE

ENGINEERING
Abstract ................................................................................... 1 19
Introduction .............................................................................. 120
Materials and methods ................................................................ 124
Results and discussion ............................................................... 127
Conclusions .............................................................................. 132
References .............................................................................. 134

CHAPTER 8 ...................................................................................... 137

THE EFFECT OF GRAIN BOUNDARY GROOVING ON MC3T3—E1
OSTEOBLAST ATTACHMENT ON DENSE HYDROXYAPATITE

Introduction .............................................................................. 137
Materials and methods ............................................................... 138
Specimen fabrication ......................................................... 138
Cell culture ...................................................................... 140
Cell imaging .................................................................... 141
Results and discussion ............................................................... 143
References .............................................................................. 147
CHAPTER 9 ....................................................................................... 148
SUMMARY AND CONCLUSIONS
Summary of work completed for this dissertation .............................. 148
Conclusions .............................................................................. 162
References .............................................................................. 165
CHAPTER 10 .................................................................................... 167
FUTURE WORK
Surface charge variation and OB attachment on HA ......................... 168
Future work in microcracking and OB attachment on HA ................... 169
References ............................................................................... 173
APPENDICES
APPENDIX A .................................................................................... 176

THREE DIMENSIONAL MICROSTRUCTURAL CHARACTERIZATION OF
POROUS HYDROXYAPATITE USING CONFOCAL LASER SCANNING

MICROSCOPY (CLSM)
Abstract ................................................................................... 176
Introduction .............................................................................. 1 77
Background .............................................................................. 1 79
Porous HA ceramics: biological and mechanical implications of
porosity ............................................................................................ 179
Biological implications of porosity ......................................... 179

viii

Impact of porosity on mechanical properties of bone scaffolds. . .180

Using the CLSM to characterize porous HA specimens ............ 180
Experimental procedure .............................................................. 183
Materials and specimen preparation ..................................... 183
CLSM observation ............................................................ 187
Scanning electron microscope comparison ............................ 188
Results and discussion ............................................................... 189
CLSM characterization of HA ceramic bone scaffold specimens.192
Summary and conclusions ........................................................... 202
Acknowledgements .................................................................... 203
References .............................................................................. 204
APPENDIX B ..................................................................................... 209

CONFOCAL LASER SCANNING MICROSCOPY (CLSM) AS A TOOL FOR
IMAGING CANCELLOUS BONE

Abstract ................................................................................... 209
Introduction .............................................................................. 210
Background .............................................................................. 21 1
Bone structure ................................................................. 211
Imaging porous structures using CLSM ................................. 212
Electron microscopy for imaging bone ................................... 213
Materials and Methods ............................................................... 214
Sample preparation ........................................................... 214

CLSM imaging ................................................................. 215
SEM imaging ................................................................... 216
Results .................................................................................... 216
Reflection mode ............................................................... 216
Topographical fluorescence ................................................ 218

SEM comparison .............................................................. 221
Discussion ............................................................................... 224
Conclusion ............................................................................... 227
References .............................................................................. 228
APPENDIX C .................................................................................... 233

SAMPLE CALCULATION FOR PERCENT OF OSTEOBLASTS ATTACHED TO
DENSE HYDROXYAPATITE SPECIMEN (CONTROL vs. MICROCRACKED, 4
HR ATTACHMENT)

APPENDIX D .................................................................................... 236
X-RAY DIFFRACTION OF DENSE HYDROXYAPATITE USED IN THIS STUDY

APPENDIXE .................................................................................... 237
ELECTROSTATIC INTERACTIONS AS A PREDICTOR FOR OSTEOBLAST
ATTACHMENT TO BIOMATERIALS
Synopsis ................................................................................. 237
Introduction .............................................................................. 238

ix

Materials and methods ............................................................... 240

Suspension preparation ..................................................... 240
Metals ................................................................... 240

Bone ..................................................................... 241

Osteoblasts ............................................................ 241

Bioceramics ............................................................ 242

Results .................................................................................... 243
Metal alloys ..................................................................... 243
Bioceramics ..................................................................... 246
Discussion ............................................................................... 248
Metal alloys ..................................................................... 248
Bioceramics ..................................................................... 250
Acknowledgements .................................................................... 251
References .............................................................................. 253
APPENDIX F ..................................................................................... 255

SURFACE POTENTIAL AND OSTEOBLAST ATTRACTION TO CALCIUM
PHOSPHATE COMPOUNDS IS AFFECTED BY SLECTED ALKALINE
HYDROLYSIS PROCESSING

Abstract ................................................................................... 255
Introduction .............................................................................. 256
Materials and Methods ............................................................... 259
Powder processing methods ............................................... 259
Powder analysis ............................................................... 260
Suspension Preparation ..................................................... 260

Bone ..................................................................... 260

Osteoblasts ............................................................ 261

Bioceramics ............................................................ 261

Results and discussion ............................................................... 262
Powder analysis ............................................................... 262
C—potential ....................................................................... 263
Stability .......................................................................... 265
Acknowledgements .................................................................... 271
References .............................................................................. 272

LIST OF TABLES

Table 1 ............................................................................................. 233

Cell counts taken from 10 fields of view per specimen (9 control
specimens, 9 microcracked specimens)

Table 2 ............................................................................................. 234

Average percent attached osteoblasts for both the control and
microcracked group

xi

LIST OF FIGURES

CHAPTER 1

Figure 1 ............................................................................................... 2
Schematic of the Haversian systems within cortical bone tissue

Figure 2 ............................................................................................... 3
Histological section showing trabeculae within cancellous bone tissue

CHAPTER 2

Figure 1 ............................................................................................. 26

Schematic detailing the cell splitting process. Media is removed by
aspiration (1) and the cells are rinsed twice with 1x PBS. Trypsin/EDTA and
complete media are added in a 1:1 ratio (2) and each plate is rocked back and
forth (3) to aid in cell detachment. Detached cells are pipetted away and placed
into a 50 ml centrifuge tube, along with the cells from other plates (4). 10 pl of
this suspension is placed into each of the two sides of a hemacytometer (5).
Cells are then viewed under the optical microscope (6) and the cells counted (7).
Using this cell population data, the remaining cell suspension is divided up
between fresh plates (60,000 - 75,000 per plate) and 10 ml fresh complete
media added to each (8).

Figure 2 ............................................................................................. 28

Silicone mold used for the foaming and casting of cylindrical porous HA
scaffolds. Both pieces are shown, including the main piece with its cylindrical
mold chambers and the solid base piece.

Figure 3 ............................................................................................. 29

Figure 3A is a view of the Buehler Isomet 1000 wafering saw used to
section the porous cylinders. Figure 3B is a close-up of the custom-built chuck
used to hold each cylinder during cutting.

Figure 4 ............................................................................................. 33

Schematic of indentation/induced microcracking and viewing/measuring of
microcracks using optical microscopy. Indents selected for measurement ("X"
pattern) are indicated by circles.

Figure 5 ............................................................................................. 33

The image at left is a schematic of a typical Vickers indentation. The
length “2a” is the width of the indentation pyramid and “2c” is the length of a
microcrack. The image at right is an SEM micrograph of a single indentation on

xii

the surface of a dense sintered HA specimen. The black arrows indicate
microcracks. Figure taken from Smith et al.3

Figure 6 ............................................................................................. 34

SEM micrograph taken at the surface of a sintered dense HA specimen
that has been thermally etched at 1100°C (1373 K) for 45 minutes to induce grain
boundary grooving.

Figure 7 ............................................................................................. 35
Schematic of a profilometry scan across the specimen surface.

Figure 8 ............................................................................................. 37

The cell scraper used to morselize porous HA scaffolds prior to counting
attached 035. The handle end, located at the left end of the photo, was placed
into each well and the specimen morselized.

Figure 9 ............................................................................................. 38

Schematic showing the process of counting cells on porous HA scaffolds.
Cells are removed from their tissue culture plates (1) and seeded onto sterile
porous HA scaffolds, placed inside of Al foil collars (2). Each collared scaffold is
placed into a 6-well polystyrene tissue culture plate and into the incubator for the
prescribed time. After incubation, each specimen is removed from its collar and
placed into a fresh 6-well plate, rinsed twice with 1x PBS, exposed to
Trypsin/EDTA and the scaffold morselized with the handle end of a cell scraper

(3). 10 ul of the resulting suspension is placed into each of the two sides of the
hemacytometer (4) and the cells (with HA particles) are viewed under the optical
microscope (6) and counted (7).

Figure 10 ............................................................................................ 41

Schematic showing the process of fabricating, microcracking, cell seeding,
staining and counting on dense HA. Powder is uniaxially pressed, sintered and
polished (1) prior to induced microcracking (2). Microcrack size is determined,
the specimens sterilized and 085 are seeded onto their surfaces (3). Following
culture times of one and four hours, the 08s are stained with
Rhodamine/Phalloidin and Hoescht (4) and viewed under an optical fluorescence
microscope (5). Fields of view are chosen as shown (grey boxes each indicate
one field of view) (6).

CHAPTER 3
Figure 1 ............................................................................................. 50
Comparison of SEM micrographs of the surface of scaffolds fabricated

using micro HA (left) and nano HA (right).

Figure 2 ............................................................................................. 51

xiii

SEM micrograph showing the surface of porous scaffold fabricated using
nano HA, including macropores (large arrows noting approximate boundary of a
macropore) and micropores (small arrows emphasizing the <60 pm
interconnecting pores).

Figure 3 ............................................................................................. 52

MC3T3-E1 OB attachment on micro HA (dashed line) and nano HA (solid
line) scaffolds as a function of time, at intervals of 0.5, 1, 2 and 4 hours. Values
are mean i SE, p < 0.05 at all time intervals for comparison of nano HA to micro
HA.

Figure 4 ............................................................................................. 53
MC3T3-E1 Osteoblast proliferation on micro HA and nano HA, after

periods of 1, 3 and 5 days. Values are mean t SE. p > 0.05 for comparison of

nano HA to micro HA porous scaffold for 1 and 3 days. p < 0.05 for 5 days (1:).

CHAPTER 4

Figure 1 ............................................................................................. 66

Figure 1A is an optical micrograph of a hemacytometer with a single OB
(indicated by circle) taken during a cell count where there were no ground HA
particles. Figure 1B is an optical micrograph of a hemacytometer with a
morselized OB-seeded micron-HA scaffold. Arrow 1 indicates an HA particle,
arrow 2 indicates an OB and arrow 3 indicates an object that is not definitively
one or the other. The general haze resulting from the ground particles, as well as
the presence of larger OB-sized particles make it difficult to identify OBs.

Figure 2 ............................................................................................. 68
MC3T3—E1 OB attachment (0.5, 1, 2 and 4 hrs) and proliferation (24, 72
and 120 hrs) on porous nano-HA and micron-HA scaffolds.

CHAPTER 5

Figure 1 ............................................................................................. 83
SEM micrograph of polished and themial etched hydroxyapatite (5000x).

Figure 2 ............................................................................................. 84
SEM micrograph of polished and thermal etched fi-TCP (3000x).

Figure 3 ............................................................................................. 84
SEM micrograph of polished and thermal etched BCP (3000x).

Figure 4 ............................................................................................. 85

SEM micrograph of as-sintered hydroxyapatite specimen surface. Image
was taken away from the fracture surface (6500x).

xiv

Figure 5 ............................................................................................. 86

SEM micrograph of as-sintered hydroxyapatite specimen surface shown in
Fig. 5. Image was taken away from the fracture surface (4500x) and at different
location on the specimen.

Figure 6 ............................................................................................. 86
SEM micrograph of as-sintered hydroxyapatite specimen surface. Image
was taken away from the fracture surface (2500x).

Figure 7 ............................................................................................. 90
Schematic showing the relationship between relative density and grain
size for ceramic powders.

CHAPTER 6

Figure 1 ............................................................................................ 108

SEM micrograph of microindentation on HA specimen surface (right). "2a"
indicates the diagonal length of the indentation impression and "20" indicates
crack length (left).

Figure 2 ............................................................................................ 1 1 1
Osteoblast attachment at 4 hours on microcracked and non microcracked
dense HA discs

Figure 3 ............................................................................................ 113

3A shows MC3T3-E1 osteoblasts with both spindle-shaped and polygonal
morphology on HA control specimen, stained to highlight actin fibers (red —
Rhodamine-Phalloidin) and nuclei (blue — Hoechst). 3B shows MC3T3-E1
Osteoblasts with spindle shaped morphology on HA microcracked specimen,
stained to highlight actin fibers (red — Rhodamine-Phalloidin) and nuclei (blue —
Hoechst).

CHAPTER 7

Figure 1 ............................................................................................ 123

SEM micrograph of microindentation on HA specimen surface. “R"
indicates radial cracks, the distance "23" indicates the diagonal length of the
indentation impression and the distance "2c" indicates crack length.

Figure 2 ............................................................................................ 128

A) MC3T3—E1 OB attachment at 1 hour on control and microcracked HA
specimens. B) MC3T3—E1 OB attachment at 4 hours on control and microcracked
HA specimens. 4-hour attachment data was taken with permission from previous
work by Smith et al.1 # indicates that difference in OB attachment between
control HA and microcracked HA at 4 hours is significant (p = 0.004).

XV

Figure 3 ......................................................................................... 131

A) MC3T3—E1 088 with spread (white circle) and polygonal (black circle)
morphology on HA control specimen. B) MC3T3-E1 083 with similar spread
(white circle) and polygonal (black circle) morphology on microcracked HA
specimen. 083 have been stained to highlight actin fibers (red — Rhodamine-
Phalloidin) and cell nuclei (blue - Hoechst).

CHAPTER 8

Figure1 ............................................................................................ 140
SEM micrograph of the surface of a grain boundary grooved HA
specimen.

Figure 2 ............................................................................................ 143
MC3T3-E1 OB attachment at 1 (Figure 2A) and 4 (Figure 28) hours on
thermally etched and control HA specimens.

Figure 3 ............................................................................................ 145
This set of four optical micrographs shows attached 083 at one hour on
the surface of a control specimen (Figure 3A) and a grain boundary grooved
specimen (Figure 3B), and also at four hours on the surface of a control
specimen (Figure 3C) and a grain boundary grooved specimen (Figure 3D).

CHAPTER 9

Figure 1 ............................................................................................ 149

Comparison of SEM micrographs of the surface of bimodally porous
scaffolds fabricated using micro HA (left) and nano HA (right). Scaffolds were
fabricated by foaming a suspension of HA powder using H202 as a foaming
agent, drying at 125°C for 1 hour and sintering either at 1360°C for 4 hours
(micro HA) or 1100°C for 1 hour (nano HA). Image taken from Smith et al., Int J
Nanomed 2006;1 2189-194.

Figure 2 ............................................................................................ 149

SEM micrograph showing the surface of porous scaffold fabricated using
nano HA, including macropores (large arrows noting approximate boundary of a
macropore) and micropores (small arrows emphasizing the <60 um
interconnecting pores). Image taken from Smith et al., Int J Nanomed
2006;1:189-194.

Figure 3 ............................................................................................ 151

MC3T3-E1 OB attachment on micro HA (dashed line) and nano HA (solid
line) scaffolds as a function of time. 085 were seeded onto porous scaffolds (n =
3 with each trial run in triplicate) at a density of 11,320 cells/cm2 and counted at
intervals of 0.5, 1, 2 and 4 hours. Counts were completed after trypsinizing,
morselizing and using a hemacytometer. Values are mean i SE, p < 0.05 at all

xvi

 

time intervals for comparison of nano HA to micro HA. Figure taken from Smith et
al., Int J Nanomed 2006;1:189-194.

Figure 4 ............................................................................................ 152

MC3T3-E1 Osteoblast proliferation on micro HA and nano HA. 083 were
seeded onto porous scaffolds (n = 3 with each trial run in triplicate) at a density of
11,320 cells/cm2 and counted after periods of 1, 3 and 5 days. Counts were
completed after trypsinizing, morselizing and using a hemacytometer. Values are
mean i SE. p > 0.05 for comparison of nano HA to micro HA porous scaffold for
1 and 3 days. p < 0.05 for 5 days (1). Figure taken from Smith et al., Int J
Nanomed 2006;1:189-194.

Figure 5 ............................................................................................ 153

MCST3-E1 OB attachment (0.5, 1, 2 and 4 hrs) and proliferation (24, 72
and 120 hrs) on porous nano-HA and micron-HA scaffolds. Again, 083 were
seeded onto porous scaffolds (n = 3 with each trial run in triplicate) at a density of
11,320 cells/cm2 and counted at intervals of 0.5, 1, 2 and 4 hours. Counts were
completed after trypsinizing, morselizing and using a hemacytometer.
Background counts were completed using scaffolds morselized in the absence of
08s and subtracted from the raw data (dotted lines), to yield adjusted values
(solid lines). Figure taken from Chapter 4.

Figure 6 ............................................................................................ 154
SEM micrograph of an as-sintered hydroxyapatite specimen surface.
Image was taken away from the fracture surface (4500x). This micrograph shows
multiple examples of microcracking due to thermal expansion in HA, indicated by
arrows. White arrows indicate cracks crossing the grain bulk and black arrows
indicate cracks along grain boundaries. This type of microcracking is not easily
controlled. Image taken from Case et al., Mater Sci Eng A 2005;390:246-254.

Figure 7 ............................................................................................ 156

SEM micrograph of microindentation on HA specimen surface (right). "2a"
indicates the diagonal length of the indentation impression and "20" indicates
crack length (left). This indent was created using a Vickers microindentor with a
load of 4.9 N, a load time of 5 secnds and a loading rate of 70 pm/sec. This
method of microcrack induction created microcracks of an average size of ZCAVE
= 124.4 pm i 21.1 pm and is more easily controlled compared to TEA-induced
microcracking. Indents such as this one were placed in a 14x14 grid across the
specimen surface. Image taken from Smith et al., Am J Biochem Biotechnol
2006;2z105-110.

Figure 8 ............................................................................................ 157

Osteoblast attachment at 1 and 4 hours on microcracked and non-
microcracked dense HA discs. 085 were seeded onto dense microcracked and
control HA scaffolds (n = 3 with each trial run in triplicate) at a density of 11,320
cells/cm2 and counted at 1 and 4 hours. Counts were completed after fixing the

xvii

cells using formaldehyde and staining with both Rhodamine/Phalloidin and
Hoescht. Stained cells were viewed under an optical fluorescence microscope
and counted in 10 fields of view. Counts were then averaged for each group of 9
specimens (control and microcracked) and the results shown here, with the error
bars indicating standard deviation (as opposed to SE, used in Figures 3 and 4).
“#” indicates that difference in OB attachment between control HA and
microcracked HA at 4 hours is significant (p = 0.004). Figure taken from Smith et
al., submitted to J Biomed Mater Res, December 2006.

Figure9 ............................................................................................ 158

A) MC3T3-E1 085 with spread (white circle) and polygonal (black circle)
morphology on HA control specimen. B) MC3T3—E1 083 with similar spread
(white circle) and polygonal (black Circle) morphology on microcracked HA
specimen. 085 have been fixed and stained to highlight actin fibers (red —
Rhodamine-Phalloidin) and cell nuclei (blue — Hoechst). Images taken from
Smith et al., submitted to J Biomed Mater Res, December 2006.

Figure 10 ......................................................................................... 158

10A shows MC3T3-E1 osteoblasts with both spindle-shaped and
polygonal morphology on HA control specimen after 4 hours culture. OBs were
fixed and stained to highlight actin fibers (red — Rhodamine-Phalloidin) and nuclei
(blue - Hoechst). 10B shows MC3T3-E1 Osteoblasts with spindle shaped
morphology on HA microcracked specimen at 4 hours, again stained to highlight
actin fibers (red — Rhodamine-Phalloidin) and nuclei (blue — Hoechst). Images
were taken using an Optical fluorescence microscope. Images taken from Smith
et al. J Am Biochem Biotechnol 2006;2z105-110.

Figure 11 ......................................................................................... 160

SEM micrograph of the surface of a grain boundary grooved HA
specimen. Grain boundary grooving was accomplished by thermally etching
sintered dense HA through heating at 1100°C for 45 minutes. Grain boundary
grooves shown here are estimated to be approximately 0.5 pm in width (hillock to
hillock). Image taken from Chapter 8.

Figure 12 ......................................................................................... 161

MC3T3-E1 OB attachment at 1 (Figure 2A) and 4 (Figure 2B) hours on
thermally etched and control HA specimens. 083 were seeded onto dense grain
boundary grooved and control HA scaffolds (n = 3 with each trial run in triplicate)
at a density of 11,320 cells/cm2 and counted at 1 and 4 hours. Counts were
completed after fixing the cells using formaldehyde and staining with both
Rhodamine/Phalloidin and Hoescht. Stained cells were viewed under an optical
fluorescence microscope and counted in 20 fields of view. Counts were then
averaged for each group of 9 specimens (control and microcracked) and the
results shown here, with the error bars indicating standard deviation (as opposed
to SE, used in Figures 3 and 4). At both time intervals there is no significant
difference between groups. Figure taken from Chapter 8.

xviii

Figure 13 ......................................................................................... 162

This set of four optical micrographs shows attached 083 at one hour on
the surface of a control specimen (Figure 3A) and a grain boundary grooved
specimen (Figure 3B), and also at four hours on the surface of a control
specimen (Figure 3C) and a grain boundary grooved specimen (Figure 30). 085
were fixed and stained to highlight actin fibers (red — Rhodamine-Phalloidin) and
nuclei (blue — Hoechst). The lesser degree of OB spread in both groups at 1 hour
is a result of the shorter culture time. Images were taken using an optical
fluorescence microscope. Images taken from Chapter 8.

APPENDIX A

Figure 1 ............................................................................................ 193

A series of four reflection-mode confocal laser scanning microscopy
images taken at a single location on hydroxyapatite (HA) specimen HA-70-AS-1:
(a) an overlay image showing both macropores (8150—300 mm) and micropores
(860—70 mm), (b) a Phi-Z scan along “cut line” A—B, with pore dimensions
labeled directly on the profile, (c) a “wire mesh” topographic three dimensional
profile highlighting the larger surface features of the specimen while supressing
the fine surface detail, and (d) a contour map that uses a gray scale to indicate
depth relative to a basal plane.

Figure 2 ............................................................................................ 194

A series of reflection-mode confocal laser scanning microscopy images of
specimen HA-70-AS-2, taken at a single location on the specimen: (a) an overlay
image, (b) a Phi-Z scan

Figure 3 ......................................................................................... 195

For a single location on the epoxy-impregnated and stained specimen HA-
70-ES-1, a series of fluorescent-mode confocal laser scanning microscopy
images: (a) an overlay image, (b) a Phi-Z profiles along the cuts “1” and cut “,"2
where the bright regions correspond to the stained epoxy that fluoresces and the
dark regions hydroxyapatite (HA), which does not fluoresce, and (c) a single 2
section image taken at a depth of about 40 mm beneath the specimen surface. A
pore throat (opening between two adjacent pores) is labeled on the image.

Figure 4 ............................................................................................ 196

A comparison between (a) an overlay confocal laser scanning microscopy
(CLSM) image and (b) a conventional scanning electron microscoope (SEM)
image taken at the same location with identical magnification. Although the SEM
image has higher lateral resolution than the CLSM image, the SEM is unable to
simultaneously image the specimen surface and within the pores because of the
limited depth of field of the SEM.

Figure 5 ............................................................................................ 197

xix

A series of reflection-mode confocal laser scanning microscopy (CLSM)
images of specimen HA-70-AS-2 taken with a higher lateral and vertical
resolution than the CLSM images in Figs. 1—5, including a: (a) Phi-Z scan and (b)
topographical profile. All images were taken at the same location on the
specimen.

APPENDIX B

Figure 1 ............................................................................................ 217
Four images taken at a single location on cancellous bone specimen,
including (Fig. 1a) a Z-stack overlay image showing macropores and
interconnecting micropores (indicated with arrows), (Fig. 1b) a Phi-Z scan cross-
sectional cut taken along the line shown, with A-B-C shown in the upper portion
indicating the features of a pore (A-B-C on the lower portion), and C-B equal to
1.2 mm, (Fig. 1c) topographical map consisting of a wire mesh of consecutive
Phi-Z cuts and (Fig. 1d) a contour map showing pore morphology and depth.

Figure2 ............................................................................................ 219
Local staining of cancellous bone imaging (a) the interior of a single
macro- and (b) micropores along the surface of a macropore.

Figure 3 ............................................................................................ 221

CLSM reflectance images showing a single micropore on the surface of a
macropore for globally stained cancellous bone. Figures 3b and 30 were used in
measuring pore size.

Figure 4 ............................................................................................ 223

A comparison between (a) CLSM image and (b) SEM image of the same
section of defatted and deproteinized cancellous bone. Fig. 4a was taken using
reflective CLSM, showing greater detail of pore depth, while Fig. 4b highlights the
limitations of SEM imaging because of lower depth of field.

APPENDIX D

Figure 1 ............................................................................................ 236
X-ray diffraction data of dense sintered hydroxyapatite (Taihei Chemical).

APPENDIX E

Figure 1 ............................................................................................ 244

Q-potential as a function of pH for as-received 316L, CoCrMo and Ti6Al4V
vs. whole bone suspended at 1% in PS.

Figure 2 ............................................................................................ 244
Q-potential as a function of pH for ground whole deer bone suspended at
1% in PS.

XX

Figure 3 ............................................................................................ 245

Q-potential as a function of pH for NaOH-washed 316L, CoCrMo and
Ti6Al4V vs. whole bone suspended at 1% in PS.

Figure 4 ............................................................................................ 246
Q-potential as a function of pH for as-received 316L, CoCrMo, Ti6Al4V and
whole bone suspended at 1% in PS with 1% (by volume) bovine serum albumin.

Fig ure 5 ............................................................................................ 246

q-potential as a function of pH for NaOH-washed 316L, CoCrMo, Ti6Al4V
and whole bone suspended at 1% in PS with 1% (by volume) bovine serum
albumin.

Figure 6 ............................................................................................ 247

q-potential as a function of pH for HA (sintered), HA (unsintered), C-TCP,
osteoblasts and whole bone suspended at 1% in PS.

Figure 7 ............................................................................................ 248

Q-potential as a function of pH for HA (sintered), HA (unsintered) and B-
TCP in 1% (by volume) bovine serum albumin in comparison to osteoblasts and
whole bone suspended at 1% in PS.

APPENDIX F

Figure 1 ............................................................................................ 266
q-potential for CDA(NaOH), CDA(NH4OH), BCP(NaOH), and
BCP(NH4OH) (2A). Q-potential for MC3T3-E1 mouse osteoblasts and ground
bone (28).Both figure 2A and 28 are shown as a function of physiologically
relevant pH at solids loadings of ~1 volume % in 0.154 M NaCl electrolyte.

Figure 2 ............................................................................................ 267
Stability calculations for CDA(NaOH), denoted W11, bone (W22) and the
interaction between CDA(NaOH) and bone (W12).

Figure 3 ............................................................................................ 267
Stability calculations for CDA(NH4OH), denoted W11, bone (W22) and the
interaction between CDA(NH4OH) and bone (W12).

Figure 4 ............................................................................................ 268
Stability calculations for BCP(NaOH), denoted W11, bone (W22) and the
interaction between BCP(NaOH) and bone (W12).

Figure 5 ............................................................................................ 268

Stability calculations for BCP(NH4OH), denoted W11, bone (W22) and the
interaction between BCP(NH4OH) and bone (W12).

xxi

CHAPTER 1

INTRODUCTION TO CALCIUM PHOSPHATE BONE TISSUE ENGINEERING

This research examined the effect of particle size, microcracking and
grain-boundary grooving in hydroxyapatite (HA) ceramics on osteoblast (OB)
attachment, with the overall goal of understanding the role of physical
characteristics in optimized scaffolds for bone tissue engineering. Images in this

dissertation are presented in color.

BONE STRUCTURE

Structurally, bone is composed of two primary types of tissue.1'2 On one
hand there is cortical bone, also known as compact bone, found prominently
around the exterior of the diaphysial regions of long bones. Cortical bone is
composed of a network of osteons, also known as Haversian systems, made up
of thin lamellae of bone tissue surrounding channels called Haversian canals
(Figure 1). These Haversian canals act as a conduit for blood vessels and nerve
fibers and travel parallel to the axis of loading in the bone. Within the lamellae
are cavities called lacuna, which are home to osteocytes (inactive osteoblast
cells). These cells exchange nutrients with the rest of the bone tissue via small
channels called canaliculi, which contain extensions of the osteocytes. Near the
outer region of the bone, lamellae spread out to form a concentric ring around the

entire bone, adjacent to the periosteum.

 

 

Figure 1: Schematic of the Haversian systems within cortical bone tissue1

Cancellous, or spongy bone consists of bone-tissue trabeculae and
interconnected pores containing the bone marrow. These trabeculae also consist
of lacunae-containing lamellae, although these lamellae are not oriented as in
cortical bone.1

The interconnected porosity of cancellous bone is vital for the transport of
nutrients via the bone marrow vasculature. Blood vessels mn from the
periosteum, outside of the cortical bone, through the Haversian canals, through
the trabelcular space and into the medullary canal.1 In order to create scaffolds
that are able to connect with the surrounding bone tissue, it is important to

reproduce this interconnected porous network. This will be discussed later.

 

Figure 2: Histological section showing trabeculae within cancellous bone tissue1

Bone remodeling and growth arise as a result of new bone production by
osteoblasts and concurrent resorption of old bone by osteoclasts. These cells are
active as bone is being continually remodeled as a result of normal stresses
applied during everyday use?"4

When an osteoblast is forming new bone, it is deemed an active
osteoblast. These cells act to synthesize type 1 procollagen molecules and
release them into the area of growth, called the osteoid. Within the osteoid,
procollagen molecules align themselves and, over the course of approximately
10 days, calcium phosphate is deposited and eventually matures into bone
apatite. As active osteoblasts are consumed by newly formed bone, the majority
die and a fraction of the remaining living cells become encased within this bone

and become osteocytes.“3

There are a variety of approaches for studying osteoblast (OB) function.
Several cell lines are available for use, including rat osteogenic sarcoma (ROC)
cells such as the UMR 106 line and the newborn mouse calveria-derived MC3T3-
E1 line.1'5’6 The MCBT3-E1 osteoblast clonal cell line is used in many cell
attachment, proliferation and differentiation experiments, in lieu of primary human
or animal 083 due to their greater availability and well-established bone forming
activity7, as well as its ability to differentiate on hydroxyapatite.1 Further, using a
clonal line eliminates the problems that are inherent to primary cell cultures, such
as difficulty in cell isolation.7’10 The use of MC3T3-E1 083 in place of primary
083 in tissue engineering has been shown to be a suitable substitute.11 We used
088 rather than osteoclasts (00$) for all of our studies because we chose to
focus on new bone production, rather than bone resorption that acts as a trigger

for bone production.

NEED FOR BONE REPLACEMENT

Each year in the United States alone there are over 500,000 bone graft
procedures performed as a result of trauma damage and resection as well as
deterioration associated with diseasem’14 Bone tissue graft options include
autograft, allografts and artificial replacement materials.

Autografts require the removal of bone tissue from a donor site within the
patient's body. This technique has inherent complications. Autografts require that
bone be removed at the donor site to replace and repair the damaged or missing

tissue. By causing damage to repair or replace missing or damaged tissue,

autografting is problematic, and therefore cannot involve load-bearing donor
sites. Additionally, retrieval of autograft tissue requires a second surgery at the
donor site, thereby increasing the morbidity/mortality rate.

Allograft tissue on the other hand, originates in a same-species,
predominately cadaveric, donor other than the patient. Cadaveric tissue is in
limited supply and carries the possibility of tissue rejection and disease
transmission.

Artificial bone tissue replacement materials are an option being developed
through an extensive range of research. These bone substitutes fall into three
categories: scaffold-based, cell-based and factor-based. This dissertation
focuses on scaffold-based bone substitutes. An ideal bone replacement scaffold
for filling a critical sized gap in bone tissue would have a bimodal pore structure
similar to that of cancellous bone. This structure allows ingrowth of new tissue
into the larger macropores and nutrient transmission through the micropores.

Interconnected macropores between 200-500 pm in diameter have been shown

to promote new tissue ingrowth and vascularization.1“"16

In addition to having a pore structure conducive to bone ingrowth, an ideal
bone replacement scaffold would also be chemically similar to natural bone.
Consequentially, ceramic materials are involved in approximately 60% of
available bone graft substitutes where, the calcium phosphate (CaP) family of
materials is used in the majority of artificial bone grafts.12 In particular,
hydroxyapatite (HA, Ca10(PO4)6(OH)2) is frequently used, because of its similarity

to the natural apatite bone mineral.67'68'35'1245'”20

One way in which bone substitute materials are being designed to better
mimic bone is through the design of nanoscale ceramic/polymer composites.

These composites consist of a nanoscale CaP ceramic, usually HA, and

21-23 24-26

collagen , a non-collagen polymer such as polyamide or a combination of
the M02728. Some of these composites show improved compressive strength
compared to HA, but the inclusion and eventual degradation of polymers such as
polylactic acid (PLA) has in some cases been shown to have a negative effect on
cell behavior.29 Also, while some of these composite systems contain
components found in living bone, they do not mimic its complex heirarchical
structure, which is compised of apatite crystals, 15 - 200 nm in length and 10 —
80 nm in width with a thickness between 2 and 7 nm.30 These crystals reside
within the collagen fibers, which form lamellar sheets in the concentric rings of
osteons in cortical bone2 and the trabeculae of cancellous bone.31 Therefore, one

current thrust of research is in the development of a bioceramic whose structure

mimics the nano-scale features of bone.32

CALCIUM PHOSPHATE BIOMATERIALS
Calcium phosphate (CaP) ceramics have been frequently studied as

L15 While these materials are not

candidates for hard tissue replacemen
osteoinductive, they are biocompatible and osteoconductive, meaning they elicit

a biologic response in vivo.33 CaP ceramics commonly used in biomedical

engineering include Ct-tricalcium phosphate (Ct-TCP), B-tricalcium phosphate (B-

 

TCP), tetracalcium phosphate (TTCP) and hydroxyapatite (HA).15 Some benefits
of using each will be discussed later in this chapter.

One use for CaP in biomedical engineering is as an osteoconductive
coating for alloys in orthopedic devices and its benefits as a coating for these
metals used as orthopedic implants have been well documented.34‘36

Because of its high degree of similarity to bone apatite, HA
(0310(PO4)6(OH)2) is a frequently investigated material for use as a bone
replacement. As a result of this inherent chemical similarity, HA exhibits
exceptional biocompatibility. Although HA shows a favorable propensity for
attracting 083, its properties as a bone implant are constantly being adjusted for
improvement.37 For example, by changing the processing parameters of the HA
itself, as well as through post-processing surface treatments OB attachment can
be altered.38 These methods and others are discussed later in this chapter.

Along with surface energy and surface composition, the biomaterial
surface charge is a factor which can be used to promote a better bone/implant

interface.39 Osteoblasts have been shown to have a propensity for attaching to

HA, beginning with an electrostatic attraction. 083 in suspension exhibit a large
enough gap in C-potential with respect to HA to exhibit an interaction deemed

unstable (meaning conducive to interparticle flocculation) by the Hogg, Healy and
Fuerstenau model for interparticle stability.“°‘42 Osteoblasts attach to, grow on
and eventually differentiate, laying down mineral matrix upon HA.7'8 083 have
been shown to bind more effectively to HA than other common biomaterials,

including Ti and steel, because of increased fibronectin and vitronectin

adsorption in vivo.43 HA belongs to the group of materials dubbed “bioactive”,
because of their tendency to bond with surrounding bone tissue and to promote
bone formation in vivo. 39"“ However, the use of this term here is incorrect
because bioactivity, strictly speaking, refers to the ability of a material to produce
living matter. This is something that HA cannot accomplish on its own.

The initial reaction between a bioceramic material and its environment is
determined by the rate of dissolution from the ceramic and the precipitation back
onto the ceramic, prior to the attachment of 033 to the specimen surface. This
rate of dissolution varies based on the Ca/P ratio of HA-based materials, and is
slowest in HA (~1:90 vs TOP“). This lack of resorbability shown by HA in vivo
Offers a trade-off though, in that HA implants maintain their mechanical integrity

over time and even result in improved shear strength at the bone/implant

interface compared with Ot- and B-TCP, exhibiting a 10x increase, determined

using the push-out mechanical test.45

Hydroxyapatite, a ceramic, exhibits good compressive but poor tensile
mechanical properties. Dense HA has a Young’s modulus (E) of 40 — 117 GPa, a
compressive strength of 294 MPa and a bending strength of 147 MPa.46 In

comparison, a typical femur possesses a Young’s modulus (E) of 13.5 GPa and
a bending strength of 250 MPa.46 This disparity in strength is what currently limits
the use of HA bone-replacement scaffolds to primarily non-load-bearing
applications.46 In order for a HA scaffold to effectively promote bone ingrowth, it

must be porous, and therefore have an ever further disparity in strength to

CO rtical bone.16

OSTEOBLAST RESPONSE TO HYDROXYAPATITE

The scope of our research within this dissertation is focused mainly on the
interaction between OBs and the HA surface, either within a bimodally porous HA
ceramic network or on the dense HA ceramic substrates. The way in which this
interaction is affected depends on the properties of that surface, those that are
intrinsic in nature and those that are extrinsic. The difference between the two
will become more apparent later in this chapter, but intrinsic properties are those
that are inherent to the particular type of HA being studied such as crystallinity,
grain size and inherent surface roughness, chemical additives during processing
and porosity, and are a result of processing parameters. Extrinsic properties refer
to those which result from further treatment of the as-processed ceramic and
include, for example the surface functionalization of the as-processed HA
substrate.

By varying the crystallinity of HA, the dissolution rate is altered.47 For
example, it was recently shown that nano-crystalline and nano-amorphous HA
elicited improved OB adhesion than conventional crystalline HA and this
response was similar to that found on RGD functionalized conventional HA.47
This is consistent with earlier findings that amorphous materials, such as
Bioglass® and apatite-wollastonite (A-W) glass ceramic better encouraged bone
ingrowth when compared to crystalline HA.19'48 For example, particles of
Bioglass® and A-W glass ceramic brought about bone ingrowth rates 2 and 1.3

tithes that associated with crystalline HA particles, respectively, after 24 weeks in

W'vo (rabbit model). Surface functionalization using RGD-containing peptides is

another example of a treatment which improves OB attachment. Functionalizing
the HA surface with EEEEEEEPRGDT improved OB attachment49 and using
GRGDSPC and cycIo-DfKRG increased attachment at 3 and 24 hours50 versus
HA control. (The nomenclature for these peptides is designed to indicate their
specific sequence of amino acids, with each letter referring to one of these amino
acids.)

The relationship between nano-grained ceramics versus micro-grained
dense bioceramics as related to DB behavior on nano-grained versus micro-
grained dense bioceramics has been investigated by Webster et al. Using dense

1-
t5 53

substrates with nano-scale grains increases OB attachmen and

51,54 51,54

proliferation , alkaline phosphatase (AP) activity and calcium

5"54 and decreases OB motility51'54. For example, OB proliferation on

deposition
nanograined HA (67 nm average grain diameter) versus conventional HA (123
nm average grain diameter) is ~50% higher after 5 days.54 Additionally, AP
activity was ~50% higher and calcium deposition ~100% higher after 28 days in
vitro.54 This work suggests that the increase in surface roughness inherent to the
decreased grain size is a contributing factor to this improvement in OB behavior.
This is consistent with improved OB behavior shown with incidence of surface
microroughing on various biomedical alloys, including Ti alloys‘c’f"56 and 316L
stainless steel,” bioactive glasses“, HA and glass-reinforced HA59 For example,

microroughing induced by chemical etching in three different kinds of bioactive

glass resulted in a ~60-300% increase in DB attachment after 2 hours.58

10

However, Webster et al. does not investigate how nano-grained porous ceramics
affect OB behavior. We address this relationship in Chapters 3 - 4.
Chemical additions to HA during processing include the use of sintering

60,61

aids, such as N33PO4 , and the substitution of silicon ions for phosphate ions

in order to increase dissolution rate62'65.
The inclusion of pores in HA scaffolds is important for successful bone

tissue ingrowth and will be discussed in its own section.

MICROCRACKING

While experiencing the rigors of daily activity, bone undergoes repeated
cyclic strain”67 The repetitive nature of this strain causes an accumulation of
microcracks in the bone tissue.46 The idea that this microdamage results from
fatigue was originally presented by Frost“, who also suggested that bone
microdamage is related to bone remodeling.68 Fatigue microdamage in bone
likely triggers targeted bone remodeling,69 which occurs near microdamagem'74

Microcracks in bone were observed using various methods, including
radiography75 and fuschin or fluorescent dye staining with or without optical
microscopy68' 7678 The dimensions of observed cracks were not always reported,
but in those instances that they were these values varies based on location,
species and the nature of activity undertaken by the patient. For example,

microcracks found in human rib bones were between 296-404 pm in length with a

Crack opening displacement of 88-97 pm.76 Microcracks found in equine subjects

11

had mean lengths between 37-60 pm, depending on the number of races the
horses had undertaken.77

Chapters 5 - 7 of this dissertation addresses microcracking in HA as a
property which may improve OB attachment. Originally observed as occurring

intrinsically, microcracking was later controlled and induced in dense HA.

POROUS HYDROXYAPATITE FOR BONE TISSUE ENGINEERING
While unimodal porosity in HA is controlled by varying parameters such as
the forces applied during cold or hot pressing, sintering time and sintering
temperature, bimodal porosity requires more complex methods of fabrication.
Bimodal porosity in HA is achieved in a number of ways, including foaming (used
by our group”), by burning out a fugitive phase such as polymethyl methacrylate
(PMMA)8°'81, the use of rapid prototyping to build up a periodic matrix, which
does not mimic the natural structure of bone, but rather is comprised of
orthogonal "struts” of HA82, or by removing the organic component of natural
cancellous bone, followed by conversion of the bone mineral into HA.”18
As mentioned above, the inclusion of naturally-occurring pores and canals
within bone is important in the transport of nutrients throughout the bone, as well
as for communication between cells and innervation of the tissue.2 An advantage
of using a porous HA implant is that the interface available for cell attachment is
far greater and more complex, resulting in an improved interlocking interface.83

In comparison, a more dense material would offer a smooth interface with no

available features with which the bone can ingrow.16' 83 Pore sizes above 100 pm

12

lead to bone ingrowth, while sizes between 200-500 pm have been shown to

15.16.83 This range of porosity

optimize the vascularization of the new tissue.
includes interconnected pores, referred to as macropores, coupled with the
interconnections themselves, called micropores. It is this interconnectivity that is
key in promoting bone ingrowth and subsequent vascularization.84 Bimodal
interconnected porosity is detailed in our work with confocal laser scanning

microscopy (CLSM) as a method for imaging cancellous bone and porous HA,

included in Appendix A-B.

OVERALL AIMS

We chose to examine the effect of grain size, microcracking and grain
boundary grooving in order to design a better bone tissue engineering scaffold.
Such a scaffold would be bimodally porous and possess properties that improve
cell behavior and tissue ingrowth.‘ The aim of the first set of studies is to
determine whether the observed relationship between nano-grained dense HA
and improved OB attachment and growth translates into similar behavior in
bimodally porous HA implants. The aim of the studies described in Chapters 5 - 8
is to investigate the role played, if any, by two different surface features on dense
HA: microcracking and grain boundary grooving. Ultimately, should these
features prove beneficial to 08 behavior, they would be included in the design of

an ideal porous scaffold.

13

 

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54.Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R. Enhanced
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55.Silva CC, Pinheiro AG, de Oliveira RS, Goes JC, Aranha N, de Oliveira
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56.Huang, J, Best, SM, Bonfield, W, Brooks RA, Rushton N, Jayasinghe SN,
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58.Itala A, Ylanen HO, Yrjans J, Heino T, Hentunen T, Hupa M, Aro HT.
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59.Prado da Silva MH, Lemos AF, Ferreira JMF, Lopes MA, Santos JD.
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62.Porter AE, Patel N, Skepper JN, Best SM, Bonfield W. Comparison of in
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63.Porter AE, Best SM, Bonfield W. Ultrastructural characterisation of
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71.Burr DB, Martin RB. Calculating the probability that microcracks initiate
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73.Stepnik MW, Radtke CL, Scollay MC, Oshel PE, Albrecht RM, Santschi
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74.Bronckers ALJJ, Goei W, Luo G, Karsenty G, D'Souza RN, Lyaruu, DM,
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75.Blickenstaff LD, Morris JM. Fatigue fracture of the femoral neck. J Bone
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76.Lee TC, O'Brien FJ, Taylor D. The nature of fatigue damage in bone. Int J
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77.Da Costa Gomez TM, Radtke CL, Kalscheur VL, Swain CA, Scollay MC,
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79. McMullen, R. An in vitro investigation of MC3T3-E1 osteoblast proliferation
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80.Flautre B, Descamps M, Delecourt C, Blary MC. Porous HA ceramic for
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81.Li SH, De Groot K, Layrolle P. Bioceramic scaffold with controlled porous

structure for bone tissue engineering. Key Engin Mater 2002;218-220z25-
30

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82. Wilson CE, de Bruijn JD, van Blitterswijk CA, Verbout AJ, Dhert WJA.
Design and fabrication of standardized hydroxyapatite scaffolds with a
defined macro-architecture by rapid prototyping for bone-tissue-
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83.Bucholz RW. Properties of osteoconductive biomaterials: calcium
phosphates. Clin Orthop Relat Res 2002;395:81-98.

84.deGroot K. Effect of porosity and physicochemical properties on the
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Acad Sci 1988;523:227-233.

21

CHAPTER 2

EXPERIMENTAL PROTOCOLS

I. Overview

This chapter is a detailed outline of all experimental procedures used in
the studies contained in this dissertation, including that work which has been
published previously."3 It is divided into sections which each describe a general
area of the work: cell culture and care, specimen fabrication and cell culture and
finally quantitation. Within each section, each individual protocol is described. I
have included an example of a statistical quantitation using actual data in

Appendix C.

II. MC3T3-E1 osteoblast culture"3

This section describes the basic cell culture protocol for the upkeep of
MC3T3-E1 murine (mouse) osteoblasts (088). These protocols are vital prior to
a specific trial, when they are used to expand the OB population to a number
suitable for that study.

The stock media used in MC3T3—E1 OB culture is referred to as

“complete” and consists of 90% alpha minimum essential media (Ct-MEM,

Invitrogen, Carlsbad, CA) and 10% fetal bovine serum (FBS, Atlanta Biologicals,

Lawrenceville, GA).

22

 

035 were removed from the tissue culture substrates enzymatically, using
10% Trypsin-ethylenediaminetetraacetic acid (Trypsin-EDTA, Invitrogen,

Carlsbad, CA).

A. Freezing cell stock

Long-terrn storage of 083 requires storage under extremely cold
conditions. Prior to storage at -180°C (93 K), 085 must be quickly frozen. A
suspension of 1,000,000 OBs in complete media is placed into a 50 ml centrifuge
tube and spun in a centrifuge (3000 rpm, 10 minutes) to separate the cells from
the media. The large number of cells included in each freezing batch takes into
account an expected post-thaw survival rate of < 50%. This population results in
a workable thawed batch cell population. The volume of suspension used here is
irrelevant and varies depending on the cell population density. The complete
media is removed by slowly pouring it into a waste container and 1.5 ml of
freezing media is added. Freezing media contains 90% complete media and 10%
dimethyl sulfoxane (DMSO). The 08s are then resuspended by pipetting up and
down 10 times and transferred to a 2 ml cryovial (Corning, Wilkes-Barre, PA).
Cryovials are placed into a -80°C (193 K) freezer for a period from a minimum of
12 hours up to 2 weeks. Twelve hours allows sufficient time for the cells to
freeze, while the maximum time of 2 weeks prevents unnecessary loss of cell
viability due to storage at a temperature higher than -180°C (93 K). Vials are then
transferred to a liquid nitrogen (LN2) storage tank (-180°C/ 93 K) for long-terrn

storage.

23

 

B. Thawing from stock

Cell stock is kept at -180°C (93 K) for long-term storage (longer than 2
weeks). Cells are divided into groups of 1-million, each in a 1.5 ml polypropylene
cryovial. Henceforth, each group of cells will be referred to as a vial.

Prior to each study, MC3T3-E1 murine osteoblasts were thawed from
frozen stock by placing each vial between the hands and rolling it back and forth
until thawed (approximately 2 minutes). Each vial was then submerged in 70%
EtOH (prepared by combining 90%. EtOH [Sigma-Aldrich, St. Louis, MO] with
enough RO water to yield a final concentration of 70%) for one minute, removed
with long handled stainless steel forceps and placed onto a piece of dry cotton
gauze to dry. Using a piece of gauze in each hand, each vial was then opened
and poured into the center of a 100 mm polystyrene tissue culture plate (Petri
dish).

Prior to the thawing procedure described above, complete media was
warmed by placing the 500 ml stock bottle in a water bath at 37°C (310 K, normal
body temperature) for approximately 15 minutes. Ten ml of complete media was
then added to each culture dish containing the thawed contents of a cell vial.

After 24 hours, the complete media was removed by aspiration, non-
adherent cells rinsed away by twice pipetting onto and removing 1X phosphate-
buffered saline (PBS) from the culture plate surface, in order to eliminate'dead

cells from the counts, and the thawed 085 were removed from their plates,

I‘

. 100% EtOH (200 proof) is available but is not used in this case. This is due to the presence of
benzyne in 200 proof EtOH, which is used as a stabilizing agent and is toxic to living cells.

24

" '“wfi-IJ. It“ >0“ ‘1' A, L . ,m “In.
s......-..1m‘ .é..rmmamx.m it.“ 2113:: : .~ .

I:

.
3:.
;1
.51..
'319'
‘:
;a
‘1
, .:.

'I'J'é.” nun-5“
: «new»... ‘1 I'v‘bq-O 'v.‘
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_ . ,

 

 

p

I o
n
V

1.4., . x
"v o N"
in

1; .. - 11' :-“.'..'::;.;;.:.
"1|!!1N0- Noun-m: 4 ,4

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counted and replated at 60,000 — 75,000 cells/culture dish. This density is

sufficient for the plated cells to divide without reaching confluency.

C. Splitting/Feeding

The term “splitting” is used when an OB culture is removed from its
substrate, divided into smaller groups and plated at a lower cell density. Healthy
08s are normally split every 2 — 3 days when using 100 mm cell culture plates.
Experience shows that this time interval leads to a population suitable for
splitting. It should be noted that more active cells (lower passage number) may
split sooner than this. Cells are ready to be split when ~25 cells are visible in
each field of view under the microscope (10x objective). This population is large
enough to be nearing, but not at, cell confluency. Confluency is a state where
adjacent cells are touching each other, and may lead to subcloning of the cells.

Plates selected for splitting are moved to the biosafety cabinet (positive
pressure, HEPA filtered, officially certified) and their media removed by
aspiration. Aspiration uses a narrow glass tube (9 inch Pasteur pipette)
connected to the central vacuum system draining into a 4 liter waste flask. Once
the media is removed, the plate is rinsed twice with sterile 1X PBS and 2 ml of a
1:1 mixture of Trypsin-EDTA and complete media per plate in order to free the
CBS from the culture surface. This process takes less than 2 minutes and is
assisted by rocking the plate back and forth and side to side (90° apart) several
times. Finally, the newly liberated cells are pipetted in suspension, up and down

five times, washing the entire surface of the plate. OBs from each plate are

25

combined in a 50 ml centrifuge tube prior to counting. Typically, a batch from 10
— 20 plates will yield 1.5 — 2 million cells. This population is then partially plated
(~20 plates) and partially frozen (1 — 2 cryovials). For lower passage cells, the
amount that is frozen is ~ 90%, in order to preserve cell stock.

Cells are counted with a hemacytometer and replated at 60,000 — 75,000
per 100 mm plate. Ten ml of complete media is added to each plate and they are

placed into the incubator at 37°C (310 K) and 5% 002 and the process repeated.

 

 

 

 

 

 

:" T 2' 4. .r""”j-I. I! I: I .
_. 1- TrypsinfEDTA-_ _--_ 1713' 3_
a... “1 e 3. ‘1 , a it

L' 2 l ix”; i—‘tI, ."
' ‘“ , I —* f»~-———-~——~-—~ ———--~— ;_'-.-: 5
/ Cel| /\'4 '1' .
Cells . 1‘ ;_; Suspension ,3
Complete Media l c 'f
l/ .. xx. 6. ____p Replated Cells
eIIs
Hemacytometer

Figure 1: Schematic detailing the cell splitting process. Media is removed by aspiration (1) and
the cells are rinsed twice with 1x PBS. Trypsin/EDTA and complete media are added in a 1:1
ratio (2) and each plate is rocked back and forth (3) to aid in cell detachment. Detached cells are
pipetted away and placed into a 50 ml centrifuge tube, along with the cells from other plates (4).
10 pl of this suspension is placed into each of the two sides of a hemacytometer (5). Cells are
then viewed under the Optical microscope (6) and the cells counted (7). Using this cell population
data, the remaining cell suspension is divided up between fresh plates (60,000 - 75,000 per
plate) and 10 ml fresh complete media added to each (8).

III. Specimen fabrication:
This section will be divided into two subsections, one for porous and
another for dense hydroxyapatite (HA). Further divisions will detail the various

steps taken to create specific sample types.

A. Porous hydroxyapatite scaffolds1

26

Two groups of porous scaffolds were fabricated by foaming using
hydroxyapatite (HA) with two different average particle sizes (diameter).

Micron scale HA powder was obtained from Hitemco Medical (Old
Bethpage, NY) and had a manufacturer-reported particle size of 10 pm. This
powder was used in its as-received form. As received powder arrived sintered.

Porous specimens were formed by first suspending HA powder in a 10'3 M
aqueous solution of KNO3. Powder and electrolyte were combined in a beaker
and mixed with a stainless steel spatula. A typical batch of suspension would
consist of 33 g of HA and 17 ml of KN03 solution. This ratio yielded a suspension
of suitable consistency (similar to pancake batter) for eventual casting of
cylinders. To this suspension, 1.8 ml of 30% H202 was added'as the foaming
agent and the suspension stirred once again. This electrolyte was used because

it aids in the dispersion of HA powder particles by creating a large double layer
thickness (1/K), resulting in uniform porosity across the cylindrical specimens.4

The HA + H202 slurry was scooped into cylindrical silicone molds, each
with a diameter of 15 mm. These molds were fabricated out of a two-part silicone
kit (SI-25 SlL—MOLD, Hyperlast North America, Chattanooga, TN) and consisted
of two pieces. The main piece contained the cylindrical mold chambers and the
second piece was placed underneath, to prevent the slurry from dripping out.
Each mold was covered with aluminum foil and placed into a gravity convection
oven (Lindberg/Blue) at 125°C (398 K) for one hour. This time and temperature

combination effectively dried the specimens sufficiently prior to sintering. Each

27

 

batch of suspension yielded approximately 6 green (unfired) specimens with a

diameter of 1.5 cm and a length between 2 and 4 cm.

 

Figure 2: Silicone mold used for the foaming and casting of cylindrical porous HA scaffolds. Both
pieces are shown, including the main piece with its cylindrical mold chambers and the solid base
piece.

After one hour, these green HA specimens were gently ejected from their
mold. In most cases, this was accomplished by carefully lifting the main piece of
each mold away from the base, allowing the cylinders to remain behind. In some
cases though, the HA cylinders would remain in their mold. To eject these HA
specimens, the flat end of a syringe plunger was used and gentle pressure
applied to eject each specimen out. The green HA specimens were then placed
into a high-purity Al203 sintering dish, with Al203 cover, and immediately sintered
in air at 1360°C (1633 K) for four hours, with a heating/cooling rate of

10°C/minute. This sintering time and temperature produces a specimen that is

chemically HA and the prescribed heating/cooling rate prevents thermal shock,

28

both in the specimens and in the furnace heating elements. Sintering was
performed immediately in order to maintain residual moisture in the green
specimens. This moisture allows the particles to remain in their post-foamed
positions and leads to a final product with structural integrity and consistent pore
structure. Sintering was completed using a high temperature CM Rapid Temp
sintering furnace (CM, Bloomsfield, NJ) equipped with molydisilicide heating
elements.

Sintered specimens were sectioned into 1 mm slices using a Buehler
Isomet 1000 diamond abrasive wafen’ng saw (Buehler, Lake Bluff, IL, USA)
operated at 200 rpm and reverse osmosis (RO) water as a lubricant. Each
cylinder was gripped using a custom fabricated padded aluminum specimen

holder and slice thickness was determined using the built-in controls of the saw.

 

Figure 3: Figure 3A is a view of the Buehler Isomet 1000 wafering saw used to section the porous
cylinders. Figure 3B is a close-up of the custom-built chuck used to hold each cylinder during
cutting.

Nano-scale HA powder with a manufacturer reported average particle size

of 20 nm was purchased from Berkeley Advanced Biomaterials (San Leandro,

29

CA, USA). The powder arrived in an ammonium hydroxide suspension, which
was subsequently centrifuged (3000 rpm, 10 minutes) to separate the powder.
The nano HA was then resuspended in 90% stock EtOH, centrifuged at 3000
rpm for 10 minutes and vacuum dried for 12 hours at room temperature in a
vacuum furnace (lsoTemp 280A, Thermo Fisher, Watham, MA). This method of
separating the powder was prescribed by the manufacturer.

Foaming and sintering were accomplished in the same manner described
above for micron HA powder, with some specific exceptions. First, each batch of
nano HA slurry contained 22 g of HA powder and 18 ml of 10'3 M KNO3. This
ratio yielded the desired consistency for foaming and casting of cylindrical
scaffolds. Second, green specimens were sintered at 1100°C (1373 K) for one
hour with a 10°C/min heating/cooling rate, using the same CM high temperature
furnace. This sintering time and temperature has been used in the past to
successfully yield nano-grained HA.5'10

The average grain size in porous HA scaffolds was determined by
scanning electron microscopy (SEM). Five representative images were taken of
both micron and nano-scale HA scaffolds and the line-intercept method (n = 50)
used to determine average grain size and standard deviation. Pore size was
determined using the same method. Specimen porosity was determined using
Archimedes method with n = 5 for both micron and nano-scale HA scaffolds.
Sample size is chosen to aid in observing any variances in porosity between

specimens.

30

 

B. Dense hydroxyapatitez"3

Dense discs were fabricated using medical-grade HA powder (Taihei
Chemical, Osaka, JP) with vendor-reported mean particle diameter of 1 — 3 pm.
HA powder was uniaxially pressed into discs at 35 MPa for 1 minute using a
hydraulic press (Carver Inc, Wabash, IN) and KBr pellet die (International Crystal
Laboratories, Garfield, NJ). This force/time combination yielded green specimens
with uniform density and minimal incidence of breakage. The resulting discs had
a diameter of 32 mm and were 2 — 3 mm thick. These discs were placed in a
single layer inside of Al203 sintering boats and sintered in air at 1360°C (1633 K)
for 4 hours with a heating/cooling rate of 10°C/min (CM furnace) and the resulting
HA discs were typically 23 mm in diameter and 1 — 2 mm thick, with a >99%
theoretical density. These sintering specifications yield specimens determined to
be HA and this thickness led to minimal warping.

Average grain size for dense HA specimens were determined using SEM
and the same line intercept technique described in the previous section.
Specimen density was calculated by comparing actual specimen mass for their
measured volume (n = 5) to theoretical values expected for that volume.

Any macroscopic imperfections on the HA discs were ground flat using 30

pm (600 grit) SiC wet abrasive paper (Carbimet 8” abrasive discs, Buehler, Lake

Bluff, IL) using a sequence including 5 minutes each at 30 pm, 12 pm, 9 pm, 6

pm and 1 pm diamond abrasive suspension or paste. Grinding and polishing

were performed with a Buehler EcoMet 3 variable speed grinder/polisher

equipped with an AutoMet 2 power head. After polishing, specimens were

31

ultrasonically cleaned by submersion in reverse osmosis (RO) water and pulsing
with an ultrasonic probe at 20 kHz and 250 watts for 5 minutes.

In both microcracking and grain boundary grooving studies, control
specimens were dense HA discs with average grain sizes equal to those of the
non-control specimens. These HA discs were also polished, sterilized and
seeded in identical fashion to the microcracked and grain boundary grooved
specimens.

Microcracks were induced by systematically indenting each specimen 196
times (14 x 14 grid with 1 mm spacing) using a Buehler Micromet ll microindentor

(Buehler, lake Bluff, IL), with a load of 4.9 N, a load time of 5 seconds and a

loading rate of 70 pm/sec.2'3 We found these indentation parameters to induce
microcracks with minimal spelling and chipping of the HA specimens.
Microcracks were formed on the HA surface by the radial cracks originating from

the comers of each indentation impression. lndent spacing was assured by

turning the stage micrometer two full turns, equaling 1 mm, between each indent.

32

Vickers V- lndentor
“a“ Indentation

’,-

Sinteredi’Polished
HA Disc

  

,1" Measurement
1

1' /' Pattern
+

 

Indents

Figure 4: Schematic of indentation/induced microcracking and viewing/measuring of microcracks
using optical microscopy. Indents selected for measurement ("X" pattern) are indicated by circles.

21:

 

Figure 5: The image at left is a schematic of a typical Vickers indentation. The length “2a" is the
width of the indentation pyramid and “20" is the length of a microcrack. The image at right is an
SEM micrograph of a single indentation on the surface of a dense sintered HA specimen. The
black arrows indicate microcracks. Figure taken from Smith et al.3

In order to accelerate crack growth saturation, each HA disc was
submerged in R0 water for 8 hours after indentation.3 The average indentation

size and crack size were then determined by optical microscopy by measuring a

33

 

representative selection of indentations (n = 20) for each HA specimen. These
indents were taken in a pattern resembling a large “X" across the entire
specimen surface (See Figure 4).

Thermal etching was used to induce grain boundary grooving in HA in
order to determine its effect on OB attachment. Polished HA discs were heated
to 1100°C (1373 K) for 45 minutes, with a heating/cooling rate of 10°C/minute.
SEM was used to determine whether etching was successful and to estimate the
width of the grain boundary grooves, using AnalySlS software (Olympus Soft
Imaging Solutions Corp Lakewood, CO). Due to its magnitude, the depth of these

grooves could not be estimated using SEM.

 

Figure 6: SEM micrograph taken at the surface of a sintered dense HA specimen that has been
thermally etched at 1100°C (1373 K) for 45 minutes to induce grain boundary grooving.

34

The average surface roughness (Ra) was measured using a Veeco
DEKTAK 6M profilometer, with a stylus radius of 2.5 pm and a measurement
length of 3000 pm. Three measurements were taken on the specimen surface, to

determine average Ra and standard deviation.

    

—>

Direction of Profilometry Scan

 

Figure 7: Schematic of a profilometry scan across the specimen surface.

IV. Cell culture/counting:

This section is divided into subsections, for porous and dense
hydroxyapatite. Since cell culture, imaging and counting protocols used for
microcracked and thermally etched (grain boundary grooved) hydroxyapatite
were the same, they have both been included under “dense hydroxyapatite”.

Prior to cell culture, all HA specimens were sterilized using the same
protocol. Specimens were sterilized by steam autoclave (Tuttnauer, Hauppauge,

NY) at 125°C (398 K) for 45 minutes with 30 minutes drying time. Specimens

35

were placed into beakers, which were covered with aluminum foil and were

autoclaved in the absence of any other materials.

A. Porous hydroxyapatite1

MC3T3-E1 subclone 7 08s of passages 22-23, cultured in complete
media were seeded across the surface of both micron and nano-scale porous HA
scaffolds inside of AI collars at 11,320 cells/cm2 and incubated in air at 37°C (310
K), 5% CO2 and high humidity. Subclone 7 cells were used because they are
known to differentiate into 085 (Reported by supplier). Complete media was
replenished every two days, following the first 24 hours. Three specimens were
seeded at 0.5, 1, 2 and 4 hours, 1, 3 and 5 days and three sets of cultures were
examined, for a total of 63 samples.

OB attachment was assessed at intervals of 0.5, 1, 2 and 4 hours and OB
proliferation at 1, 3 and 5 days. Prior to cell counting, scaffolds were removed,
placed into separate wells, rinsed twice with 1x phosphate-buffered saline (PBS),
trypsin/EDTA: complete media (1 :1, a ratio which allow detachment at a sufficient

rate) added and the scaffolds morselized to facilitate OB detachment, using a
mortar-and-pestle motion with the handle end of a cell scraper. Ten pl of cell

suspension was pipetted from the center of each well and placed into each of the

two sides of a hemacytometer for counting.

36

 

Figure 8: The cell scraper used to morselize porous HA scaffolds prior to counting attached 08s.
The handle end, located at the left end of the photo, was placed into each well and the specimen
morselized.

Statistical differences in OB counts on micron and nano-scale HA
scaffolds were determined by using the Student’s t-test. This test is useful in
determining significance in the difference between means of two groups. A p
value of 0.05 or lower was considered to be significant. This level of p value is

commonly associated with significance.

37

Collar

 

 

 

 

 

 

 

 

Figure 9: Schematic showing the process of counting cells on porous HA scaffolds. Cells are
removed from their tissue culture plates (1) and seeded onto sterile porous HA scaffolds, placed
inside of Al foil collars (2). Each collared scaffold is placed into a 6-well polystyrene tissue culture
plate and into the incubator for the prescribed time. After incubation, each specimen is removed
from its collar and placed into a fresh 6-well plate, rinsed twice with 1x PBS, exposed to
Trypsin/EDTA and the scaffold morselized with the handle end of a cell scraper (3). 10 pl of the
resulting suspension is placed into each of the two sides of the hemacytometer (4) and the cells
(with HA particles) are viewed under the optical microscope (6) and counted (7).

Issues which arose while using the hemacytometer for cell counting are
addressed in detail in Chapter 4. In short, we attempted to correct the error
observed in our cell counts by subtracting away background HA particles. This
involved repeating the above procedure without the cells present. Morselized
scaffolds were placed into the hemacytometer and false-positive “cells” were
counted. If a particle resembled a cell, as it appeared during a normal cell count
with scaffold particles present, it was tallied.

In future studies, cells should be counted by first staining to differentiate
them from the surrounding scaffold particles and then by counting, either directly
by optical microscopy or confocal laser scanning microscopy (CLSM), or

indirectly by spectrophotometry.

38

B. Dense hydroxyapatite3

MC3T3-E1 subclone 14 0B5 were obtained from ATCC (Manassas, VA).
Sterile microcracked HA and control HA specimens were placed into Al collars
inside 6-well polystyrene cell culture plates and OBs were seeded onto the HA
discs at a density of 11,320 cells/cmz. Two separate trials were conducted for
incubation times of 1 and 4 hours each. OB seeding was done on specimens in
triplicate, with n = 3 for a total of nine microcracked HA and nine control HA

specimens for each trial. The triplicate model is commonly used to improve

 

statistical significance for in vitro cell culture studies. The same protocol was
used for seeding OBs onto grain boundary grooved (and corresponding control)
HA specimens. The only difference is that MC3T3-E1 subclone 7 cells were used
instead.

In both microcracking and grain boundary grooving studies, OBS were
fixed using 3.7% formaldehyde (H2CO) (JT Baker, Phillipsburg, NJ) in order to
terminate the life cycle and preserve the physical state of the cells prior to their
death, as well as perrneabilized with a solution of 0.1% Triton X-100 (Sigma-
Aldrich, St. Louis, M0) to open the cell membranes. Prior to staining, OB—seeded
HA specimens were pre-incubated in 1% bovine serum albumin (BSA,
Invitrogen). Pre-incubation improves the uptake of Rhodamine-Phalloidin into the
cells.

The actin cytoskeletin was stained using Rhodamin-Phalloidin while the
nucleus was stained using Hoescht 3342 (Invitrogen). These stains penetrate the

cell membrane and selectively stain particular organelle (in this case the

39

_—

 

 

cytoskeleton and nucleus, respectively). Stained OB-seeded HA specimens were
mounted in glycerol solution (JT Baker) to seal the specimen surfaces from the
air, and viewed using a Leica DM IL fluorescence microscope (Leica
Microsystems, Wetzlar, GER) and optical micrographs recorded using the Spot
RT camera system (Diagnostic Instruments, Inc. Sterling Heights, MI).
Rhodamine-Phalloidin stain fluoresces at a wavelength of 565 nm when exposed
to light of a wavelength of 542 nm while the Hoescht stain fluoresces at a
wavelength of 465 nm when exposed to light with wavelength equal to 355 nm.

For the microcracking study, OB attachment was quantified by counting
OB nuclei over the microscope’s field of view (h = 1000 pm, w = 1500 pm). Ten

fields of view were selected and counted for each specimen and the Student’s t-
test was used to determine statistical significance, with p < 0.05 taken as
significant. A similar protocol was used for the grain boundary grooving study,
with the only difference being that 20 fields of view were selected per specimen.
The effect of induced microcracking in HA specimens on OB morphology
was assessed by staining the actin fibers and observing the OB shape using
optical fluorescence microscopy. Actin fibers are a component of the

cytoskeleton and spread throughout the cell, following its shape.

40

 

Induced Microcracking

L.
_-r- - .
(- m ““3 M i \_
‘-___ 'L ‘4" ' _" '\~_

1 .____ L..-» 2. ---.__. ._._---

 

 

 

 

"A Powder Pressed, Sintered and
Polished HA Disc
3.
( al.-“7:27;. 4. {xi-2:312: 7:"
' ' -—;:.i-——— ._ -'3--"—-;i -—
Stained 085 on Disc 8 Seeded Disc
5. ._
l, .............. .

 

6. Field ofView

Figure 10: Schematic showing the process of fabricating, microcracking, cell seeding, staining
and counting on dense HA. Powder is uniaxially pressed, sintered and polished (1) prior to
induced microcracking (2). Microcrack size is determined, the specimens sterilized and CBS are
seeded onto their surfaces (3). Following culture times of one and four hours, the OBs are stained
with Rhodamine/Phalloidin and Hoescht (4) and viewed under an optical fluorescence
microscope (5). Fields of view are chosen as shown (grey boxes each indicate one field of view)

41

REFERENCES

1. Smith IO, McCabe LR, Baumann MJ. MC3T3-E1 Osteoblast attachment
and proliferation on porous hydroxyapatite scaffolds with nanophase
powder. Int J Nanomed 2006;1:189-194.

2. Case ED, Smith IO, Baumann MJ. Microcracking in hydroxyapatite and
the implications for bone tissue engineering. Mater Sci Engin A
2005;390:246-254.

3. Smith IO, Baumann MJ, Case ED. Induced microcracking affects
osteoblast attachment on hydroxyapatite scaffolds for tissue engineering.
J Am Biochem Biotechnol, 2006;2z105-110.

4. Hiemenz PC, Rajagopalan R. Princliples of colloid and surface chemistry.
New York: Marcel Dekker; 1997. 650p.

5. Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R. Nanocrystalline
hydroxyapatite enhances osteoblast function. In: Proceedings of the First
Joint BMES/EMBS Conference, Atlanta, GA; 1999. p744.

6. Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R. Enhanced
functions of osteoblasts on nanophase ceramics. Biomaterials
2000;21:1803-1810.

7. Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R. Specific
proteins mediate enhanced osteoblast adhesion on nanophase ceramics.
J Biomed Mater Res 2000;51:475-483.

8. Webster TJ, Siegel RW, Bizios R. Design and evaluation of nanophase
alumina for orthopaedic/dental applications. Nanostruct Mater
1999;12:983-986.

9. Webster, TJ, Siegel, RW and Bizios, R. Osteoblast adhesion on
nanophase ceramics. Biomaterials 1999;20:1221-1227.

10.Webster TJ, Siegel RW, Bizios R. Nanoceramic surface roughness

enhances osteoblast and osteoclast functions for improved
orthopaedic/dental implant efficiency. Scripta Mater 2001;44:1639-1642.

42

CHAPTER 3
MC3T3-E1 OSTEOBLAST ATTACHMENT AND PROLIFERATION ON
POROUS HYDROXYAPATITE SCAFFOLDS FABRICATED WITH
NANOPHASE POWDER
IAN 0. SMITH, LAURA R. MCCABE2 AND MELISSA J. BAUMANN1
1Department of Chemical Engineering and Materials Science, Michigan State
University, East Lansing, MI, USA
2Department of Physiology, Michigan State University, East Lansing, MI, USA

PUBLISHED IN: Int J Nanomed 2006;1:189-194.

ABSTRACT
Porous bone tissue engineering scaffolds were fabricated using both nano

hydroxyapatite (nano HA) powder (20 nm average particle size) and micro HA

powder (10 pm average particle size), resulting in sintered scaffolds of 59 vol%

porosity and 8.6 i 1.9 pm average grain size and 72 vol% porosity and 588 :l: 55

nm average grain size, respectively. Scanning electron microscopy (SEM) was
used to measure both the grain size and pore size. MC3T3-E1 osteoblast (OB)
attachment and proliferation on both nano HA and micro HA porous scaffolds
was quantified. As expected, OB cell number was greater on nano HA scaffolds
compared to similarly processed micro HA scaffolds 5 days after seeding, while

OB attachment did not appear greater on the nano HA scaffolds (p < 0.05).

INTRODUCTION
Bone is a composite material having a calcium phosphate (CaP) mineral
component and a collagen-based organic matrix. CaP is a bioceramic resembling

hydroxyapatite (HA, Ca1o(PO4)6(OH)2), where the individual crystals have a plate-

43

like morphology, 15 - 200 nm in length and 10 — 80 nm in width with a thickness
between 2 and 7 nm. (Fratzl et al, 2004) These apatite crystals are imbedded
within the collagen fiber component of the organic matrix, which forms lamellar
sheets comprising the concentric rings of osteons in cortical bone (Young and
Heath, 2000) and the more woven structure of trabeculae in cancellous bone.
(Hassenkam et al, 2004) The collagen matrix is composed primarily of Type I
collagen fibers of approximately 100 nm in diameter. (Fratzl et al, 2004)

These nano-components play an active role in bone development and
growth. Collagen fibers act as the scaffold on which newly formed nanocrystals
of bone apatite form and grow, leading eventually to the formation of woven
bone. This primitive early phase of bone is eventually replaced by the lamellae of
adult bone. (Young and Heath, 2000) Additionally, collagen fibers have been
shown to incorporate into porous HA tissue engineering scaffolds where the
interfacial zone between bone and scaffold has been found to be a nano HA-
reinforced collagen matrix. (Chen et al, 2004)

One area of current research in biomimetic nano-scale ceramic bone
substitute materials is therefore focused on the design of nanoscale
ceramic/polymer composites to mimic some of the morphological features of
natural bone. These systems can be grouped into three categories: nano-
CaP/collagen (Du et al, 1998; Du et al, 1999; Sakane et al, 2004), nano-
CaP/non-collagen polymer (Wei et al, 2003; Wei et al, 2004a; Wei et al, 2004b)
and a combination of the two (Cui et al, 2004; Liao et al, 2004). Although there

are several types of CaP bioceramics, HA is the most commonly used, having a

44

chemistry that closely matches normal bone apatite. (Hench, 1998; Jarcho, 1981;
Li et al, 2004)

Nanocrystalline HA has been shown to be biocompatible with a minimal
inflammatory response (Silva et al, 2004), (Huang et al, 2004), although there is
a potential for cytotoxicity. (Huang et al, 2004) While some engineered
nanoCaP/polymer composites such as HA/collagen polylactic acid (PLA)
scaffolds from Liao et al reported improvements in compressive strength using
nano HA, they did not replicate the hierarchical structure of bone which may be
crucial in encouraging bone ingrowth into a tissue engineered scaffold. (Du et al,
1998; Liao et al, 2004) Also, Yoneda et al. confirm that the degradation of PLA
adversely affects cell behavior specifically, that increased degradation leads to a
decrease in cell spreading, adhesion and cell survival. (Yoneda et al, 2004)
Therefore, current research focuses on developing a bioceramic with a
morphological structure similar to normal bone to promote bone regeneration in
vitro and in vivo. (Longsworth and Eppell, 2002)

Webster et al compared changes in osteoblast (OB) behavior on nano-
grained versus micro-grained dense bioceramics, including AI203, TiO2 and HA,
demonstrating increases in OB attachment (Webster et al, 19993; Webster et al,
1999b; Webster et al, 1999c) and proliferation (Webster et al, 1999a; Webster et
al, 2000a), alkaline phosphatase (AP) activity (Webster et al, 1999a; Webster et
al, 2000a) and calcium deposition (Webster et al, 19993; Webster et al, 2000a),

as well as decreased cell motility (Webster et al, 1999a; Webster et al, 2000a).

45

Osteoblast attachment on a bioinert ceramic like AI203 can be attributed to
the presence of physical surface irregularities that lead to increased
morphological fixation. (Hench, 1998) Such fixation in turn plays an important
role in the use of TiO2 as a coating for orthopedic implants, improving the
integrity of the implant/bone interface. (Areva et al, 2004) Fixation also plays a
role in OB attachment to HA, but increased osteoblast adhesion may also occur
because of improved protein interactions that have been linked to decreases in
grain size.25

Enhanced OB behavior for HA having a finer grain size is further
illustrated in a study by Xie et al, which recently reported that MC3T3-E1 083
seeded onto dense HA discs showed differences in gene expression related to

grain boundary surface area. (Xie et al, 2004) Gene expression of OBs on HA
discs fabricated with an average particle size of 35 pm showed a decrease in
cellular metabolism and an increase in differentiation while HA discs processed
from smaller particle sizes (~11 um) showed gene expression profiles that

signified increased cell growth. Each of the previous studies was conducted
using dense ceramic scaffolds, designed specifically for cell culture purposes,
while the current study assesses whether porous bone replacement scaffolds
fabricated with nano HA powder rather than micro HA improves OB attachment

and proliferation.

46

MATERIALS AND METHODS
HA powder, of 10 pm average particle size, was purchased from Hitemco

Medical (Old Bethpage, NY, USA) and used in its as-received form. Nano-scale
HA powder having an average particle size of 20 nm was purchased from
Berkeley Advanced Biomaterials (San Leandro, CA, USA). The nano HA powder
arrived in an ammonium hydroxide suspension which was subsequently
centrifuged, resuspended in ethanol, centrifuged again to separate out the nano
HA powder and then vacuum dried for 12 hours before use (per manufacturer
instructions).

Methods used in this study can be divided into four parts: scaffold
fabrication, imaging of scaffold surfaces, osteoblast cell culture and assessment
of OB attachment and proliferation.

Porous cylindrical samples of each HA powder particle size were foamed
by suspending the HA powders in KN03 electrolyte to which H202 had been
added, pouring the resulting foamed suspension into cylindrical molds and drying
at 125°C. The dry, green HA samples were then sintered at 1100°C for 1 hour
(nano HA) following the procedure established by Webster et al. (Webster et al,
1999a; Webster et al, 2000a; Webster et al, 2000b) or 1360°C for 4 hours (micro
HA) using a protocol (McMullen, 2004) developed by McMullen. Both the micro
and nano HA sintered scaffolds were next sectioned using a Buehler lsoMet
1000 diamond abrasive wafering saw (Buehler, Lake Bluff, IL, USA) into discs of

approximately 1 mm thickness and autoclaved.

47

 

 

Nano HA and micro HA scaffold surfaces were observed using scanning
electron microscopy (SEM). Samples were coated with a thin layer of Au,
approximately 21 nm thick, using an Emscope SC500 sputter coater (Emscope
Laboratories Ltd, Ashford, UK), mounted onto Al stubs and imaged using a JEOL
6400 SEM (JEOL Corporation, Tokyo, Japan) operated at an accelerating
voltage of 5 keV. SEM photomicrographs were collected and exported using the
AnalySlS software package.

MC3T3-E1 OBs of passage 22-23, cultured in alpha Minimum Essential

Medium (Oi-MEM) supplemented with 10% Fetal Bovine Serum (FBS), were

trypsinized, resuspended in a-MEM and seeded evenly across the surface of
porous HA scaffolds at a density of 11,320 cells/cm2 and incubated in air at 37°C

(310 K), 5% CO2 and high humidity. (Shu et al, 2003) 2 ml of Ot-MEM was added

per specimen and was replenished every two days. Three specimens were
statically cultured per time interval and two sets of cultures were examined,
resulting in 6 samples per time interval. MC3T3-E1 OBs are a widely used 0B
cell line (Cerroni et al, 2002; lyer et al, 2004; Shu et al, 2003; Werner et al, 2002;
Xie et al, 2004) that have been shown to behave in a similar manner to primary
OBs. (Griffin et al, 2005)

QB attachment was assessed at intervals of 0.5, 1, 2 and 4 hours.

Scaffolds were removed, placed into separate wells, rinsed with phosphate-
buffered saline (PBS), trypsin/ot-MEM (50/50) added and the scaffolds

morselized to facilitate osteoblast detachment. Media was drawn from each well

48

and the 083 counted using a hemacytometer. OB cell counts were also

conducted at 1, 3 and 5 days following attachment to assess OB proliferation.
Statistical differences between attachment and differentiation on micro HA

and nano HA scaffolds were calculated using the Student ttest. A p value of 0.05

or lower was considered to be significant.

RESULTS
Porous HA scaffolds were successfully fabricated by foaming and

sintering both micro HA and nano HA powder as described in the previous

section. SEM showed that the micro HA scaffolds have a grain size of 8.6 pm i

1.9 pm which is approximately 15 times greater than the nano HA grain size of

588 nm 1 55 nm, (representative micrographs given in Figure 1). The average
grain size was determined using a line-intercept method (n = 50) based on ASTM
standard E 112.33 The volume porosity of micro HA and nano HA scaffolds using
Archimedes method were found to be 59 vol% and 72 vol %, respectively.27 Both
micro and nano HA scaffold surfaces consisted of grains which are the sintered

particles interspersed with frequent gaps along with larger

macropores/micropores of 390 pm i 210 um/ 38 pm i 22 pm for micro HA and
480 pm .1: 275 pm/ 30 pm i 21 pm for nano HA. These size ranges fall within

standard macropore (200-400 pm) and micropore (<60 pm) ranges reported by

deGroot et al. (deGroot et al, 1988) A representative SEM photomicrograph

49

 

 

 

showing the nature and distribution of the porosity within the pore wall of a nano

HA scaffold is given in Figure 2.

 

Figure 1: Comparison of SEM micrographs of the surface of scaffolds fabricated using micro HA
(left) and nano HA (right).

50

 

Figure 2: SEM micrograph showing the surface of porous scaffold fabricated using nano HA,
including macropores (large arrows noting approximate boundary of a macropore) and
micropores (small arrows emphasizing the <60 pm interconnecting pores).

Contrary to published data by Webster et al using dense HA scaffolds,
these results show that initial 0B attachment does not increase as a function of
decreased grain size on porous HA scaffolds. (Webster et al, 1999a; Webster et
al, 2000a; Webster et al, 2000b; Webster et al, 2001) MC3T3-E1 OB attachment
on porous scaffolds is plotted as a function of time for 0.5, 1, 2 and 4 hours
(Figure 3). Based on hemacytometer analysis, OB attachment on the micro HA
scaffolds appears greater than on the nano HA scaffolds and differences were

statistically significant (with p < 0.05).

51

% Cells Plated

 

600

500 ~

400 1

300 4

200 1

100 ~

 

 

 

 

1.5

T

2

fl

2.5

Time (Hours)

4.5

Figure 3: MC3T3—E1 0B attachment on micro HA (dashed line) and nano HA (solid line) scaffolds
as a function of time, at intervals of 0.5, 1, 2 and 4 hours. Values are mean t SE, p < 0.05 at all
time intervals for comparison of nano HA to micro HA.

When OB cell number is examined, the mean values for OB proliferation

on nano HA scaffolds are higher when compared to the mean values on micro

HA scaffolds (Figure 4) at 1, 3 and 5 days culture time. Differences in

proliferation between micro HA and nano HA scaffolds at 1 and 3 days were not

found to be statistically significant, with p > 0.05 in both cases (p = 0.64 and p =

0.12, respectively). However, at 5 days, differences in osteoblast number were

found to be significant (p < 0.05).

52

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

250000

200000 i
2 U HA El nanoHA 'l'
3 150000 -
q.-
° I
h . .
E 100000 ‘ I
a .
z

50000 I I
0 , - ....... .
1 3 5

Time (days)

Figure 4: MC3T3-E1 Osteoblast proliferation on micro HA and nano HA, after periods of 1, 3 and
5 days. Values are mean i SE. p > 0.05 for comparison of nano HA to micro HA porous scaffold
for 1 and 3 days. p < 0.05 for 5 days (:t).

DISCUSSION

Porous bioceramic scaffolds offer advantages over dense substrates for
bone tissue engineering applications, because they present a more complex
interface for bone ingrowth that more closely mimics natural bone. Hench, 1998;
Hulbert et al, 1972; Hulbert et al, 1970) Additionally, porous CaP ceramics are
capable of promoting osteoprogenitor attachment and new bone formation.
(Ohgushi et al, 1990) Osteoprogenitor cells cultured on porous HA have been
shown to produce extracellular matrix in vivo, a precursor to bone formation.

(Ohgushi et al, 1990) Therefore, it is important to observe the effect of HA grain

53

size reduction on 0B behavior and to examine these differences for porous

scaffolds which are more suited for bone tissue engineering applications.

Foaming and sintering nano HA with a particle size of 20 nm yielded
porous scaffolds of submicron grain size (588 i 55 nm). While this grain size is
not truly nanoscale (< 100 nm), it is small enough to assess the effect of grain
size reduction (15X) on OB behavior in vitro. This reduction in grain size means
that there is an increased grain boundary surface area, which has been tied to
increased OB attachment (Perla et al, 2004) and enhanced gene expression (Xie
et al, 2004) indicative of increased cell growth and differentiation. Reduction in
grain size has also been linked to increased surface roughness, which in tum has
been shown to increase 0B adhesion, proliferation, AP activity and calcium
production. (Webster 2001) In future studies, atomic force microscopy (AFM) will
be used to measure the surface roughness of these scaffolds, to determine

whether these results are repeatable for porous HA.

Webster et al have repeatedly shown that a large decrease in HA grain
size produces changes in osteoblast activity, including improvements in
attachment and proliferation. (Webster et al, 1999a; Webster et al, 2000a;
Webster et al, 2000b; Webster et al, 2001) Results of this study both contradict
(OB attachment) and reinforce (0B proliferation) these findings. Our study shows
that OB attachment did not improve on scaffolds fabricated from nano HA
powder as compared to those produced from micro HA powder. However, each
of Webster’s studies used dense HA having a sintered grain size of 67 nm

(Webster et al, 20003; Webster et al, 2000b), which is ~11% of our sintered grain

54

size and only 1/3 larger than their unsintered particle size. In contrast, our
scaffolds have grains, which are 30 times larger than our unsintered particles.
This difference in grain growth in our scaffolds, from 20 nm to 588 i 55 nm, is
likely an artifact of the highly porous nature of our scaffolds in which the
predominant sintering mechanism is coarsening. (Barsoum, 1997) Future studies
will investigate the sintering behavior of these foamed nano HA scaffolds with the
goal of reducing grain growth to yield true high volume porosity nanograined HA

scaffolds.

At each time period, 0.5, 1, 2 and 4 hours, hemacytometer analyses
suggest an increase in cell number (attachment) on micro HA scaffolds
compared to nano HA scaffolds. This suggests that increased grain is not
positively related to 0B attachment behavior in highly porous scaffolds. However,
the cell attachment values for both scaffolds are much larger in this study than
expected. This suggests that HA particles themselves may have been included in
the cell counting analyses. Additional approaches are currently being used to

address this issue.

While OB attachment did not increase with decreased grain size, the
mean 0B proliferation, measured at 1, 3 and 5 days, on nano HA scaffolds was
higher than the mean values for the micro HA scaffolds. This indicates that
proliferation may be driven by increased grain boundary area resulting from the
smaller grain size, in addition to the optimum morphology in terms of pore size
and pore interconnectivity (eg macro and microporosity) presented by our highly

porous scaffolds. The basis for OB/HA adhesion and the effect of grain boundary

55

volume on OB behavior are not completely understood, but may be linked in part
to protein interactions at the surface. (Webster et al, 2000b) Webster et al found
that a decrease in grain size promoted an increase in vitronectin and collagen
concentration and a decrease in albumin, laminin and fibronectin concentration.
(Webster et al, 2000b) They also linked the protein stereochemical structure to
grain size and surface topography of bioceramic substrates. (Webster et al,
2000b) Based on these findings, FBS proteins may adsorb more readily to nano
HA scaffolds compared to micro HA scaffolds in this study, but this was not
verified. The effect of FBS protein adsorption in vitro, contrasted with protein

adsorption in an in vivo environment should be investigated.

Other approaches to increasing OB attachment and proliferation on HA
surfaces include peptide coating. For example, EEEEEEEPRGDT coating
increased attachment at 30 minutes, as compared to an HA control. (Itoh et al,
2002) GRGDSPC and cycIo-DfKRG increase osteoprogenitor cell attachment at
3 and 24 hours, as compared to non-treated hydroxyapatite. (Durrieu et al, 2004)
While these coatings increased OB attachment and proliferation, the smaller
grain size found in our HA scaffolds is an intrinsic property, reducing the need for
post-fabrication scaffold treatments such at the peptide coatings of Itoh et al. and

Durrieu et al.

Although the use of nano HA in porous bone scaffolds does not initially
improve OB attachment by 4 hours culture time, the apparent increase in
proliferation of OBs at day 5 shows the potential for more rapid osteoblast

growth, and presumably more rapid calcification and bone formation in vitro.

56

While further studies must be conducted to measure OB differentiation and
calcium production over time in vitro, nano HA porous bone scaffolds, because of
their increased grain boundary area and optimum pore size and interconnectivity,
show promise as an effective bone scaffold for use in bone tissue engineering

applications.

ACKNOWLEDGEMENTS

The authors acknowledge the use of the SEM facilities at the Advanced
Microscopy Center, Michigan State University along with the assistance of Mr J

Quast with the SEM imaging.

57

REFERENCES

Areva, S, Paldan, H, Peltola, T, et al. 2004. Use of sol-gel-derived titania
coating for direct soft tissue attachment. J Biomed Mater Res. 70A:169-78.

ASTM. 2004. E112-96 Standard Test Methods for Determining Average Grain
Size. ASTM International.

Barsoum, M. 1997. Fundamentals of ceramics. McGraw-Hill, New York.

Cerroni, L, Filocamo, R, Fabbri, M, et al. 2002. Growth of osteoblast-like cells
on porous hydroxyapatite ceramics: an in vitro study. Biomol Eng. 192119-24.

Chen, QZ, Lu, W, Wong, CT, et al. 2004. Role of collagen fibres in bone-
bonding of crystalline hydroxyapatite implant. 7th World Biomaterials
Congress, Sydney, Australia. 1439.

Cui, FZ, Zhang, W and Liao, SS. 2004. Hierarchical self-assembling of nano-
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Du, C, Cui, FZ, Zhu, XD, et al. 1999. Three-dimensional nano-HAp/collagen
matrix loading with osteogenic cells in organ culture. J Biomed Mater Res.
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Durrieu, MC, Pallu, S, Guillemot, F, et al. 2004. Grafting RGD containing
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Fratzl, P, Gupta, HS, Paschalis, EP, et al. 2004. Structure and mechanical
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Griffin, MA, Smith, IO and Baumann, MJ. 2005. Comparison between primary
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Hassenkam, T, Fantner, GE, Cutroni, JA, et al. 2004. High-resolution AFM
imaging of intact and fractured trabecular bone. Bone. 35:4-10.

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Hench, LL. 1998. Bioceramics. J Am Ceram Soc. 81 :1705-28.

Huang, J, Best, SM, Bonfield, W, et al. 2004. In vitro assessment of the
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Hulbert, SF, Morrison, SJ and Klawitter, JJ. 1972. Tissue response to three
ceramics of porous and non-porous structures. J Biomed Mater Res. 6:347-
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Hulbert, SF, Young, FA, Mathews, RS, et al. 1970. Potential of ceramic
materials as permanently implantable skeletal protheses. J Biomed Mater
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ltoh, D, Yoneda, S, Kuroda, S, et al. 2002. Enhancement of osteogenesis on
hydroxyapatite surface coated with synthetic peptide (EEEEEEEPRGDT) in
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lyer, W, Kadakia, TB, McCabe, LR, et al. 2004. CCAAT/enhancer-binding
protein-beta has a role in osteoblast proliferation and differentiation. Exp Cell
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Jarcho, M. 1981. Calcium phosphate ceramics as hard tissue prosthetics. Clin
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structural characterizations of synthetic nano-hydroxyapatite and mineral
phase from porcine femoral bone. High technology Letters. 10:48-51.

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698:158-65.

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osteogenesis in porous hydroxyapatite and tricalcium phosphate: a
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59

Perla, V, Ejiofor, JU and Webster, TJ. 2004. Directed osteoblast adhesion at
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60

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61

CHAPTER 4
A BRIEF COMMUNICATION REGARDING THE INACCURACIES OF
HEMACYTOMETER MEASUREMENTS WHEN ASSESSING OSTEOBLAST
ATTACHMENT ON POROUS NANO- AND MICRON-SCALE
HYDROXYAPATITE SCAFFOLDS
IAN 0. SMITH‘, LAURA R. MCCABE2 AND MELISSA J. BAUMANNI
1Dept Chemical Engineering and Materials Science, Michigan State University,

East Lansing, MI, USA
2’Dept PhysiologY, Michigan State University, East Lansing, MI, USA

INTRODUCTION

In our paper entitled “MC3T3-E1 osteoblast attachment and proliferation
on porous hydroxyapatite scaffolds fabricated with nanophase powder”,
published in the lntemational Journal of Nanomedicine‘, we found that porous
scaffolds fabricated using nano-scale hydroxyapatite (HA) powder Significantly
improved osteoblast (OB) proliferation when compared to HA scaffolds fabricated
with micron-scale powder after 5 days in vitro. However, we also Showed that OB
attachment on nano-HA porous scaffolds was Significantly lower on the nano-HA
powder scaffolds in comparison to the micron-HA porous scaffolds. This finding
is contrary to those Shown in multiple studies by Webster et al., who found that a
large decrease in grain size led to increased OB attachment and proliferation“.
This apparent discrepancy led us to examine the accuracy of the techniques
used in our previous study, namely OB counts using a hemacytometer.

In this letter, we address the accuracy of hemacytometer counts for
porous HA scaffolds and use a Simple method to reduce the error. The

hemacytometer relies on transmitted light to identify cells. In order to trypsinize

62

and count the cells, we had to morselize the porous scaffolds. Therefore, despite
our best efforts to eliminate extraneous morselized scaffold particles, the HA
scaffold particles were commingled with the OBS and, because these particles
reflected light in a Similar manner and were of approximately the same size, it
was difficult to discern between an OB and an HA morsel. By employing a
background count to quantify the number of HA particles which could be
statistically generated during the sample processing procedure, we can subtract
these values from the raw counts leaving us with a normalized and corrected OB

cell count.

MATERIALS AND METHODS

Cylindrical porous HA scaffolds were fabricated using the foaming
technique described in earlier studiesI'6 Sintered HA Specimens were
approximately 2 cm in length and 1.25 cm in diameter with an average porosity of
59 i 2.2% (micron-HA) or 72 i 8.3% (nano-HA), determined by Archimedes
method. These HA scaffolds had an average grain diameter of 8.6 i 1.9 pm for
micron-HA and 588 i 55 nm for nano-HA, with a bimodal pore distribution
(macro/micropore diameter of 390 i 210 um/38 :t 33 pm for micron-HA and 480 :l:
275 um/30 + 21 pm for nano-HA) determined by scanning electron microscopy.1
These cylindrical specimens were sectioned into 2 mm-thick discs using a
Buehler Isomet 1000 diamond abrasive wafering saw. Three micron-HA discs
and three nano-HA discs were used at each time point and the entire study was

run in triplicate, requiring a total of 63 micron-HA discs and 63 nano-HA discs.

63

Prior to 08 seeding, scaffolds were autoclave sterilized and placed into
commercially pure Al collars.
MC3T3-E1 subclone 7 murine osteoblasts (OBS) of passage 26-28 were

used for this study. OBS were cultured on polystyrene (PS) tissue culture plates,
with complete media (0t minimum essential media [Ct-MEM] supplemented with

10% fetal bovine serum [FBS]) until a suitable population size was reached. Prior
to seeding onto the porous HA scaffolds, the OBS were removed from the PS
using trypsin-EDTA and counted with a hemacytometer. The OBS were then
placed into 24-well plates, one per well and the OBS seeded onto the scaffolds at
11,320 cells/cm2 to ensure confluency7 with 2 ml of complete media added to
each well. These Specimens were then incubated at 37°C, 5% CO2 and high
humidity.

OB attachment was assessed at 0.5, 1, 2 and 4 hours and 0B proliferation
measured at 1, 3 and 5 days. At each time interval, the scaffolds were removed,
rinsed with 1X phosphate-buffered saline (PBS), to remove any non-adherent
OBs, and placed onto fresh 24-well tissue culture plates, one per well. One ml
each of trypsin-EDTA and complete media were then added to each well.
Scaffold homogenization was used to assist in the release of OBS from surface
during trypsinization.8 Each scaffold was morselized using the handle-end of a
cell scraper (Benton Dickenson) to encourage OB detachment. Samples of
suspended OBS were then drawn from each well and the cells counted using a
hemacytometer. The student's t-test was used to determine statistical difference

at each time point, with p < 0.05 taken to indicate a significant difference.

64

Because the remnants of the morselized scaffold can be of a Similar size
and shape as the OBS when viewed under the hemacytometer (Figure 1),
background counts on morselized HA scaffolds which were not seeded with OBS
were also conducted. Three micro-HA specimens and three nano-HA Specimens
(run in triplicate for a total of 9 each) were submerged in complete media,
removed, rinsed with 1X PBS, trypsin was added and the scaffolds morselized.
Media was drawn from each well and any cell-like particles were counted using a
hemacytometer. These background counts were then subtracted from the raw

hemacytometer counts at each time interval.

65

 

Figure 1: Figure 1A is an optical micrograph of a hemacytometer with a single OB (indicated by
circle) taken during a cell count where there were no ground HA particles. Figure 1B is an optical
micrograph of a hemacytometer with a morselized OB-seeded micron-HA scaffold. Arrow 1
indicates an HA particle, arrow 2 indicates an OB and arrow 3 indicates an object that is not
definitively one or the other. The general haze resulting from the ground particles, as well as the
presence of larger OB-sized particles make it difficult to identify OBS.

66

RESULTS AND DISCUSSION

Figure 2 is a plot of OB attachment (0 — 4 hours) and proliferation (24 -
120 hours) versus time with, and without, the background HA count subtracted
away. The data shows that the background-adjusted OB attachment on micron-
HA porous scaffolds increases from 66 i 259% at 0.5 hours to 193 i 327% at 4
hours. In comparison, OB attachment increases to a lesser degree on nano-HA,
from 0% to 28 i 155% between 0.5 and 4 hours. However, OB proliferation
increases much more noticeably, from 239 i 261% to 1004 :t 697% on micron-
HA and from 239 1- 261% to 1044 i 301% on nano-HA. These results are
different than those reported in our earlier work, in that there is no Significant
difference in OB attachment or proliferation on nano-HA porous scaffolds
compared to micron-HA porous scaffolds.1

The non-adjusted data is consistent with our earlier findings Showing OB
attachment to be significantly higher (p < 0.05) for micron-HA porous scaffolds
with no significant difference (p > 0.05) in OB proliferation between 1 and 3 days
for nano-HA relative to micron-HA porous scaffolds. The non-adjusted data does
differ from the earlier data in that at 5 days OB proliferation is not Significantly

higher on nano-HA scaffolds compared to micron-HA scaffolds.

67

1400 ~

+ nano HA adjusted .
1200 - —-I— micron HA adjusted ‘ , . ' i I

--O--nanoHA .' '
1000 - - - I - - micron HA

 

% Cells Attached

 

 

0 20 40 60 80 100 120
Time (hrs)

Figure 2: MC3T3-E1 OB attachment (0.5, 1, 2 and 4 hrs) and proliferation (24, 72 and 120 hrs) on
porous nano-HA and micron-HA scaffolds.

Both the nano-HA and micron-HA scaffolds promote OB proliferation at
times approaching five days. This indicates that the controlling factor in whether
a porous scaffold potentially promotes bone growth may be porosity, rather than
grain size. This finding is contrary to those shown at length by Webster et al. and
to a lesser extent by Xie et al. in studies of OB behavior on dense ceramic
scaffolds, including HA, which showed that decreased grain Size improved OB
attachment, proliferation, differentiation, AP activity and calcium production""5 as
well as enhanced gene expressiong. However, in each of these cases the
scaffolds studied were highly dense, meaning that grains and surface roughness
were the only topographical features present. In comparison, a bimodally porous

scaffold, such as those used in the current study, more closely mimics the

68

structure of natural bone”12

and presents a more biomimetic interface for OB
interactions. These findings further strengthen the argument that bimodal
porosity, similar to that found in natural bone, is a primary factor in promoting

bone ingrowth. Bone tissue engineering scaffold materials, including HA, must

include both large and small pores, which allow ingrowth and nutrient flow.

CONCLUSIONS

The purpose of examining hemacytometer counting techniques for
assessing OB attachment and proliferation on porous HA scaffolds was to
determine whether inherent problems with this technique could be corrected by
utilizing background counts. By using a simple background count, where
scaffolds were ground and OB-like objects were counted, we attempted to
correct for any inaccuracy in hemacytometer values. While this method reduced
interference in hemacytometer counts due to the inclusion of HA particles, the
results indicate further inaccuracies. Results show roughly a doubling of the OB
population by 10 hours, which is not typical on HA.7 This high number of OBS is
likely due to counting non-adherent cells, which were trapped in the porous
matrix and not effectively rinsed away. To avoid these problems, a quantifiable
stain could be used in lieu of hemacytometer counts, and a more thorough
rinsing regimen used.

Inaccurate OB counts from hemacytometer data for porous scaffolds is
currently unavoidable. The scaffolds used here are porous in nature as opposed

to the dense scaffolds used in past studiesz's' 9, which have looked solely at the

69

effect of grain size and not the role of porosity in OB attachment. Hence, these
measurements have been made on dense samples that have not needed to be
morselized. It is not surprising then that our results from porous scaffolds are
different when compared to the dense scaffolds used in these earlier studies.
While using dense samples makes cell attachment measurements more
straightforward, the dense surface of these samples is less biomimetic and not

applicable in many in vivo applications.

70

 

REFERENCES

1.

Smith IO, McCabe LR, Baumann MJ. MC3T3-E1 Osteoblast attachment
and proliferation on porous hydroxyapatite scaffolds with nanophase
powder. Int J Nanomed 2006;1:189-194.

Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R. Nanocrystalline
hydroxyapatite enhances osteoblast function. In: Proceedings of the First
Joint BMES/EMBS Conference, Atlanta, GA; 1999. p744.

Webster TJ, Siegel RW Bizios R. Design and evaluation of nanophase
alumina for orthopaedic/dental applications. Nanostruct Mater
1999;12:983-986.

Webster TJ, Siegel RW Bizios R. Osteoblast adhesion on nanophase
ceramics. Biomaterials 1999;20:1221-1227.

Webster TJ, Ergun C, Doremus RH, Siegel SW, Bizios R. Enhanced
functions of osteoblasts on nanophase ceramics. Biomaterials
2000;21:1803-1810.

McMullen, R. An in vitro investigation of MC3T3-E1 osteoblast proliferation
and differentiation on hydroxyapatite based tissue-engineered scaffolds.
Masters Degree Thesis. Department of Chemical Engineering and
Materials Science. Michigan State University, East Lansing, MI; 2004.
181p.

Shu R, McMullen R, Baumann MJ, McCabe LR. Hydroxyapatite
accelerates differentiation and suppresses growth of MC3T3-E1
osteoblasts. J Biomed Mater Res. 2003;67A:1196-1204.

Lu HH, Tang A, Oh SC, Spalazzi JP, Dionisio K. Compositional effects on
the formation of a calcium phosphate layer and the response of
osteoblast-like cells on polymer-bioactive glass composites. Biomaterials
2005;26:6323-6334.

Xie J, Baumann MJ McCabe LR. Osteoblasts respond to hydroxyapatite
surfaces with immediate changes in gene expression. J Biomed Mater
Res. 2004;71A:108-117.

10. Hench LL.. Bioceramics. J Am Ceram Soc 1998;81:1705-1728.

11.Hulbert SF, Morrison SJ, Klawitter JJ. Tissue response to three ceramics

of porous and non-porous structures. J Biomed Mater Res 1972;6z347-
374.

71

12.Hulbert SF, Young FA, Mathews RS, Klawitter JJ, Talbert CD, Stelling FH.
Potential of ceramic materials as permanently implantable Skeletal
protheses. J Biomed Mater Res. 1970;4:433-456.

72

CHAPTER 5

MICROCRACKING AND POROSITY IN CALCIUM PHOSPHATES AND THE
IMPLICATIONS FOR BONE TISSUE ENGINEERING

E.D. CASE, I.O. SMITH AND M.J. BAUMANN
Department of Chemical Engineering and Materials Science,
Michigan State University, East Lansing, MI 48824-1226 USA

PUBLISHED IN: Mater Sci Engin A 2005 390:246-254.

ABSTRACT

Calcium phosphates including hydroxyapatite (HA) have been widely
studied as bone scaffold materials. However their mechanical properties are
highly variable and may be a function of thermal expansion anisotropy (TEA)
induced stresses and microcracking. There is confusion concerning the
mechanisms and microscopic identification of microcracks in HA. This study
presents clear evidence of microcracking from micrographs of aS-sintered HA
surfaces which avoids the complications of identifying TEA-induced microcracks
on fracture surfaces. Additionally, the existing literature of TEA-induced
microcracking iS briefly reviewed and pertinent papers involving likely
microcracking in HA are analyzed. The recent realization in the biomedical
literature notes that microcracks are of critical importance in remodeling and
repair of damaged bone tissue. Hence, the similarities between microcracking in
HA used for scaffolds in bone tissue engineering and that in the normal bone
repair process is of importance in designing HA scaffolds with improved

mechanical properties and biocompatibility.

73

INTRODUCTION

When a material exhibits thermal expansion anisotropy (TEA), the
coefficient of thermal expansion is a function of crystallographic direction. TEA
occurs in all crystalline systems of lower symmetry than cubic [1]. In general, any

material property described by a second rank tensor (such as the thermal
expansion coefficient, 00]) is isotropic in an amorphous material and in a cubic
atomic lattice, but for tetragonal, hexagonal, orthorhombic, monoclinic and
triclinic systems, 00,- is no longer isotropic.

Hydroxyapatite (HA) often crystallizes in a hexagonal lattice with space
group P63/m [2, 3], although HA also may crystallize in a monoclinic lattice with
space group P21/b [4]. Thus, in either the hexagonal or monoclinic form, HA
exhibits thermal expansion anisotropy. However, the important practical effects of
TEA arise in polycrystalline bodies, since as polycrystalline materials cool from
elevated temperatures, thermal expansion anisotropy results in contraction of
individual grains from neighboring grains at differing rates, depending upon
crystallographic direction. These differences in thermal contraction rates give rise
to mechanical stresses.

The physical consequences of the thermal expansion anisotropy are

mainly determined by the magnitude of the thermal expansion differences along

the crystallographic axes, say on, O2 and 03. If the differences among the (11, a2

and a3 values are sufficiently large, the TEA driven stresses can generate

microcracks.

74

In this study, we briefly review the concept of TEA-induced microcracking
and the critical grain size for microcracking. We also discuss the biomedical
implications of microcracking in bone and HA and the concomitant implications
for bone tissue engineering. Finally, we present a series of micrographs of dense
HA Specimens and carefully consider studies in the literature that deal with

microcracking in HA.

BACKGROUND
Microcracking and critical grain size

An energy balance between a volumetric energy term (the stored elastic
strain energy) and an area energy term (the surface energy corresponding to the
newly created crack surface area) gives rise to a critical grain Size for TEA-
induced microcracking in brittle polycrystalline materials. If the specimen grain
size is G, then the condition that G > GCR leads to microcracking in the Specimen
while if G < GCR, then the Specimen does not microcrack. The critical grain size,

GCR, can be estimated by [5, 6]

_ 14.4y,
E(AT)2 (Aa

 

(1)

CR .05)
where Yf is the fracture surface energY. E is the Young's modulus of the material,

, and Aamax is the maximum difference in the values of the thermal expansion
coefficients along the crystallographic axes. The effective temperature interval

over which TEA stresses accumulate in the specimen is given by AT. At

75

elevated temperatures, stresses can relax in the specimen, but during cooling the

specimen will reach a temperature, TU, at which the intergranular stresses no
longer relax. AT is thus given by (TU - TL), where TL is the low temperature limit of

the heating cycle (typically TL is room temperature).
The value of the critical grain size for thermal expansion microcracking

varies from ceramic to ceramic. For example, literature values of GCR for ceramic

oxides span nearly two orders of magnitude with a GCR of approximately 100 pm

for alumina [7], about 15 pm for gadolinum oxide [8], and roughly 3 pm for
FeTi205 [6]. As is predicted by equation 1, the trend in GCR tends to be dominated

by Aamax, with GCR decreasing rapidly as the degree of thermal expansion

anisotropy (measured by Aamax) increases.

In addition to the existence of a critical grain size, a crucial aspect of TEA-
induced stress and microcracking is related to the mechanical behavior based on
whether G is smaller or larger than GCR. As an example, consider grain Size G31,
where G51 << GCR. The residual stresses due to TEA would be relatively minimal
in this case.

However, as the grain size increases toward GCR, the local stresses can
significantly alter the strength and fracture toughness of the material [9, 10]. If the
grain size increases to G52 where G32 is Slightly smaller than GCR, then ideally the
specimen would not microcrack but the residual stresses induced by TEA could
be quite large and would in fact become sufficient to generate microcracks if the

grain size is increased Slightly, such that 652 becomes equal to GCR [10, 11].

76

For a grain size GL1 slightly larger than the critical grain size, microcrack
damage begins to accumulate in the Specimen [11, 12]. If the grain size is further
increased such that GL2 > GCR, then the observed magnitude of TEA-induced
microcrack damage increases as the grain Size increases [7, 8, 13, 14]. For G >
GCR, the resulting microcracking in polycrystalline specimens impacts the fracture
behavior of the specimen, causing either a net increase or decrease in fracture
toughness[15,16]

In addition to changes in strength and fracture toughness when G > GCR,
volumetrically distributed microcrack damage within the specimen leads to
changes in a wide variety of material parameters, including the elastic modulus

[6-8, 10, 17] and thermal conductivity [18, 19].

Microcracking in Bone

It is well accepted that fatigue injuries such as stress fractures result in
microcracking in the bone. There is in turn ample evidence linking these
microcracks to bone healing and remodeling in both the human and veterinary
literature [20, 21]. The ability of bone to microcrack improves its toughness by
absorbing energy, thereby preventing or Slowing down the propagation of a
failure-inducing macrocrack [22, 23].

Examining the distal ends (the end furthest from the body) of the third
metacarpal and metatarsal bones of Thoroughbred racehorses with condylar
fractures, Radtke found that macrofracture and failure of the bone was correlated

to the presence of microcracks [24]. In a separate study investigating the fatigue

77

damage in the greyhound central tarsal bone, Muir notes the presence of both

linear microcracks (~30-110 pm in length) and clusters of ultra-microcracking

(<10 pm in length) [21]. Linear microcracking was found to exist around areas of

remodeled cortical and trabecular bone adjacent to a vascular supply, while ultra-
microcracking under tensile loading was found at the attachment Site of the
plantar ligament.

While the exact functional relationships are not known between fatigue
and stress fractures, bone remodeling and repair have been associated with the
initiation of microcracks arising from fatigue damage and altered osteocytes
activity in a study by Bentolila in adult rat long bones [25]. In humans, Burr et al.
find that microcrack density increases with age, where the microcrack damage is
initiated at the collagen fiber to bone matrix attachment sites and rapidly
accelerates after 40 years of age [26].

There is a proven Significant increase in remodeling Sites that develop
after the initiation of microcracking. However, if fatigue damage continues
concomitant with bone remodeling (which involves an increase in the local bone
porosity), the resulting loss of mechanical properties, such as strength and
stiffness, likely is a cause of further microcracking. This would in turn, stimulate
further remodeling and a loss of mechanical integrity until fracture finally occurs.
Such a site specific response to fatigue damage is only just gaining acceptance
in the biological community and has implications for understanding fracture and

healing in the osteoporotic elderly and in human and animal athletes [27, 28].

78

HA in Bone Tissue Engineering

Each year ~300,000 bone grafts are performed in the US. [29].
Applications include bone loss from neoplasia, fracture repair, Spinal surgery,
and fracture treatment in the elderly. The model bone substitute is 1) osteogenic
(direct bone formation via transplanted bone cells called osteoblasts, 2)
osteoinductive (the ability to recruit bone forming cells), 3) osteoconductive (the
scaffold supports bone in-growth), 4) mechanically stable and 5) readily
available. Although autografts (bone transplants from the patients own bone
supply) and allografts (same species transplants from a banked supply of bone)
seem ideal to repair bone defects, they are limited in supply, carry additional
surgical/patient risks and possess diminished properties [30, 31].

Because the mineral content of bone has a composition and crystal

structure closely matching HA, porous calcium phosphates like HA (non-

resorbable in vivo [32]), B-TCP (resorbable in vivo [32]) and BCP (a mixture of

HA and B-TCP) are being investigated for use in bone tissue engineering

applications. Dense calcium phosphates cannot support osteoblast in-growth or
bone formation while porous calcium phosphates are osteoconductive, but lack

the strength of dense HA.

EXPERIMENTAL PROCEDURE
Commercial HA and B-TCP powdersI, of vendor-Specified medical grade

purity, were obtained and used as manufactured. BCP powder was obtained by

 

I Berkeley Advanced Biomaterials Inc., San Leandro, Ca., USA

79

mixing 60% HA (particle diameter of 0.85 pm) with 40% B-TCP (particle diameter
of 1.15 pm) powders by volume. This mixture was then shaken for one minute
and thoroughly mixed by placing the container in an ultrasonic bath for 30
minutes. Green disc specimens were fabricated by uniaxially pressing 1.5 1:
0.002 g of dry powder in a 32 mm diameter pellet die at 6.9 MPa for one minute
using a standard laboratory press. Discs were Sintered in air for four hours at
either 1360°C (for HA) or 1200°C (for B-TCP and BCP) in a high temperature
sintering furnace (CM Rapid Temp Furnace) by heating from room temperature
to the sintering temperature at a rate of ~10°Clmin. Following sintering, samples
were cooled to room temperature by a controlled furnace cool at a rate of
~10°Clmin. The pore volume was determined using Archimedes method and the
pore and grain Size determined using the line-intercept method.

To determine phase composition and purity, X-ray diffraction (XRD) was

performed on (1) unsintered powder Specimens from each of the three starting
powders (HA, B-TCP and BCP) as well as (2) powder specimens obtained by
grinding the sintered disc specimens. All powders were mounted on a

borosilicate glass slide and scanned in a powder diffractometer over the 20

range from 20° to 50° using filtered CU-KO. radiation. Sintered discs were ground

with an agate mortar and pestle and the resulting powder was sieved to remove

large particles.

Sintered Specimens used for grain Size analysis were polished to 1 pm

with diamond paste followed by ultrasonic cleaning and thermal etching at

80

1200°C for four hours. Additional sintered discs were prepared for examination in
the scanning electron microscope (SEM) by Slicing each disc perpendicular to
the circular face using a low speed diamond wafering saw or by fracturing the
discs by the impact of a steel wedge on the specimen surface. The specimens
were sectioned in order to produce samples small enough to fit into the SEM
chamber. Specimens were cut using a Beuhler Isomet 1000 low speed saw with
a diamond wafering blade at 200 rpm. This dry procedure was used to suppress
possible microcracking or microcrack extension induced by liquid immersion. The
surfaces were then prepared for examination in the SEM by mounting sections of
each material on an Al stub with C adhesive tape and gold coating using an
EMSCOPE 80500 Sputter Coater operated at 20 mA for 3 minutes resulting in a
gold coating of 21 nm. SEM micrographs were taken at an accelerating potential
of 15 KeV and a working distance of 15 mm on a JEOL JSM6400 SEM equipped
with a LaBe emitter. During imaging, care was taken to ensure that micrographs
were imaged of representative areas, at least 2 mm away from the cut or

fractured edges.

RESULTS AND DISCUSSION
Microstructure of calcium phosphate specimens

Representative SEM photomicrographs of the as-sintered, polished and
thermally etched surfaces of HA, B-TCP and BCP specimens are shown in Figs.
1 to 3. The line-intercept method was used to assess the grain Size of these

materials yielding values of 7.9 i 1.5 pm for HA, 2.9 i 0.6 pm for B-TCP and 1.3

81

i 0.1 pm for BCP. The density of the samples was measured using Archimedes
method resulting in densities of 97.7% .+_ 0.7%, 63.9% i 3.2% and 69.1% i 1.4%
for HA, B-TCP and BCP, respectively. The trend in relative density and grain size

in the CaP specimens is consistent with the general observation that during
processing Iof ceramic powders, densification typically proceeds with limited grain
growth up to a relative density of about 0.9 and at higher relative densities, grain
growth accelerates [33—36].

Broad peaks in the XRD patterns (data not shown) for the unsintered HA,
BCP and B-TCP powders likely are due to the small particle size of the starting
powders. Diffraction peaks were much sharper for the sintered specimens. The

minor phases observed were CI-TCP and B-TCP for the HA discs and OL-TCP and
HA for the B-TCP discs. For BCP, minor phases of HA and B-TCP, and 0t-TCP

were present. The HA/B-TCP discs were determined to be 61 % HA using

 

I
%HA : IOOJ‘IA #100 (2)

IIOOJIA + IOO,fl—TCP

where I100, HA and I1oo,p-mp are the intensities of the maximum intensity peaks of

HA and B-TCP, respectively [37].

82

 

Figure 1: SEM micrograph of polished and thermal etched hydroxyapatite (5000x).

Microcracking in HA specimens

Porosity can suppress TEA-induced microcracking in ceramics when the
volume fraction porosity (VFP) exceeds roughly 20% - 30% [14]. Since the [3-
TCP and BCP specimens included in this study have a VFP of about 31% and
39%, respectively, it iS unsurprising that microcracks were not observed in these
calcium phosphates.

In contrast, the SEM micrographs of the aS-Sintered HA surfaces (Figs. 4-
6) clearly Show microcracking. The microcracking is abundant and not correlated
with areas of porosity, rather it is uniform across the HA surface. A careful
examination of the micrographs illustrates that the microcracks are associated
with what appear to be separated grain boundaries, emerging from the grain

boundary corners to continue into adjacent grains as noted by the arrows. The

83

maximum crack opening displacement is roughly 0.35 pm with an average

microcrack length of approximately 12.5 :I: 9.9 pm.

 

Figure 3: SEM micrograph of polished and thermal etched BCP (3000x).

84

For the aS-Sintered surface of the HA specimens sectioned by fracture
(Fig. 4), microcracks transverse the grains (white arrows) and run along the grain
boundaries (black arrows). Some of the microcracks that transverse grains
appear to be connected to the grain boundary microcracks. For the as-sintered
surface of the HA specimens sectioned by cutting (Figs. 5 and 6), the microcrack
damage was similar to that observed in Fig. 4 for the specimen sectioned by
fracture. Figs. 5 and 6 represent two different areas on a Single specimen.
Therefore, the microcrack damage appears to be independent of sectioning

technique and in each of Figs. 4-6, the microcrack damage is quite similar.

 

Figure 4: SEM micrograph of as-sintered hydroxyapatite specimen surface. Image was taken
away from the fracture surface (6500x).

85

 

Figure 5: SEM micrograph of as-sintered hydroxyapatite specimen surface shown in Fig. 5.
Image was taken away from the fracture surface (4500x) and at different location on the
Specimen.

 

Figure 6: SEM micrograph of as—sintered hydroxyapatite specimen surface. Image was taken
away from the fracture surface (2500x).

86

A re-evaluation of key results from the literature on microcracking in HA
AS discussed by Hoepfner and Case [38], the critical grain Size for TEA

induced microcracking in HA is difficult to estimate, in large part due to the large
uncertainty in the value of Aamax that arises from a broad range of single crystal
thermal expansion values found in the literature. From Eq. (1), we see that the

critical grain size is proportional to 1/(A0Lmax)2. Thus, as reviewed by Hoepfner

and Case [38], the available x-ray diffraction data yields Aamax values for HA that

vary by as much as a factor of 10, and therefore it is difficult to make a precise

estimate of HA'S critical grain size. (Hoepfner and Case [38] base an estimate of
GCR on the value of Aamax that is most consistent with the average thermal

expansion values.)

For the HA literature, damage that has been attributed to other
mechanisms can perhaps be attributed to TEA-induced microcrack damage. For
example, in a study of Sintered HA, Ruys et al. [39] observed that as the
specimen density increased from that of the original powder compact to a relative
density of ~ 0.95, the specimens' fracture strength increased monotonically.
However, beginning at a relative density of about 0.95, the fracture strength
dropped with further increases in density. Ruys et al. [39] attributed the observed
strength decrease to dehydroxylation of the HA (where dehydroxylation refers to
the loss of hydroxyl groups during heating, which can lead to the decomposition
of HA to oxyhydroxyapatite or in the extreme, to TCP [2].)

To evaluate the concept advanced by Ruys et al. [39] that dehydroxylation

during sintering was responsible for the microcracking observed in their HA

87

Specimens, it is critical that we examine several aspects of sintering, including
the generation of volatiles during sintering and related questions of the evolution
of pore Shape, pore size and grain Size during Sintering. In general, it iS true that
if volatiles are generated during Sintering (via mechanisms such as thermal
decomposition [36, 40] or debinding [41]) then those volatiles must be removed
before local pore closure commences, othenivise pressure from the encapsulated
gases can result in cracking, bloating and blistering of the sintered component
[36,40,41]

However, it is important to note that the initiation of local pore closure is
just one step along the evolutionary path experienced by the ceramic compact as
pore Shape, pore Size and grain Size change during sintering [33-36]. In

particular, the process of local pore closure begins early in the intermediate stage
of sintering (at relative density, p, values of approximately 0.7 to 0.8 [33]) as
tubular pores along the grain edges begin pinching off into Closed, spherical
pores. When the sintered component reaches a p value of about 0.9 to 0.95, the

intermediate stage of sintering is complete and essentially the entire inventory of
pores within the specimen has pinched off into isolated Spherical pores [33-36].
Thus, for the sake of interpreting the experimental results of Ruys et al.

[39], it is critical to distinguish between (1) the beginning of the pore pinch-off

process (at roughly 0.7 < p < 0.8) which can lead to cracking due to entrapped

gases and (2) the end of the pore pinch-off process at p values of about 0.9 to

0.95. If dehydroxylation did spawn the microcracks in the HA Specimens studied

by Ruys et al. [39], then the microcracking (and possibly blistering and bloating of

88

the Specimens) would have likely commenced at the relatively modest p values

of 0.7 to 0.8, rather than at the much higher p value of 0.95 at which Ruys et al.

[39] actually observed both the microcracking and the associated drop in
strength.

These considerations of pore pinch-off and microcracking beg the
following question, "If dehydroxylation did not cause the microcracking observed
by Ruys et al. [39] for their HA Specimens with densities > 0.95, then what did”?
The answer likely is again found within the context of microstructural evolution

during sintering, namely the widely observed phenomenon that grain growth
accelerates when the Sintered compact reaches a p value of about 0.90 to 0.95

[42-45].
As depicted in Fig. 7, during the initial and intermediate stages of
sintering, considerable densification occurs (especially during the intermediate

stage). Despite the densification, very little grain growth occurs until the end of
the intermediate stage. Thus for densities below p ~ 0.90, grains in the

densifying specimen are roughly comparable in size to the powders that made up
the original powder compact [44]. Only at the onset of the final stage of sintering
(at a relative density of ~ 0.90 - 0.95) does grain growth accelerate dramatically
(Fig. 7). This behavior was been widely observed in metals and ceramics, for
example, in powder compacts of TiO2 [42], Al203, BeO, ZnO and copper [43],

MgO-doped alumina [44] and Silica doped, ultrapure Al203 [45].

89

 

Grain Size

PGG

 

 

 

PINT Density (%)

Figure 7: Schematic showing the relationship between relative density and grain size for ceramic
powders.

Thus, in the study by Ruys et al. [39], at a relative density of about 0.95
the grain growth process likely accelerated such that the grain size G became
larger than the critical grain Size GCR. In general, for G > GCR, microcracking
ensues and the strength and fracture toughness drop, which also is consistent
with the particular observations recorded for HA by Ruys et al. [39]. Given these
considerations of pore/grain Size evolution during Sintering, it is likely that the
microcracking observed in HA by RuyS et al. [39] was due to TEA-induced stress
rather than dehydroxylation.

In addition to the Ruys et al. paper [39], another paper from the literature
that is critical to examine is one by Halquani et al. [46] concerning microcracking
in HA. Although Halquani et al. [46] made strength and fracture toughness
measurements on HA specimens that are consistent with the presence of

microcrack-induced damage, there are some questions associated with their

90

microscopy of the surface features they identified as TEA-generated
microcracks.

Via hot pressing, Halquani et al. [46] obtained nearly-dense polycrystalline

HA specimens with grain sizes ranging from roughly 0.15 -1.2 pm. Two grain-
size-dependent regimes of mechanical behavior were observed. For G < 0.4 pm,
the fracture strength, 0, increased monotonically from O ~ 75 MPa (grain size of
0.15 pm) to a maximum strength of 0' = 137 MPa (grain Size of 0.4 pm). Thus,
the strength (6) increased by about 83% over the grain Size interval from 0.15
pm to 0.4 pm. In contrast, for G > 0.4 pm, 6 decreased monotonically from the
maximum at G = 0.4 pm to a 0' of ~100 MPa at G = 1.2 pm. Also at a grain Size

of 0.4 pm, the fracture toughness, Kc, passed through a maximum, although the

maximum Kc value of 1.2 MPam“2 was only roughly 20% larger than the Kc
values at the smallest and the largest grain sizes.

Halquani et al. [46] interpreted the maximum in both fracture strength and
fracture toughness as being due to TEA-induced microcracking with a critical
grain Size of about 0.4 pm. However, Halquani et al. [46] only identified structures
they interpreted as being microcracks on the fracture surfaces of their HA
specimens. Furthermore, the microstructural features that Halquani et al. [46]
identify as microcracks in fact seem very similar to cleavage steps that are nearly
always present on the transgranular fracture surfaces in brittle materials.

Furthermore, Halquani et al. [46], state that microcracks in their HA specimens

91

are difficult to observe since they "can only be seen when perpendicular to the
rupture plane on a mirror area (i.e. on regions of transgranular failure)". Thus, it is
possible that the features identified as microcracks by Halquani et al. [46] may
actually be cleavage steps.

In contrast to the Halquani et al. study [46], in this study, the micrographs
that document the microcracks in our HA specimens (Figs. 4-6) were taken on
the aS-sintered specimen surfaces, thus avoiding the possible complication of
fracture-surface artifacts. Thus, the micrographs presented in this study (Figs. 4-
6) are the first unequivocal evidence for microcracking in polycrystalline HA to
appear in the literature.

If we continue to examine the observations of Halquani et al. [46], we find
that the magnitude of the fracture toughness increase in HA observed by
Halquani et al. [46] iS consistent with other researchers’ theoretical calculations
and experimental observations of TEA-stress and microcracking induced
changes in ceramics. For example, for Specimens with grain sizes smaller than
the critical grain size, Shum [47] calculated that stresses produced by thermal
expansion anisotropy and elastic anisotropy stresses can increase the fracture
toughness by up to about 1.3 or perhaps as much as 3 in some cases. In Samo
and Tomozawa's [48] study of a TEA-induced microcracking in a zirconia-lithium
aluminosilicate glass ceramic , the experimentally-determined increases in
fracture toughness are of the same magnitude that Shum predicts for G < GCR.
As was the case for the HA specimens studied by Halquani et al. [46], for G >

GCR the Samo and Tomozawa [48] glass ceramic materials experienced a

92

dramatic decrease in toughness accompanied by a fracture strength that
dropped to essentially zero. Similar behavior (that is a rise in strength and
fracture toughness for G < GCR and a drop for G > GCR) also has been observed
for TEA-induced behavior in alumina, TiO2, iron titanate, and Nb205 [15].

In addition to TEA as the mechanism for inducing microcracks in CaP
materials, two additional mechanisms to be considered are (i) thermal Shock [10,
49] and (ii) thermal expansion mismatch [10, 50, 51]. For thermal shock, the
stresses required for microcracking are generated by a rapid change in the
Specimen’s ambient temperature [10, 49]. In this study, the rate of change in
ambient temperature is relatively modest, in that the most rapid temperature
changes experienced by the Specimens occur during processing when the
specimens are heated at ~10°Clmin up to the Sintering temperature, followed by
cooling, again at about 10°C/min to room temperature. As calculated by Halquani
et al. [46], our cooling rate is not sufficient to induce the microcracking observed
in the HA specimens included in this study.

Additionally, microcracking may be generated by thermal expansion
mismatch, which is similar to TEA, except that thermal expansion mismatch
involves the differing coefficients of thermal expansion between two or more
phases in a given material [10, 51]. Thermal expansion mismatch microcracking
has been both modeled theoretically [51] and observed experimentally in a
variety of ceramic systems [50, 52-54]. To investigate the possibility that thermal

expansion mismatch caused microcracking in, for example, HA specimens with a

minor B-TCP phase, we can examine studies of BCP (biphasic calcium

93

 

phosphate - HA/B-TCP) materials. As an example, for hot-pressed BCP (90 wt%

HA/10 wt% B-TCP) specimens having a homogeneous distribution of 8-TCP

grains within the HA matrix, Raynaud et al. [55] obtained a fracture strength of
150 MPa, which was double the strength Raynaud et al. measured in the same

study [55] for similarly prepared pure HA specimens. Thus, Raynaud et al.'s [55]
two-phase mixture of HA/B-TCP was obviously not badly microcracked,

compared to the pure HA. Otherwise, the BCP strength would have been lower

than that of the strength of the pure HA specimens. Furthermore, based on the x-

ray diffraction results, the B-TCP phase fraction was likely less than 10% for the

HA specimens included in this study. Thus, for HA specimens with a minor 8-

TCP phase (such as those in this study), thermal expansion mismatch induced
stresses are not likely to result in significant microcracking,

Since thermal Shock, thermal expansion mismatch and dehydroxylation
are unlikely sources of microcracking for the specimens in this study, and
because HA is known to crystallize in a non-cubic atomic lattice, the likely source

of the microcracks observed in our specimens is therefore TEA.

CONCLUSIONS AND SUMMARY

HA is an important biomedical material. Although HA should display
thermal expansion anisotropy, the existence and effects of TEA-induced
microcracking are somewhat uncertain based on the current literature. For

example, microcracking in HA was attributed to dehydroxylation by Ruys et al.

94

 

[39], although given considerations of microstructural evolution; the observations
are more consistent with microcracking due to TEA. Also, while Halquani et al.
[46] presented micrographs of surface features that they interpreted as TEA-
induced microcracks, Halquani et al. stated that these features were only
observed on the surfaces of HA grains that had been fractured transgranularly.

In contrast, in this study, microcracks in HA are presented for a series of
specimens with aS-sintered surfaces, to avoid the complication of attempting to
separate surface features induced by TEA-microcracks from those features
generated during the macrofracture of the specimen. Furthermore, building upon
a close examination of the HA studies by Ruys et al. [39] and Halquani et al.
[46], we have employed pertinent literature on sintering in ceramics and
microcracking in ceramics to construct arguments that the observations of RuyS
et al. [39] and Halquani et al. [46] strongly support the case for TEA-induced
microcracking in HA. The micrographs included in this study (Figs. 4-6) bolster
those arguments.

A difficulty in developing HA for biomedical use has been, in part, the fact
that its properties can vary greatly from researcher to researcher as a result of for
example, differences in sintering temperature or heat treatments [39, 46, 55].
Property changes that are associated with TEA-induced microcracking have
been studied widely in the general ceramics literature [5-8, 10]. Putting the
dizzying variation of properties into context (that is, with the observation that

TEA-induced stresses and microcracking may play a major role in the variability

95

of properties) is a significant step toward fully utilizing hydroxyapatite as a

biomedical material, specifically for use as a scaffold in bone tissue engineering.

ACKNOWLEDGMENTS

The authors acknowledge J.A. Meganck (graduate of the Michigan State
University Mechanical Engineering Department) and M. Soto (a graduate of the
Chemical Engineering and Materials Science (CHEMS) Department at MSU) for
their assistance in this study with sample preparation and XRD studies. The
authors gratefully acknowledge funding from the National Science Foundation
(DMR0074439) and the MSU Foundation for support of IO. Smith as well as for
equipment and Space in the MSU College of Engineering’s new Tissue

Engineering laboratories.

96

 

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CHAPTER 6
INDUCED MICROCRACKING AFFECTS OSTEOBLAST ATTACHMENT ON
HYDROXYAPATITE SCAFFOLDS FOR TISSUE ENGINEERING

I.O. SMITH, M.J. BAUMANN, E.D. CASE

Department of Chemical Engineering and Materials Science, College of
Engineering, Michigan State University, 2527 Engineering Building, East
Lansing, Michigan, 48824, USA

PUBLISHED IN: J Am Biochem Biotechnol 2006;22105-110.

ABSTRACT

Bone microdamage caused by routine activity plays an important role in
triggering targeted and nontargeted bone remodeling. Targeted remodeling
occurs near localized areas of microdamage”. We hypothesize that bone
remodeling may be directly and positively influenced by inducing microcracks in
hydroxyapatite (HA) scaffolds for bone tissue engineering. A study by Case et al.
Showed microcracking in HA discs having >98% theoretical densitys. These
microcracks occurred without the application of external stress, likely as a result
of thermal expansion anisotropy (TEA). TEA generates microcracks only when
the Specimen grain size exceeds a critical value (GCR). Due to conflicting data in
the literature on the thermal expansion along the crystallographic axes of HA, it is
difficult to estimate GCR precisely for HA, but GCR likely ranges from a few tenths
of microns to several micronss. As the grain size of a polycrystalline Specimen
increases above the critical grain Size, the number of microcracks also increases
but the number density of microcracks is difficult to control. Therefore, in this

study we used Vickers microindentation to induce controlled microcracks in

101

>99% dense HA discs. The effect of microcracking on the osteoblast (OB)
attachment was then quantified. Control HA discs and microcracked HA were
seeded with MC3T3-E1 OBS and cultured for 4 hours. The OBS were then
stained with Rhodamine-Phalloidin and Hoechst fluorescent dyes to identify the
actin fibers and nuclei, respectively. 08 attachment was quantified using
fluorescent light microscopy. OB attachment at 4 hours was 29% i 7.6% of the
initial seeded OBS for the microcracked HA specimens and 18% i 6.1% of initial
seeded OBS for the non-microcracked control HA Specimens. The difference in
these OB attachment values was statistically significant as determined via the

Student t-test (p = 0.004, with p < 0.05 taken to indicate significance).

102

INTRODUCTION
Microcracking in Bone

In a 1960 study of cracks located primarily in the cement line of human
rib bones, Frost7 first suggested that microcracking in bone results from
mechanical fatigue. These cracks were not exclusively an artifact of
specimen preparation, but rather an inherent characteristic of the bone. More

recently, similar cracks have been observed in humana‘m, bovineIO'“,

2

canineI'1 and equine bone13'”. Zioupos et al. report the size of these cracks

as 10 — 50 pm in length with an apparent opening of 2 — 5 pm“.

Bone exposed to routine activity undergoes repeated cyclic strain, on
the order of 0.0003 to 0.0008 and 0.0008 to 0.0012 for walking and running,
respectively” 16. While these values are lower than the maximum strain
allowed by bone in a single cycle, their repetitive nature leads to an
accumulation of microcracks". Left unchecked, this microdamage coalesces
into macrocracks that can cause failure of the bone8'18'19. The process of
microcrack accumulation to produce macrocracks iS especially pronounced in
aged individualsg. Frost also suggested a relationship between bone
microcracking and subsequent bone remodeling7. Microdamage in bone likely
triggers both targeted and nontargeted remodeling”. Nontargeted remodeling
is not site dependant and occurs less frequently”, while targeted remodeling
occurs at sites of localized microdamage as observed in studies by Johnson

| 2-4. 13

et al., Burr et al. and Mori et a . In these cases, remodeling was

significantly more focused around regions of microdamage”. Remodeling

103

occurs when osteoclasts resorb damaged bone, including osteocytes”. The
resorption process is regulated by the osteocyte syncytium in local regions of

remodeling”25

. Osteocyte damage and subsequent apoptosiS leads to
increased osteoclast activity followed by subsequent osteoblast activity.
Because microcracks occur with routine fatigue loading, the dynamic repair of

microdamage iS therefore an integral part of maintaining the integrity of the

bone?

Microcracking in Hydroxyapatite

Case et al. were the first to definitively identify microcracking in dense
(>98% theoretical density) HA which occurred in the absence of an external
applied stresss. This microcracking was likely a result of thermal expansion
anisotropy (TEA), since non-cubic crystalline materials such as HA have
dissimilar thermal expansion coefficients along the crystallographic axes. As a
polycrystalline body cools following Sintering, TEA causes neighboring grains
to shrink at differing rates and in turn generates a stress that is sufficient to
cause microcracking if the mean grain size exceeds a critical grain size, GCR.

Although TEA generated microcracks were observed in the dense HA
specimens (average grain size of 7.9 pm i 1.5 pm) examined by Case et al.5,

it is difficult to precisely estimate GQR for HA. AS discussed by Hoepfner et al.,
there are large disparities among the values reported in the literature for the
thermal expansion coefficient as a function of crystallographic direction, which

lead to large uncertainties in calculating GCR for HA6.

104

In general, the number density of microcracks increases with
increasing grain size, but in the absence of extensive measurements, it is
difficult to precisely predict changes in TEA-induced microcracks.
Furthermore, there is no control over the spatial distribution of TEA-induced
microcracks. Therefore in our HA Specimens, in order to control both the
number density and placement, microcracks were induced by Vickers micro-
indentation where the microcrack length was controlled by varying the
indentation load. The following relationships were used to predict the

microcrack length (20) and the indentation impression diagonal length (2a)

 

, E V= P
Ixc. 261(5) [6%] (1)
H =1.8544(§P—2) (2)

where Kc is fracture toughness, E is elastic modulus, H is hardness, P is

indentation load, and o is a calibration constant, using values of KC = 1
MPan, E = 115 GPa and a = 0.01626' 27.

The array of Vickers indentations resulted in an areal distribution of
microcracks. Cell counting was used to quantify OB attachment on both the
microcracked and the nonmicrocracked HA specimens. The effect of
microcracking on OB morphology was assessed by direct observation using

optical fluorescence microscopy. A spindle shaped or polygonal spread OB

105

morphology was then used as an indicator of typical OB attachment to the HA

specimen surfaces28’31.

MATERIALS AND METHODS
Powder processing

Dense HA discs were fabricated using high purity powder (Taihei
Chemical, Osaka, JP) with a vendor-specified mean particle size of 1 - 3 pm.

In a binderless process developed by McMullen, HA powder was uniaxially
pressed into discs at 35 MPa for 1 minute using a KBr pellet die (lntemational
Crystal Laboratories, Garfield, NJ) or a hydraulic press (Carver Inc, Wabash,
IN)”. The resulting HA powder discs (diameter = 32 mm, thickness = 2 — 3
mm) were sintered in air at 1360°C for 4 hours, at a heating/cooling rate of
~10°Clminute resulting in HA discs of 23 mm in diameter and 1 - 2 mm in

thickness and >99% theoretical density as determined by Archimedes method
(n = 5). The Sintered HA Specimen grain Size was 4.8 pm i 1.0 pm,
determined by the line-intercept method (n = 5).

The surfaces of the HA discs were ground flat using 30 pm (600 grit)
SiC wet abrasive paper and then all were sequentially polished using 30 pm,

12 pm, 9 pm, 6 pm and 1 pm diamond abrasive suspension or paste for 5

minutes at each grade. A Buehler EcoMet 3 variable speed grinder/polisher
and an AutoMet 2 power head were used to grind and polish. Polished

Specimens were ultrasonically cleaned by submersion in reverse osmosis

106

(RO) water followed by pulsing with an ultrasonic probe at 20 kHz and 250

Watts for 5 minutes.

Microcrack Induction
Surface-breaking microcracks were induced by indenting each
specimen 196 times using a Buehler Micromet ll microindenter (Buehler, Lake

Bluff, IL), with a load of 4.9 N, a load time of 5 seconds and a loading rate of
70 pm/sec. These loading parameters were found by the authors to be the

optimum set of conditions for inducing microcracks with minimal spalling and
chipping of the HA specimens.

lndentations were made over a single 14 mm x 14 mm grid, with a one
mm indentation spacing, across each specimen surface. The square-based
pyramidal diamond indenter tip used for Vickers indentations creates
diamond-Shaped impressions on the HA disc surface. Microcracks are
formed by the radial cracks originating from the comers of these indentation
impressions (See Figure 1, where white arrows indicate radial cracks).
Following indentation, each HA disc was submerged in R0 water for 8 hours
to insure crack growth saturation. The average indentation size and crack
size were determined by optical microscopy for a representative selection of
indentations (n = 20) on each HA Specimen. A preliminary study of the effect
of HA specimen immersion in R0 water was conducted, in which 8
indentations were applied to a single dense HA disc, which was submerged in

R0 water, to determine the effectiveness of this treatment in assisting crack

107

growth. Crack length was measured at t = 30 seconds immediately following
indentation, as well as at submersion times of 8 and 24 hours. The results of
this trial are mentioned in addition to the average indent size and post-

saturation crack length in the results section.

 

 

Figure 1: SEM micrograph of microindentation on HA Specimen surface (right). "2a" indicates
the diagonal length of the indentation impression and "2C" indicates crack length (left).

MC3T3-E1 culture

M03T3-E1 osteoblasts (OBS) of subclone 14, obtained from ATCC
(Manassas, VA) were thawed and plated. The OBS were cultured using alpha
minimum essential media (Ct-MEM, Invitrogen Grand Island, NY)
supplemented with 10% fetal bovine serum (FBS, Atlanta Biologicals,
Lawrenceville, GA), incubated in a humid atmosphere at 37°C and 5% C02.

The OBS were split regularly with cell passages 25-29 used for this study.

108

OB attachment study

Microcracked HA and control HA disc specimens were autoclave
sterilized, collared and placed into 6-well polystyrene cell culture plates in
preparation for OB seeding. Prior to seeding, OBS were removed by

trypsinization (Trypsin/EDTA, Invitrogen, Grand Island, NY), counted with a
hemacytometer, centrifuged and resuspended in fresh complete a-MEM.

Following standard procedures, OBS were seeded onto the HA discs at a
density of 11,320 cells/cm2 and enough media was added to account for
evaporation”. After 4 hours incubation at 37°C and 5% CO2 in a humid
environment, the media was removed and each HA specimen washed using
1X phosphate buffered saline (PBS), to remove any non-attached OBS. OB
seeding was done in triplicate, with n = 3 for a total of nine microcracked HA
and nine control HA specimens.

OBS were fixed using 3.7% formaldehyde (H2CO) (JT Baker,
Phillipsburg, NJ) and perrnealized with a solution of 0.1% Triton X-100
(Sigma-Aldrich, St. Louis, MO). Prior to staining, OB seeded HA Specimens
were pre-incubated in 1% bovine serum albumin (BSA, Invitrogen).

The actin cytoskeleton was stained using Rhodamine-Phalloidin while the
nucleus was stained using Hoechst 33342 stain (lnvitrogen)32. Stained 0B
seeded HA specimens were mounted in glycerol solution (JT Baker) and
viewed using a Leica DM IL fluorescence microscope (Leica Microsystems,
Wetzlar, GER) and optical micrographs recorded using the Spot RT camera

(Diagnostic Instruments, Inc., Sterling Heights, MI). Rhodamine-Phalloidin

109

 

stain fluoresces at a wavelength of Aemmision = 565 nm when exposed to light
of wavelength )Vexcitation = 542 nm while the Hoechst stain fluoresces at a
wavelength of Aemmision = 465 nm when exposed to light with wavelength of

)I-excitation = 355 nm-

OB attachment was quantified by counting nuclei over the
microscope's field of view. Ten fields of view were randomly selected and
counted for each Specimen. The student's t-test was used to determine
statistical significance, with p < 0.05 taken as significant.

The effect of induced microcracking in HA specimens on OB
morphology was assessed by staining the actin fibers and observing the OB

shape using optical fluorescence microscopy.

RESULTS AND DISCUSSION

Vickers indentation impressions in the dense HA Specimens had an

average diagonal length 2aAVE = 54.7 pm i 3.6 pm (Fig. 1). After crack

saturation, the mean crack length was 2CAVE = 176.3 pm i 23.8 pm (Fig. 1).
Immersing microcracked HA specimens in R0 water resulted in an

increase in crack length from 143 pm i 23.8 pm to 158 pm i 17.4 pm after 8

hours and to 167 pm i 14.0 pm after 24 hours (n = 8). This is consistent with

previous findings that H2O assists crack growth“. Student t-test determined
that the increase in microcrack length between 8 and 24 hours was not

Significant.

110

 

Figure 2 is a plot of OB attachment at four hours for microcracked HA
specimens (n = 9) and control HA specimens with no microcracking (n = 9),
Showing the Significant difference in OB attachment between the two. Error
bars indicate standard deviation. After 4 hours in culture, OB attachment on
the microcracked HA specimens was 29% i 7.6% of initial seeded OB
number compared to 18% i 6.1% of initial on the control HA Specimens. This
represents a 60% increase in OB attachment with microcracking. Student's t-
test produced a p-value of 0.004, presence of microcracks on the HA surface

had a positive and significant influence on OB attachment.

4O

 

I Control
I Microcracked

Percent attached

 

 

 

Figure 2: Osteoblast attachment at 4 hours on microcracked and non microcracked dense HA
discs

The role of the Vickers indentations in increased OB attachment on HA

specimens should be addressed. If we consider each indentation impression

111

as a surface breaking pore of an average size 2a = 54.7 um i 3.6 pm, then

we can compare this size to that of typical pores designed to promote OB
ingrowth. OBs respond to a bimodal system of pores, with pore sizes of 200 —

400 pm for macropores and < 70 pm for interconnecting micropores35'37.

Since the Vickers indentation impressions do not approach this scale and do
not include any interconnections, it is unlikely that they are the governing
factor in increasing OB attachment.

Figure 3 includes optical micrographs showing OB attached to the
surface of a control HA specimen (Fig. 3A) and OB attachment on a
microcracked HA specimen (Fig. SB). In addition to Hoechst stain to
illuminate the cell nucleus, Rhodamine-Phalloidin stain was used to illuminate
the actin fibers within the CBS The CBS attached to the surface of control HA
have spread morphology (Fig. 3A), indicated by the either spindle-shaped or
polygonal appearance of each cell. This spread morphology is consistent with
the morphology of M03T3-E1 085 reported by Sudo et al. and other828'31,
and infers strong adhesion of the 083 to the HA substrate”. The presence of
some spindle-shaped CBS and some polygonal CBS is consistent with
observations of OB attachment on HA substrates at 4 hours by McMullen”.
Figure BB also shows the spread spindle-shaped morphology of the 035 on
the surface of microcracked HA, indicating that the induction of microcracks

on HA specimens does not change the shape of the attached CBS.

112

 

Figure 3: 3A shows MCBT3-E1 osteoblasts with both spindle-shaped and polygonal
morphology on HA control specimen, stained to highlight actin fibers (red — Rhodamine-
Phalloidin) and nuclei (blue — Hoechst). 3B shows MC3T3-E1 Osteoblasts with spindle
shaped morphology on HA microcracked specimen, stained to highlight actin fibers (red —
Rhodamine-Phalloidin) and nuclei (blue - Hoechst).

113

CONCLUSION

Microcracked HA specimens showed 60% greater OB attachment in vitro
compared to non-microcracked control HA specimens (Figure 2). To capture
early attachment events, future studies to assess OB attachment at earlier
times will be done. Additional studies are also needed to determine whether
the
induction of microcracks positively affects OB proliferation, differentiation and
mineralization over time in vitro. However, the initial findings presented in this
study indicate that microcracks enhance OB attachment activity. By
mimicking the naturally occurring damage process that triggers bone
remodeling, microcracks induced by Vickers indentation positively and

significantly influenced the initial OB response to HA scaffolds.

114

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118

CHAPTER 7

Early Stage MC3T3-E1 osteoblast attachment on microcracked
hydroxyapatite scaffolds for tissue engineering

IAN 0. SMITH, MELISSA J. BAUMANN, ELDON D. CASE
Department of Chemical Engineering and Materials Science, The College of
Engineering, Michigan State University, East Lansing, MI 48824

SUBMITTED TO: J Biomed Mater Res, Dec 2006.

ABSTRACT

The effect of Vickers indentor—induced microcracking on osteoblast
(OB) attachment on dense hydroxyapatite (HA) scaffolds for tissue
engineering was assessed at four hours by the current authors.1 This study
extends our earlier work to early stage (one hour) OB attachment on such
scaffolds. A population of > 3500 microcracks was induced across nine dense
HA discs and omitted in nine control HA discs. MC3T3-E1 murine 085 were
seeded onto these HA specimens and the OB attachment was then quantified
and OB morphology assessed at one hour using fluorescent staining and
fluorescent optical microscopy. The results were compared to our earlier four-
hour OB attachment data.1 Early stage (one-hour) OB attachment on HA was
not significantly affected by the presence of microcracks, which resulted in
attachment of 31.3% i 12.7% versus 25.1% i 8.9% on control HA surfaces.
This is in contrast to the 61.1% increase at 4 hours increasing from 18% 3:
6.1% to 29% :t 7.6% for microcracked versus control, shown in our earlier

work.1 Therefore, while microcracking does have a statistically meaningful

119

 

effect at four hours attachment time, there is no significant effect on the early

stage (1 hour) OB attachment on HA.

INTRODUCTION
The work presented in this paper differs from our earlier work‘, where
osteoblast (OB) attachment was assessed at four hours, because we are now
observing early stage (one hour) OB attachment. A triplicate model was
utilized with three microcracked HA specimens and three non-microcracked
control HA specimens in each trial, for a total of nine microcracked and nine
control specimens. MC3T3-E1 murine 083 were seeded onto these HA
specimens and the subsequently attached OBs were then counted at one
hour and their morphology assessed using fluorescent staining and
fluorescent optical microscopy. The results were compared to those
previously collected at four hours1 in order to first assess the effect of
microcracking in the early phase of OB attachment and then to evaluate the
overall potential of microcracking on OB activity.
If Bone microdamage occurs in the form of microcracks in humans“,
thoroughbred racing horses5'6 and greyhounds7'8 following repetitive physical

activity. For example, while the strain levels in human bone during walking
reach a = 3 x 10'4 to 8 x 10“, and during running reach a = 8 x 10“ to 1.2 x 10'

3, this is less than the threshold failure strain in a single cycle. However, the
repetitive nature of these activities leads to an accumulation of

microdamage.9'1O Unless such activity ceases or declines dramatically, this

120

microdamage can coalesce into macrocracks, leading to catastrophic
fracture?“12 The connection between bone microcracks and repetitive
physical activity was first postulated by Frost in 1960, who used Fuchsin
staining to detect their presence.13 More recently, microcracks have been
identified in bone using increasingly sophisticated techniques including optical
and confocal laser scanning microscopy (CLSM) using Fuchsin3'4'3'6
haematoxylin and eosin7 and fluorescent staining4'”, scanning electron
microscopy (SEM)5 and radiographyz.

In most cases, if the physical activity that causes the damage is
reduced, catastrophic failure is prevented because bone has the ability to
dynamically remodel following the onset of microdamage. Osteoclasts (OCs)
are specialized bone cells, which resorb bone while 083 generate new bone
in the vicinity of the microcracks.15 OC and OB activities are increased when
the syncytium, adjacent to areas of bone microdamage, signals for
remodeling to begin through increased osteocyte apoptosis. This coupling
between remodeling and microcracking may paradoxically result in a lower
incidence of initial catastrophic fracture because of the ability of the
microcracks to dissipate energy through crack formation, which in turn signals
bone healing.

While the links between microcracking, osteocyte damage, apoptosis
and subsequent in vivo bone remodeling in bone are known‘azo, the

relationships between the degree of microcrack damage in hydroxyapatite

(HA) and OB attachment is not yet fully understood. HA is a calcium

121

phosphate-based ceramic with a crystal structure and composition close to
that of the mineral component of bone which which leads to HA’s use as a
bone replacement scaffold material.”22 We have previously shown that OB
attachment at 4 hours increased significantly in microcracked dense HA
samples compared to dense HA control specimens.1 It is noteworthy that this
increase occurred in the absence of the osteocyte syncytium instrumental to
the in vivo bone remodeling process. These results indicate that osteocyte
signaling is likely not the sole mechanism for initiating bone remodeling.
Therefore microcracked HA scaffolds may offer promise as a synthetic bone
scaffold material for enhanced bone tissue repair.

In our earlier study, microcracks were created by placing 196 Vickers
indentations in the HA specimens. The indentations were arranged in a
14x14 grid with the indentation centers located 1 mm apart.1 Each indentation
produced two perpendicular microcracks for 392 microcracks per specimen
on a total of nine specimens.

96 Vickers indentation is a method used to assess hardness in brittle
materials. Vickers indentation uses a square-based pyramidal diamond
indenter tip, which, upon loading, leaves a diamond-shaped impression with
radial microcracks extending into the material from the comers of that
impression. Figure 1 is a SEM micrograph of a single Vickers indentation on a
HA surface, showing the induced radial cracks (labeled as “R"). The
microcrack length (2c) and indent width (23) are indicated on this micrograph

using white arrows.

122

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, .

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.56"

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I
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Figure 1: SEM micrograph of microindentation on HA specimen surface. “R" indicates radial
cracks, the distance "2a" indicates the diagonal length of the indentation impression and the
distance "20" indicates crack length.

The grid of Vickers indentations created a uniform microcrack density
across the HA scaffold surface. In contrast, thermal expansion anisotropy
(TEA) induced microcracks, such as those identified in our earlier work, are
more difficult to control.23 TEA-induced microcracks form during cooling in
non-cubic crystalline materials when the average grain size is above a critical

value (G > GCR). This value is difficult however to estimate for HA because of
variations in reported values of the thermal expansion coefficient, on.24 While

we were the first to definitively identify microcracking without an externally
applied stress in dense HA23, it was not until our more recent work that we

were able to control the size and number of these microcracks.1

123

MATERIALS AND METHODS

HA disc specimens were fabricated using medical-grade purity HA
powder with a manufacturer-reported particle diameter of 1 — 3 um (Taihei

Chemical, Osaka, Japan), which was die-pressed into 32 mm disc-shaped
green compacts. These green specimens were sintered in air at 1360°C for 4
hours, at a heating/cooling rate of ~10°Clmin. Specimens were then
subjected to further heat treatment, not related to the current study, at 1100°C
for 45 minutes at a heating/cooling rate of ~10°Clmin.2 Using this process, we
fabricated specimens of > 99% theoretical density (determined by
Archimedes method) that are ~23 mm in diameter and 1 — 2 mm thick, with an
average grain diameter of 4 um i 1 pm. A total of 18 specimens were used in
this study, nine microcracked and nine controls.

Immediately prior to their use, these HA discs were ultrasonically

cleaned, ground using 600 grit (30 um) silicon carbide abrasive paper and

sequentially polished using 30 um, 12 pm, 9 um, 6 pm and 1 pm diamond

abrasive suspensions for 5 minutes each. Grinding and polishing was
completed using a Buehler EcoMet 3 polishing wheel with an AutoMet 2
polishing head (Buehler, Lake Bluff, IL). Following the final polishing step,
specimens were cleaned by pulsing with an ultrasonic probe while submerged

in reverse osmosis distilled (RO) H20.

 

' This second heat treatment was necessary to induce grain-boundary grooving vital to an
earlier, yet unpublished study. This grooving was eliminated during the grinding process, prior
to OB seeding in the current study.

124

Microcracks were induced, following methods established in our
previous study‘, across the surface of half of the specimens using a Buehler
MicroMet ll Vickers indentor. The placement of Vickers indentations using this
established load magnitude, loading time and loading speed resulted in
cracks with the least amount of spalling and chipping, which are
counterproductive to microcrack formation. The indentations were arrayed in
a 14x14 grid, 1 mm apart, with two perpendicular radial microcracks emerging
from each indentation pyramid for a total of 392 cracks/specimen and > 3500
cracks total across all nine specimens. Following indentation, each HA disc
was submerged in R0 distilled HzO for eight hours in order to facilitate crack
growth saturation prior to cell seeding.1

For this study, MC3T3-E1 clonal murine 083 of subclone 14 (American
Type Culture Collection, Manassas, VA) were thawed from frozen stock (-

180°C) and plated onto polystyrene tissue culture dishes. OB cultures were

fed using complete media; a combination of 90% a-minimum essential media

(a-MEM, Invitrogen, Grand Island, NY) and 10% fetal bovine serum (FBS,

Atlanta Biologicals, Lawrenceville, GA). OB cultures were then placed into an
incubator at 37°C and 5% C02 in a humid atmosphere and split every two
days. 083 were split until a sufficient cell number was reached for this study
(passages 28-29) first by trypsinizing the 083 (trypsin/EDTA, Invitrogen,
Grand Island, NY) and then dividing them to be replated at a lower OB

density, to allow room for the 083 to divide.

125

The control (nine discs) and microcracked (nine discs) HA specimens
were autoclave sterilized at 125°C for one hour prior to OB seeding. The OBs
were seeded onto sten‘le HA substrates at a cell density of 11,320/cm2 and 1
ml complete media added to each well.25 Prior to seeding, each HA substrate
was surrounded by a commercially pure Al foil collar to control cell density by
preventing the 085 from falling off the specimen surface. These seeded
specimens were then incubated at 37°C and 5% C02 in a humid atmosphere
for one hour.

After incubation, the OB-seeded HA discs were removed from the
incubator and rinsed twice with 1x phosphate buffered saline (PBS) to remove
any non-attached cells. The attached 083 were then fixed and stained using
Rhodamine-Phalloidin, which acts to illuminate the cell cytoskeleton, and
Hoechst, which allows the cell nucleus to be imaged.26

Attached 083 were imaged using an optical fluorescence microscope
(Leica DM IL, Leica Microsystems, WetZlar, GER) and a Spot RT camera
(Diagnostic instruments, lnc, Sterling Heights, MI). Rhodamine-Phalloidin
fluoresces at A = 565 nm and specifically stains the actin fibers within the OB
cytoskeleton. This provides a clear image of the morphology of each OB.
Hoechst stains the cell nuclei and fluoresces at A = 465 nm, providing an
image of each OB nucleus. Ten randomly selected fields of view were imaged
per specimen surface and the number of OB nuclei counted in each.
Statistical significance in cell counts was determined using Student’s t-test,

with p < 0.05 taken to indicate a significant difference. The effect of

126

microcracking on OB morphology was assessed after staining for actin fibers
by observing the shape of the cytoskeleton, imaged using optical

fluorescence.

RESULTS AND DISCUSSION
Vickers indentations were produced as described earlier and the
indentation diagonal lengths measured.1 The indenter-induced impressions

on the surface of dense HA had an average diagonal length of 2am; = 47.3

um i 4.6 pm and yielded an average post-saturation crack length of 2cAVE =

124.4 pm i 21.1 um.

Figure 2a is a plot of OB attachment at one hour for both control HA (n
= 9) and microcracked HA (n = 9) specimens. For comparative purposes,
Figure 2b includes four-hour time point attachment data reported in our earlier
study.1 Error bars represent the standard deviation and “#” indicates a

statistically significant difference.

127

Percent Attached

Percent attached

50.0 -

 

 

 

 

 

 

 

 

 

 

 

 

D Control
I Microcracked
40.0 —
30.0 — I
20.0 ~ I
10.0 -
0.0
1 hr
Time
50.0 - #
El Control
40'0 l Microcracked
30.0 -
20.0 — [
10.0 I
0.0
4 hr
Time

Figure 2: A) MC3T3-E1 OB attachment at 1 hour on control and microcracked HA specimens.
B) MC3T3-E1 OB attachment at 4 hours on control and microcracked HA specimens. 4-hour
attachment data was taken with permission from previous work by Smith et al.1 # indicates
that difference in CB attachment between control HA and microcracked HA at 4 hours is
significant (p = 0.004).

128

OB attachment at one hour was 25.1% i 8.9% of 083 plated on the
control HA and 31.3% i 12.7% of OBs plated on the microcracked HA.
Student's t-test gave a p-value of 0.25, meaning that there was no significant
difference between the OB attachment on microcracked and non-
microcracked HA specimens. Therefore the presence of microcracks on HA
surfaces does not have a statistically meaningful effect on early stage 1-hour
OB attachment. In contrast, after four hours of culture, the percentage of
plated 085 which attached on microcracked HA was 61.1% greater, when
compared to non-microcracked HA, increasing from 18% i 6.1% to 29% i
7.6% of 033 plated.1 Student's t-test showed that this increase in the
percentage of OB attachment was significant (p = 0.004). Since data for each
time point is taken from a different study, using a different batch of the same
MC3T3-E1 OB subclone, a direct quantitative comparison between the one
and four-hour attachment data cannot be made here.

The images in Figure 3 are optical fluorescence micrographs showing
033 attached to the surface of control and microcracked HA scaffolds.
Images were enhanced by staining with Rhodamine—Phalloidin, which
illuminated the actin fibers within the OB cytoskeleton, and Hoechst stain to
highlight the OB nuclei. Figure 3A is a photomicrograph of the control HA
surface, showing six attached OBs. These cells exhibit either spread or
polygonal morphology, which is consistent with the MC3T3-E1 OB
morphology reported by Sudo et al.27'28 For OBs, both the spread and

polygonal morphology is indicative of strong adhesion to the underlying

129

substrate.” Figure SB is a photomicrograph of the microcracked HA surface
having seven attached 083, all of which also exhibit either spread or
polygonal morphology, indicating that microcracks do not change the
attached OB morphology. This OB morphology is consistent with that in our
earlier findings, showing both the spread and polygonal OB morphology was
present at four hours attachment time, indicating that microcracks do not

change OB morphology at four hours attachment time either.1

130

 

Figure 3: A) MCST3-E1 083 with spread (white circle) and polygonal (black circle)
morphology on HA control specimen. B) MCST3-E1 035 with similar spread (white circle)
and polygonal (black circle) morphology on microcracked HA specimen. 085 have been
stained to highlight actin fibers (red — Rhodamine-Phalloidin) and cell nuclei (blue — Hoechst).

CONCLUSIONS

We have assessed the effect of induced microcracking on early, 1-hour
OB. Microcracks were induced on the surface of nine HA discs by indenting
each disc 196 times using a Vickers microindenter. This resulted in 392 radial
cracks per specimen on nine specimens for a total of 3528 induced
microcracks. This large population of microcracks is in contrast to the nine

non-indented, non-microcracked control HA disc surfaces. While the grain

diameter of the specimens used in this study (4 pm i 1 um) is larger than GCR

for TEA-induced microcracking as estimated for HA (1.1 pm)“, surface

breaking microcracks were not observed in these specimens under SEM.

M03T3-E1 OB attachment data shows that this large population of
microcracks does not significantly increase OB attachment at one-hour
attachment times in comparison to non-microcracked samples. This result is
in contrast to the 61.1% improvement in mean OB attachment at four hours
reported previously.1

Additionally, the presence of microcracks did not affect OB morphology
on the HA surface. Both microcracked and control HA surfaces led to either
spread or polygonal attached osteoblasts, pointing to strong attachment to the
underlying substrate.29 This finding is in agreement with our four-hour
attachment data, which found a no significant link between microcracking and
resulting OB morphology.

While enhanced OB attachment at a time of four hours may preview

improved cell differentiation and production of mineral matrix by mimicking the

132

naturally occurring damage mechanism that brings about bone repair, one
hour is too early for this phenomenon to be observed.
Our future work will focus on assessing the effect of induced

microcracking on OB growth and differentiation.

133

REFERENCES

1. Smith IO, Baumann MJ, Case ED. Induced microcracking affects
osteoblast attachment on hydroxyapatite scaffolds for tissue engineering. Am
J Biochem Biotechnol 2006; 2: 105-110.

2. Blickenstaff LD, Morris JM. Fatigue fracture of the femoral neck. J Bone
Joint Surg 1966; 48A: 1031-1047.

3. Schaffler MB, Choi K, Milgrom C. Aging and matrix microdamage
accumulation in human compact bone. Bone 1995; 17: 521-525.

4. Lee TC, O'Brien FJ, Taylor D. The nature of fatigue damage in bone. Int J
Fatigue 2000; 22: 847-853.

5. Stepnik MW, Radtke CL, Scollay MC, Oshel PE, Albrecht RM, Santschi
EM, Markel MD, Muir P. Scanning electron microscopic examination of third
metacarpal/third metatarsal bone failure surfaces in thoroughbred racehorses
with condylar fracture. Vet Surg 2004; 33: 2-10.

6. Da Costa Gomez TM, Radtke CL, Kalscheur VL, Swain CA, Scollay MC,
Edwards RB, Santschi EM, Markel MD, Muir P. Effect of focused and radial
extracorporeal shock wave therapy on equine bone microdamage. Vet Surg
2004; 33: 49-55.

7. Johnson KA, Skinner GA, Muir P. Site-specific adaptive remodeling of
Greyhound metacarpal cortical bone subjected to asymmetrical cyclic loading.
Am J Vet Res 2001; 62: 787-793.

8. Muir P, Johnson KA, Ruaux-Mason CP. In vivo matrix microdamage in a
naturally occurring canine fatigue fracture. Bone 1999; 25: 571-576.

9. Lanyon LE, Hampson WGJ, Goodship AE, Shah JS. Bone deformation
recorded in vivo from strain gauges attached to the human tibial shaft. Acta
Orthop Scand 1975; 46: 256-268.

10. Burr DB, Milgrom C, Fyhrie DP, Fon~ood M, Nyska M, Finestone A,
Hoshaw S, Saiag E, Simkin A. In vivo measurement of human tibial strains
during vigorous activity. Bone 1996; 18: 405-410.

11. Schaffler MB, Radin EL, Burr DB. Mechanical and morphological effects
of strain rate on fatigue of compact bone. Bone 1989; 10: 207-214.

12. Schaffler MB, Radln EL, Burr DB. Long-term fatigue behavior of compact
bone at low strain magnitude and rate. Bone 1990; 11: 321-326.

134

13. Frost HL. Presence of microscopic cracks in vivo in bone. Henry Ford
Hospital Med Bull 1960; 8: 27-35.

14. Zioupos P, Currey JD. The extent of microcracking and the morphology of
microcracks in damaged bone. J Mat Sci 1994; 29: 978-986.

15. Burr DB. Targeted and nontargeted remodeling. Bone 2002; 30: 2-4.

16. Burr DB, Martin RB, Schaffler MB, Radin EL. Bone remodeling in
response to in vivo fatigue microdamage. J Biomech 1985; 18: 189-200.

17. Bronckers ALJJ, Goei W, Luo G, Karsenty G, D'Souza RN, Lyaruu DM,
Burger EH. DNA fragmentation during bone formation in neonatal rodents

assessed by transferase-mediated end labelling. J Bone Miner Res 1996; 11:
1281-1291.

18. Colopy SA, Benz-Dean J, Barrett JG, Sample SJ, Lu Y, Danova NA,
Kalscheur VL, Vanderby Jr. R, Markel MD, Muir P. Response of osteocyte
syncytium adjacent to and distant from linear microcracks during adaptation
to cyclic fatigue loading. Bone 2004; 35: 881-891.

19. Noble BS, Peet N, Stevens HY, Brabbs A, Mosely JR, Reilly GC, Reeve
J, Skerry TM, Lanyon LE. Mechanical loading: biphasic osteocyte survival
and targeting of osteoclasts for bone destruction in rat cortical bone. Am J
Physiol Cell Physiol 2002; 284: 934-943.

20. Verborgt 0, Gibson GJ, Schaffler MB. Loss of osteocyte integrity in
association with microdamage and bone remodeling after fatigue in vivo. J
Bone Miner Res 2000; 15: 60-67.

21. Hench LL. Bioceramics: From Concept to Clinic. J Am Ceram Soc 1991;
74: 1487-1510.

22. DeGroot K. Effect of porosity and physicochemical properties on the
stability, resorption, and strength of calcium phosphate ceramics. In:
Bioceramics: Material characteristics versus in vivo behavior, Vol. 523. Ed.
Duscheyne , lemons J. Annals of the New York Academy of Sciences, New
York, 1988.

23. Case ED, Smith IO, Baumann MJ. Microcracking and porosity in calcium
phosphates and the implications for bone tissue engineering. Mat Sci and
Eng A 2005; 390: 246-254.

24. Hoepfner TP, Case ED. An estimate of the critical grain size for
microcracks induced hydroxyapatite by thermal expansion anisotropy. Mater
Lett 2004; 58: 489-492.

135

25. Shu R, McMullen R, Baumann MJ, McCabe LR. Hydroxyapatite
accelerates differentiation and suppresses growth of MC3T3-E1 osteoblasts.
J Biomed Mater Res 2003; 67A: 1196-1204.

26. McMullen R. An in vitro investigation of MC3T3-E1 osteoblast proliferation
and differentiation on hydroxyapatite based tissue engineered scaffolds.
Masters Degree Thesis, Michigan State University 2004: p. 86.

27. Sudo H, Kodama HA, Amagai Y, Yamamoto S, Kasai S. In vitro
differentiation and calcification in a new clonal osteogenic cell line derived
from newborn mouse calvaria. J Cell Biol 1983; 96: 191-198.

28. ltakura Y, Kosugi A, Sudo H, Yamamoto S, Kumegawa M. Development
of a new system for evaluating the biocompatibility of implant materials using
an osteogenic cell line (MCBT3-E1). J Biomed Mater Res 1988; 22: 613-622.

29. Locci P, Becchetti E, Pugliese M, Rossi L, Belcastro S, Calvitti M,

Pietrarelli G, Staffolani N. Phenotype expression of human bone cells cultured
on implant substrates. Cell Biochem Funct 1997; 15: 163-170.

136

CHAPTER 8

THE EFFECT OF GRAIN BOUNDARY GROOVING ON MC3T3-E1
OSTEOBLAST ATTACHMENT ON DENSE HYDROXYAPATITE

INTRODUCTION

Grain boundary grooving is an artifact of thermal etching, a
phenomenon that occurs in polycrystalline materials that have been heated to
elevated temperatures (T ~ 1/2 - 2/3TM). Thermal etching is the result of the
system moving towards equilibrium, where the grain boundary and surface

energy are related as shown below in Equation 1
1’35 = 273vCOS(\lI/2) (1)

where yss = the energy of the solid-solid interface between grains, st = the

solid-vapor interfacial energy and w = the angle of thermal etching.1

While the surface roughness associated with nano-scale ceramic
substrates has been associated with improved osteoblast (OB) attachment“,
the contribution of grain boundaries to this relationship is not fully understood.
We investigated the relationship between grain boundaries and OB
attachment, by determining whether the presence of grain boundary grooves,
Independent of a change in grain size, significantly affects OB attachment and

morphology on dense hydroxyapatite (HA) substrates.

137

MATERIALS AND METHODS
Specimen Fabrication

We used medical grade hydroxyapatite (HA) powder (Taihei Chemical,
Osaka, JP) with a manufacturer’s reported mean particle diameter of 1 - 3 pm

to fabricate dense HA discs. Eighteen specimens were fabricated using a
binderless powder pressing technique developed earlier.5 For each HA disc,
powder was poured into a 32 mm diameter KBr pellet die (lntemational
Crystal Laboratiries, Garfield, NJ) and uniaxially pressed at 35 MPa for one
minute, yielding a green powder compact 2 — 3 mm in thickness. As-pressed
discs were then sintered in air at 1360°C (1633 K) for 4 hours, at a
heating/cooling rate of ~10°C per minute. Sintered specimen density was
determined to be > 99%, determined by Archimedes method and the average
grain diameter was determined to be 4 pm i 1 pm using the line-intercept
method. X-ray diffraction verified that the sintered specimens were in fact HA
(Appendix B).

Sintered specimens were wet ground using 600 grit (30 um) SIC
abrasive paper and sequentially polished using 30 um, 12 pm, 9 pm, 6 pm

and 1 pm diamond abrasive paste and/or polishing suspension for 5 minutes

at each level. Grinding and polishing were done at 200 rpm using a Buehler

EcoMet 3 polishing wheel equipped with an AutoMet 2 automated head

138

(Buehler, Lake Bluff, IL). Specimens were thoroughly rinsed with reverse
osmosis (RO) water between steps and cleaned by pulsing (20 kHz, 250
Watts) with an ultrasonic probe in R0 water for five minutes following the final
step.

Half (9) of the polished discs were thermally etched at 1100°C (1373
K) for 45 minutes, with a heating/cooling rate of 10°/min in order to induce
grain boundary grooving. All specimens were then stored in a vacuum
dessicator until cell culture.

This processing protocol yielded two groups of equally dense HA

discs, both with an average grain diameter of 4 pm 1r 1 um, in which one

group is polished to 1 pm (control) and the other includes grain boundary

grooves. Figure 1 is a scanning electron micrograph showing the surface of a

grain boundary grooved specimen.

139

 

Figure 1: SEM micrograph of the surface of a grain boundary grooved HA specimen.

Cell Culture
MCBT3-E1 osteoblasts (033) of subclone 14 obtained from ATCC

(Manassas, VA) were used for this study. 085 were thawed from frozen stock

(-180°C/ 93 K) and plated onto polystyrene tissue culture dishes. a-minimum

essential media (a-MEM, Invitrogen, Grand island, NY) supplemented with
10% fetal bovine serum (FBS, Atlanta Biologicals, Lawrenceville, GA) was
added to the thawed 083 (complete media). 083 were then incubated at
37°C (310 K) and 5% 002 in a humid atmosphere and split approximately
every two days. Cells were split by trypsinization (trypsin/EDTA, Invitrogen,

Grand Island, NY) and re-plated onto polystyrene plates until the OB

140

population reached the level needed for the number of samples in this study
(passages 24-26).

The control and grain boundary grooved HA discs wereremoved from
the dessicator and then autoclave sterilized at 125°C (398 K) for 1 hour
immediately prior to cell seeding. 083 were plated onto these discs at a
density of 11,320 cells/cmz, where cell confluency has been shown to occur,
after which 1 ml of complete media was then added to each well.6
Commercially pure AI collars were used during cell seeding to prevent the
media from leaving the specimen surfaces, giving a constant surface area for
OB attachment. OB-seeded discs were then placed into commercially pure AI
collars and incubated in a humid atmosphere of 5% C02 at 37°C (310 K) for

four hours.

Cell imaging

Following incubation, the media was removed and each OB-seeded
HA specimen was rinsed twice with 1X phosphate buffered saline (PBS) to
remove any non-adherent 03s. The adherent cells were then fixed using
3.7% formaldehyde (H2CO) (JT Baker, Phillipsburg, NJ), permeabilized for 5
minutes using 0.1% Triton X-100 (Sigma-Aldrich, St. Louis, MO) and pre-
incubated in 1% bovine serum albumin (BSA, Invitrogen, Grand Island, NY).
The 083 were then stained using two different dyes; Rhodamine-Phalloidin
was used to stain the actin fibers of the OB cytoskeleton and Hoescht was

used to stain the nuclei (Invitrogen, Grand Island, NY). Rhodamine-Phalloidin

141

stain fluoresces at a wavelength of Xemmision = 565 nm when exposed to light
of wavelength )‘vexcitation = 542 nm, while the Hoechst stain fluoresces at a
wavelength of lemmision = 465 nm when exposed to light with wavelength of

Xexcnauon = 355 nm. This characteristic absorption/emission behavior is suited

for use with our imaging equipment.

Stained specimens were mounted in glycerol solution (JT Baker) and
the cells imaged using a Leica DM lL optical fluorescence microscope (Leica
Microsystems, Wetzlar, GER) equipped with a Spot RT camera (Diagnostic
Instruments, Inc., Sterling Heights, MI). Twenty systematic fields of view,
selected In a pattern covering the entire specimen surface, were imaged and
the cell nuclei were counted and the counts tallied in each. Students t—test
was used to determine the statistical significance in the cell counts, with p <
0.05 taken as significant.

The effect of thermal etching of HA specimens on OB morphology was
assessed by imaging the actin fibers to assess the OB shape using optical
fluorescence microscopy.

Following OB imaging, the HA discs were cleaned, repolished and half
(9) of the discs were thermally etched once again. The entire procedure was

then repeated for an attachment time of one hour.

142

RESULTS AND DISCUSSION

70 _ IControl

El Grain Boundary Grooved

8
——I

 

% 08 Attachment
8

20-

 

 

 

 

 

90 - I Control
El Grain Boundary Grooved

% 08 Attachment
(A) 8 0| 0) V (D
o o o o o

N
O

 

_I
O
I

 

 

O
I

 

Figure 2: MC3T3-E1 OB attachment at 1 (Figure 2A) and 4 (Figure 28) hours on thermally
etched and control HA specimens.

Figure 2 shows the percent of MCBT3-E1 OB attachment as a fraction

of the total number of 083 plated as a function of time for both control HA and

143

thermally etched HA discs. Since the one-hour and four-hour data sets
represent separate experiments using different OB populations, the data
cannot be directly compared. However, attachment data for both the control
and grain boundary grooved specimens can be compared at each time point.
The mean OB attachment on HA surfaces with grain-boundary grooving is
45.6% :t 12.2% at one hour which is not significantly different from the mean
OB attachment on the polished control HA substrate with no grain-boundary
grooving at one hour (47.3% i 12.8%). Similarly, at four hours no significant
change in the mean OB attachment exists between the HA substrates with
grain boundary grooving (59.0% i 13.0%) and the polished control HA
substrates (59.3% i 23.9%) with no grain-boundary grooving.

These results show that the presence of grain boundary grooving on
dense HA substrates does not significantly affect OB attachment. Earlier
studies by Webster et al. show an increase in OB attachment as a function of
increased surface roughness associated with substrates processed from
nano-scale ceramic powders. Webster et al. found that decreasing the grain
size from 177 nm to 23 nm, surface roughness increased from 17 nm to 20
nm (AI203).4 Our findings indicate that the grain boundary grooving does not
result in as pronounced a surface roughness (R8 = 10.23 i 1.8 nm) and are
not a factor in this increased OB attachment. Rather, decreasing grain size in
unpolished specimens increases surface roughness by creating a larger
number of surface asperities. This phenomenon is apparent in AFM images

published by Webster et al.‘

144

200 um

 

Figure 3: This set of four optical micrographs shows attached 085 at one hour on the surface
of a control specimen (Figure 3A) and a grain boundary grooved specimen (Figure 3B), and
also at four hours on the surface of a control specimen (Figure 3C) and a grain boundary
grooved specimen (Figure 30).

These results show that grain boundary grooving also has no effect on

the morphology of attached 083. The attached 033 shown in Figure 3A and

33 exhibit a similar morphology (average spicule length of 13.6 urn :t 10 pm
for control and 10.4 um i 5.4 urn for grain boundary grooved), as do the 08s
shown in Figure BC and 3D (average spicule length of 24.8 pm 1 10.3 pm for
control and 28.8 pm 1: 9.4 pm for grain boundary grooved). The lower degree

of OB spreading (Average spicule length of 12 um :t 8 um versus 18 um :I: 8

145

pm) on the one-hour as compared to the four-hour attachment

photomicrographs is likely due to the earlier stage of OB attachment. At four
hours attachment, 088 on the control specimen (Figure 3C) and the grain
boundary grooved specimen (Figure 3D) both exhibit a spindle-shaped
spread morphology. This is consistent with that of two of our earlier studies,
where the microcracks on the HA substrate surface did not change the OB
morphology.

While the mechanism by which microcracks form in HA is
fundamentally different than that by which grain boundary grooves are
formed, these surface features are of the same order of magnitude in size. By

looking at the micrograph in Figure 1, the grain boundary groove width can be

measured to be approximately 0.1 pm. From our earlier work7 induced

microcracks in HA are approximately 0.5 pm in width. These dimensions
should be verified statistically and using a higher magnification, but from
these estimates it is apparent that surface features between 0.1 and 0.5 um

in width do not have an effect on OB morphology. In these previous studies,
as well as the present one, the 083 have a spread polygonal and/or spindle-
shaped morphology. which is consistent with that shown for healthy 083 by

Sudo et al.8

146

REFERENCES

1.

Kingery WD, Bowen HK, Uhlmann DR. Introduction to Ceramics. New
York: John Wiley and Sons; 1960. 1032p.

. Webster TJ, Ergun C, Doremus RH, Siegel RW, Bizios R.

Nanocrystalline hydroxyapatite enhances osteoblast function. In:
Proceedings of the First Joint BMES/EMBS Conference, Atlanta, GA;
1999.p744.

Webster TJ, Siegel RW, Bizios R. Design and evaluation of nanophase
alumina for orthopaedic/dental applications. Nanostruct Mater
1999;12:983-986.

Webster TJ, Siegel RW, Bizios R. Osteoblast adhesion on nanophase
ceramics. Biomaterials 1999;20:1221-1227.

McMullen, R. An in vitro investigation of MC3T3-E1 osteoblast
proliferation and differentiation on hydroxyapatite based tissue-
engineered scaffolds. Masters Degree Thesis. Department of Chemical
Engineering and Materials Science. Michigan State University, East
Lansing, MI; 2004. 181p.

Shu R, McMullen R, Baumann MJ, McCabe LR. Hydroxyapatite
accelerates differentiation and suppresses growth of MC3T3-E1
osteoblasts. J Biomed Mater Res. 2003;67A:1196-1204.

Smith, I.O., McCabe, L.R., Baumann, M.J., MC3T3-E1 Osteoblast
attachment and proliferation on porous hydroxyapatite scaffolds with
nanophase powder, Int J Nanomed, 1[2], 189-194 (2006).

Sudo H, Kodama HA, Amagai Y, Yamamoto S, Kasai S. In vitro

differentiation and calcification in a new clonal osteogenic cell line
derived from newborn mouse calvaria. J Cell Biol 1983; 96: 191-198.

147

CHAPTER 9

SUMMARY AND CONCLUSIONS

Summary of Work Completed for this Dissertation

This research quantified the effect of two phenomena associated with
grain boundaries in hydroxyapatite (HA), microcracks and grain boundary
grooving, on osteoblast (OB) attachment. These findings tie into the larger
question of why a reduction in particle size in polycrystalline ceramics (HA in
this case) leads to an increase in OB attachment in vitro.

While Webster et al. have studied the link between decreased grain
size and improved OB behavior on dense scaffolds, an ideal bone tissue
engineering scaffold should be a more biomimetic surface where
interconnected macroporosity is of primary importance in bone tissue
ingrowth. Therefore, we investigated whether grain size affects OB
development in foamed porous HA scaffolds with bimodal porosity (Figure 1).
Bimodal porosity encompasses both macropores and micropores that link the
macropores in order to provide nutrient transmission (Figure 2), in a manner
similar to that found in cancellous bone. In Chapters 3 and 4, porous scaffolds
with an average porosity of 59% i 2.2% and 72% i 8.3% were fabricated with
micron-HA powder (having an average manufacturer-reported particle
diameter of 10 pm and an average sintered grain diameter of 8.6 um i 1.9
pm) and nano-HA powder (having an average manufacturer-reported particle
diameter of 20 nm and an average sintered grain diameter of 588 nm i 55

nm), respectively.

148

 

Figure 1: Comparison of SEM micrographs of the surface of bimodally porous scaffolds
fabricated using micro HA (left) and nano HA (right). Scaffolds were fabricated by foaming a
suspension of HA powder using H202 as a foaming agent, drying at 125°C for 1 hour and
sintering either at 1360°C for 4 hours (micro HA) or 1100°C for 1 hour (nano HA). Image
taken from Smith et al.. Int J Nanomed 2006;1:189-194.

 

Figure 2: SEM micrograph showing the surface of porous scaffold fabricated using nano HA,
including macropores (large arrows noting approximate boundary of a macropore) and
micropores (small arrows emphasizing the <60 pm interconnecting pores). Image taken from
Smith et al., Int J Nanomed 2006;1:189-194.

149

Using a hemacytometer, MC3T3-E1 murine OB attachment was
quantified at 0.5, 1, 2 and 4 hours and OB proliferation (growth) was
quantified at 1, 3 and 5 days. This original data is shown in Figures 3 and 4.
Background counts were also taken to eliminate the contribution of OB-sized
HA particles and to increase the OB count accuracy. Adjusted data from
background counts is shown in Figure 5. We found that OB attachment is not
significantly different between 0.5 and 4 hours and that OB proliferation was
not significantly different from 1 to 5 days. This suggests that surface
topography within the walls of these bimodally porous scaffolds (on the scale
of the grain size) is not a significant contributor to OB attachment and
proliferation and that bimodal porosity, is a greater factor in promoting OB
attachment. However, high hemacytometer cell counts indicate that further
inaccuracies exist in these results and a more accurate method must be

employed to measure OB attachment.

150

 

600

500 4

400 - a

I

300 ’ .

°/o Cells Plated

200 ~ ’

 

100 «

 

 

 

.1
—4
—I

0 7 I T I T
O 0.5 1 1.5 2 2.5 3 3.5 4 4.5

Time (Hours)

Figure 3: MC3T3-E1 OB attachment on micro HA (dashed line) and nano HA (solid line)
scaffolds as a function of time. 083 were seeded onto porous scaffolds (n = 3 with each trial
run in triplicate) at a density of 11,320 cells/cm2 and counted at intervals of 0.5, 1, 2 and 4
hours. Counts were completed after trypsinizing, morselizing and using a hemacytometer.
Values are mean t SE, p < 0.05 at all time intervals for comparison of nano HA to micro HA.
Figure taken from Smith et al., Int J Nanomed 2006;1:189-194.

151

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

250000

200000 4
1’ III HA El nanoHA I
3 150000 -
I.-
° I
h
3
E 100000 - I
3
Z I

50000 I .I. I
O I _. .. T
1 3 5
Time (days)

Figure 4: MC3T3-E1 Osteoblast proliferation on micro HA and nano HA. 088 were seeded
onto porous scaffolds (n = 3 with each trial run in triplicate) at a density of 11,320 cells/cm2
and counted after periods of 1, 3 and 5 days. Counts were completed after trypsinizing,
morselizing and using a hemacytometer. Values are mean t SE. p > 0.05 for comparison of
nano HA to micro HA porous scaffold for 1 and 3 days. p < 0.05 for 5 days (1:). Figure taken
from Smith et al., Int J Nanomed 2006;1:189-194.

152

+ nano HA adjusted .
1200 - + micron HA adjusted . ' ' ' i I

--O--nanoHA
1000 - --l--micron HA . .'

 

% Cells Attached

 

 

0 I I l I I j
0 20 40 6O 80 100 120

Time (hrs)

Figure 5: MC3T3-E1 OB attachment (0.5, 1, 2 and 4 hrs) and proliferation (24, 72 and 120
hrs) on porous nano-HA and micron-HA scaffolds. Again, 083 were seeded onto porous
scaffolds (n = 3 with each trial run in triplicate) at a density of 11,320 cells/cm2 and counted at
intervals of 0.5, 1, 2 and 4 hours. Counts were completed after trypsinizing, morselizing and
using a hemacytometer. Background counts were completed using scaffolds morselized in
the absence of CBS and subtracted from the raw data (dotted lines), to yield adjusted values
(solid lines). Figure taken from Chapter 4.

The question remains, however, as to what feature(s) associated with
reduced grain size in HA lead(s) to increased OB attachment. Therefore, two
different grain boundary-associated phenomena in dense HA scaffolds were
investigated. In Chapters 5 - 7, the effect of microcracking on OB attachment
was evaluated, by inducing microcracks of a known size using a Vickers
indentor and in Chapter 8 at the contribution of grain boundary grooving on
OB attachment was assessed.

Microcracks resulting from thermal processing of HA were first
identified in our work reported in Chapter 5, Part A.1 In this study, three

different CaP materials were sintered and examined using scanning electron

153

microscopy (SEM). Dense HA specimens (97.7 i 0.7% of theoretical) with an
average grain diameter of 7.9 i 1.5 pm contained microcracks both along

grain boundaries and traversing individual grains.

 

Figure 6: SEM micrograph of an as-sintered hydroxyapatite specimen surface. Image was
taken away from the fracture surface (4500x). This micrograph shows multiple examples of
microcracking due to thermal expansion in HA, indicated by arrows. White arrows indicate
cracks crossing the grain bulk and black arrows indicate cracks along grain boundaries. This
type of microcracking Is not easily controlled. Image taken from Case et al., Mater Sci Eng A
2005;390:246—254.

Naturally, we wished to study the effect of these microcracks on OB
behavior in vitro. Bone microdamage including microcracking, has been
shown to play a significant role in triggering bone remodeling and we
therefore hypothesize that microcracked scaffolds would show improved OB
activity. However, the microcracks found in this study resulted from thermal
expansion anisotropy (TEA), and only occur above a critical grain size (GCR)

for a given material. For HA in particular, the GCR is difficult to estimate and,

154

therefore, the degree of microcracking due to TEA is difficult to control. We
needed to create microcracked HA specimens, where the microcrack number
and size could be controlled.

Therefore, we induced controlled microcracks (2cAVE = 124.4 pm i
21.1 pm) in dense HA, using Vickers microindentation (Figure 7). In Chapters
6 and 7, we observed the effect of this induced microcracking on OB
attachment for dense HA at one and four hours. These results are shown
again here, in Figure 8. At one hour, the OB attachment was 25.1 i 8.9% of
the total number of 033 plated on the control HA and 31.3% i 12.7% of the
total number of 083 plated on the microcracked HA, with a p-value of 0.25 as
determined by Student’s t-test. This shows that there was no significant
difference between the OB attachment on microcracked and non-
microcracked HA specimens. In contrast, in a separate trial at four hours,
29% i 7.6% of the total 083 attached to the microcracked HA versus 18% i-
6.1% on the control HA, a 61.1% increase. Student's t-test showed that this

was a significant increase in the percentage of OB attachment (p = 0.004).

155

 

 

Figure 7: SEM micrograph of microindentation on HA specimen surface (right). "2a" indicates
the diagonal length of the indentation impression and "2c" indicates crack length (left). This
indent was created using a Vickers microindentor with a load of 4.9 N, a load time of 5
secnds and a loading rate of 70 um/sec. This method of microcrack induction created
microcracks of an average size of 2cm; = 124.4 um i 21.1 um and is more easily controlled
compared to TEA-induced microcracking. Indents such as this one were placed in a 14x14
grid across the specimen surface. Image taken from Smith et al., Am J Biochem Biotechnol
2006;21105-110.

156

50.0 - 40.00 - IControl

 

 

 

 

   

I Control I Microcracked
I Microcracked #
40-0 '0 30.00(
D 2
O
z 0
§ 30.0 g
< 2 20.00 'I
E C
8 20.04 8
g 8
0- 10.00 ~
10.0 -
0‘0 _ 0.00 1—
1 hr 4 hrs

Figure 8: Osteoblast attachment at 1 and 4 hours on microcracked and non-microcracked
dense HA discs. 085 were seeded onto dense microcracked and control HA scaffolds (n = 3
with each trial run in triplicate) at a density of 11,320 cells/cm2 and counted at 1 and 4 hours.
Counts were completed after fixing the cells using formaldehyde and staining with both
Rhodamine/Phalloidin and Hoescht. Stained cells were viewed under an optical fluorescence
microscope and counted in 10 fields of view. Counts were then averaged for each group of 9
specimens (control and microcracked) and the results shown here, with the error bars
indicating standard deviation (as opposed to SE, used in Figures 3 and 4). “#" indicates that
difference in OB attachment between control HA and microcracked HA at 4 hours is

significant (p = 0.004). Figure taken from Smith et al., submitted to J Biomed Mater Res,
December 2006.

At both 1 and 4 hours, the presence of microcracks at the HA surface
did not have a significant effect on OB morphology was shown in figures 9
and 10. In Figure 9, 083 possess either a spread or polygonal morphology on
both control and microcracked specimens. In Figure 10, this morphology is

also apparent for both groups.

157

100 um

 

Figure 9: A) MCST3-E1 033 with spread (white circle) and polygonal (black circle)
morphology on HA control specimen. B) MC3T3-E1 OBs with similar spread (white circle)
and polygonal (black circle) morphology on microcracked HA specimen. 035 have been
fixed and stained to highlight actin fibers (red - Rhodamine-Phalloidin) and cell nuclei (blue —
Hoechst). Images taken from Smith et al., submitted to J Biomed Mater Res, December
2006.

 

Figure 10: 10A shows MC3T3-E1 osteoblasts with both spindle-shaped and polygonal
morphology on HA control specimen after 4 hours culture. 085 were fixed and stained to
highlight actin fibers (red — Rhodamine-Phalloidin) and nuclei (blue — Hoechst). 108 shows
MC3T3-E1 Osteoblasts with spindle shaped morphology on HA microcracked specimen at 4
hours, again stained to highlight actin fibers (red — Rhodamine-Phalloidin) and nuclei (blue -
Hoechst). Images were taken using an optical fluorescence microscope. Images taken from
Smith et al. J Am Biochem Biotechnol 2006;2z105—110.

These findings indicate that microcracking does indeed have a positive
effect on OB attachment in vitro at four hours, and should be investigated
further in order to understand its effect on later stage OB behavior, including

proliferation, differentiation and calcium production.

158

While the presence of surface-breaking microcracks in HA scaffolds
positively influences OB attachment, it is not likely that this phenomenon
contributed to increased OB attachment on the nano-grained dense scaffolds
investigated by Webster et al. who found increased OB attachment in nano-
scale grain diameter materials (grain diameter < 67 nm). This grain size is
below the GCR for TEA-induced microcracking in HA. Since Webster et al. did
not induce microcracking, it is unlikely that any microcracks were present in
their specimens, due to their very small grain size (< 100 nm).

The second grain boundary phenomenon investigated in this
dissertation is grain boundary grooving in dense HA, brought on by thermal

etching at 1100°C for 45 minutes following sintering (Figure 11). Grain
boundary grooving is estimated to be on the order of 0.5 pm. The surface

roughness of grain boundary grooved HA was 10.23 i 1.8 nm, which is lower

than the Ra = 19 - 32 nm reported earlier for nanophase ceramics“.

159

 

Figure 11: SEM micrograph of the surface of a grain boundary grooved HA specimen. Grain
boundary grooving was accomplished by thermally etching sintered dense HA through
heating at 1100°C for 45 minutes. Grain boundary grooves shown here are estimated to be
approximately 0.5 um In width (hillock to hillock). Image taken from Chapter 8.

MCST3-E1 033 were seeded onto grain boundary grooved and control
dense HA discs for one and four hours. The results are shown in Figure 12.
At one hour, the mean OB attachment was 45.6 i 12.2% on control HA and
47.3 i 12.8% on grain-boundary grooved HA, a non-significant difference (p =
0.78). Similarly, at four hours no significant change in the mean OB
attachment occurs between the control HA (59.0 i 13.0%) and grain
boundary grooved HA (59.3 i 23.9%) (p = 0.98). This finding indicates that
the contribution of grain boundary grooving to the increased OB attachment

found for nano-scale HA versus micron-scale HA is not significant.

160

7o - IControl 90 _ Icmtm'

 

 

 

 

 

 

 

 

 

 

IZIGrain Boundary Grooved
IGrain Boundary Grooved 80 . --
60 -
70 -
E 50 - E 60 - _-
E E F
5 40 ' 5 so -
g 3:"
< < 40 _
e 3°~ 3 ..
3 g 30 -
20 4 20 _
10 ~ 1o -
o 0 1
1 hr 4 I'll‘

Figure 12: MC3T3-E1 OB attachment at 1 (Figure 2A) and 4 (Figure 2B) hours on thermally
etched and control HA specimens. 085 were seeded onto dense grain boundary grooved
and control HA scaffolds (n = 3 with each trial run in triplicate) at a density of 11,320 cells/cm2
and counted at 1 and 4 hours. Counts were completed after fixing the cells using
formaldehyde and staining with both Rhodamine/Phalloidin and Hoescht. Stained cells were
viewed under an optical fluorescence microscope and counted in 20 fields of view. Counts
were then averaged for each group of 9 specimens (control and microcracked) and the
results shown here, with the error bars indicating standard deviation (as opposed to SE, used
in Figures 3 and 4). At both time intervals there is no significant difference between groups.
Figure taken from Chapter 8.

This study also found that grain boundary grooving had no significant
effect on OB morphology, as shown in Figure 13. While the OB spicule length
is shorter at 1 hour than at 4 hours (12 pm i 8 pm for control and 18 um i 8
gm on grain boundary grooved at 1 hour. 24.8 pm i 10.3 pm for control and
28.8 um :I: 9.4 pm for grain boundary grooved at 4 hours.) there is no
difference in OB morphology between control and grain boundary grooved

groups at either time point.

161

 

Figure 13: This set of four optical micrographs shows attached 085 at one hour on the
surface of a control specimen (Figure 3A) and a grain boundary grooved specimen (Figure
3B), and also at four hours on the surface of a control specimen (Figure 30) and a grain
boundary grooved specimen (Figure 3D). 085 were fixed and stained to highlight actin fibers
(red — Rhodamine-Phalloidin) and nuclei (blue - Hoechst). The lesser degree of OB spread in
both groups at 1 hour Is a result of the shorter culture time. Images were taken using an
optical fluorescence microscope. Images taken from Chapter 8.

Conclusions

These results rule out microcracking and grain boundary grooving as
contributing factors in increased OB attachment on nano-scale HA substrates,
as microcracking due to TEA does not occur at the grain size reported and
grain boundary grooving does not significantly Increase OB attachment.
However, this does mean that a porous nano-grained scaffold may not exhibit
increased OB attachment were cracks introduced via other means, such as

an applied mechanical force. This indicates that the role that grain boundaries

162

play in increased OB attachment for nano-grained HA is related to either the
presence of a grain boundary phase or variation in surface charge associated
with those regions.

Ceramic grain boundaries often possess a precipitated solute or a
separate phase.5'6 An amorphous phase has been shown to form during
sintering in several polycrystalline ceramics including Si3N47, high alumina
ceramics8 and ZI'029. As the grain size is reduced, the grain boundary length
per unit area increases. With the presence of a continuous amorphous grain
boundary phase the surface area fraction of this glassy phase increases as
well. Glassy biomaterials have been shown to increase appositional bone
growth when compared to crystalline biomaterials10 where the presence of a
continuous amorphous phase has been linked to enhanced OB attachment
and eventual bone growth. Glassy biomaterials, such as 45S5 Bioglass, have
been shown to promote bone formation in vivo and are being evaluated for
potential applications.11 Recently, It was discovered that nano-crystalline and
nano-amorphous HA elicited similar OB responses to RGD functionalized
conventional HA.12 Even without the presence of an amorphous grain
boundary phase, the decrease in crystallinity associated with decreased grain
size may promote OB attachment and should be investigated further.

Using transmission electron microscopy (TEM), an amorphous, glassy
grain boundary phase was found on nano-grained HA taken from dental

enamel.13 This amorphous phase did not develop as a result of sintering and

observation of non-biological HA using TEM found no evidence of an

163

amorphous grain boundary phase. This is consistent with the lack of
amorphous grain boundary grain observed in sintered HA by Kleebe et al,
where high-resolution electron microscopy (HREM) was used to image grain
boundaries in sintered undoped HA.14

Increased surface roughness resulting from decreased grain size Is a
likely contributor to improved OB behavior. In addition to several studies by
Webster et al., in which increased surface roughness in ceramic materials led
to increases in OB attachment and proliferation, OB behavior has been

shown to improve with increased surface roughness on various biomedical

15,16 I 17

alloys, including Ti alloys and 316L stainless stee

When viewed in concert with prior work, and by effectively ruling out
microcracking and grain boundary grooving as determining factors in OB
attachment on nanograined dense HA, our research suggests that increased
OB attachment may result from increased porosity and surface topography.
The nanometer-scale roughness found in nano-grained ceramics is most
similar to that found in living bone.18 Therefore, while grain boundaries may

not directly improve OB attachment, they are another surface feature within a

more biomimetic topography.

164

REFERENCES

1.

Case ED, Smith IO, Baumann MJ. Microcracking and porosity in
calcium phosphates and the implications for bone tissue engineering.
Mat Sci and Eng A. 2005;390:246-254.

Webster TJ, Ergun C, Doremus RH, Siegel, RW Bizios R.
Nanocrystalline hydroxyapatite enhances osteoblast function. In: The
Proceedings of the First Joint BMES/EMBS Conference, Atlanta, GA;
1999. p744.

Webster TJ, Siegel RW, Bizios R. Design and evaluation of nanophase
alumina for orthopaedic/dental applications. Nanostruct Mater
1999;12:983-986.

Webster TJ, Siegel RW, Bizios R. Osteoblast adhesion on nanophase
ceramics. Biomaterials 1999;20:1221-1227.

Kingery WD. Plausible concepts necessary and sufficient for
interpretation of ceramic grain-boundary phenomena: I, Grain-
boundary characteristics, structure, and electrostatic potential. J Am
Ceram Soc 1974;57:1-8.

Kingery WD. Plausible concepts necessary and sufficient for
interpretation of ceramic grain-boundary phenomena: ll, Solute
segregation, grain-boundary diffusion, and general discussion. J Am
Ceram Soc 1974;57:74 - 83.

Jin Q, Ning XG, Wilkinson DS, Weatherly GC. Redistribution of a grain-
boundary glass phase during creep of silicon nitride ceramics. J Am
Ceram Soc.1997;80:685—691.

Powell-Dogan CA, Heuer AH. Devitrification of the grain boundary
glassy phase in a high-alumina ceramic substrate. J Am Ceram Soc
1994;77:2593—2598.

Kobayashi S. X-ray microanalysis of the boundary phase in partially
stabilized zirconia (PSZ). J Mat Sci Letters 1985;42268-270.

10.Fujishiro Y, Hench LL, Oonishi H. Quantitative rates of in vivo bone

generation for Bioglass and hydroxyapatite particles as bone graft
substitute Journal of Materials Science: Materials in Medicine
1997;8:649 - 652.

165

11.Lu HH, Pollack SR, Ducheyne P. Temporal zeta potential variations of
4585 bioactive glass immersed in an electrolyte solution J Biomed
Mater Res 2000;51:80-87.

12.Anselme K, Bigerelle M, Noel B, Dufresne E, Judas D, lost A, Hardouin
P. Qualitative and quantitative study of human osteoblast adhesion on
materials with various surface roughnesses. J Biomed Mater Res
2000;49: 155-166.

13.Kleebe HJ, Bres EF, Berache-Assolant D, Ziegler G. High-resolution
electron microscopy and convergent-beam electron diffraction of
sintered undoped hydroxyapatite. J Am Ceram Soc 1997;80:37—44.

14.Balasundaram G, Sato M, Webster TJ. Using hydroxyapatite
nanoparticles and decreased crystallinity to promote osteoblast
adhesion similar to functionalizing with RGD. Biomaterials
2006;27:2798-2805.

15.Wen 8, Liu Q. High resolution electron microscopy investigations of
interface and other structure defects in some ceramics. Microsc Res
Tech 1998;40:177-186.

16.Webster TJ, Ejiofor JU. Increased osteoblast adhesion on nanophase
metals: Ti, Ti6Al4V, and CoCrMo. Biomaterials 2004;25:4731—4739.

17.Hao L, Lawrence J, Phua YF, Chian KS, Lim GC, Zheng HY.
Enhanced human osteoblast cell adhesion and proliferation on 316 L8
stainless steel by means of C02 laser surface treatment J Biomed
Mater Res B: Appl Biomat 2005;73: 148-156.

18.Kaplan FS, Hayes WC, Keaven TM, Boskey A, Einhom TA, Iannotti

JP. In: Simon, SR, editor. Orthopaedic basic science. Columbus, OH:
American Academy of Orthopaedic Surgeons, 1994: p127.

166

CHAPTER 10

FUTURE WORK

Our results naturally suggest two paths of future study. The first is a
further investigation into why a decrease in grain size leads to an increase in
OB attachment. Although this work has ruled out microcracking and grain
boundary grooving as contributing factors to increased OB attachment in HA,
and others have shown no evidence of a separate grain boundary phase in
non-biological HA, this reduction in grain size has still been shown to
positively affect OB attachment in vitro."3 Specifically future work will
determine whether this relationship is due to differences in surface charge
associated with grain boundaries in HA.

The second path is to further study the effect of microcracking on bone
regeneration. We have determined that the presence of induced microcracks
in HA increases OB attachment and we hypothesize that this is an indication
that the presence of microcracks is initiating the bone remodeling process. In
living bone, remodeling begins with damage to adjacent osteocytes, which
apoptose, signaling continued remodeling. Future work will investigate
whether induced microcracking leads to osteocyte apoptosis. Additionally, in
vitro/in vivo studies using porous, nano-grained scaffolds with induced

microcracking will be conducted.

167

Surface charge variation and DB attachment on HA

A factor which has not yet been linked to increased OB attachment is
nanoscale surface charge variation associated with grain boundaries in
ceramics. In the 1974 Sosman Memorial Lecture at the American Ceramics
Society, Kingery supported the concept that grain boundaries carry an electric
charge, by measuring associated space-charge clouds in NaCl (-), AI203 (+)
and MgO H.4

In our own work‘T8 (Appendix E-F), as well as numerous other studiesg'

1‘ electrophoresis has been used to measure the C-potential, which is an

estimate of the surface potential, associated with particulate biomaterials.
This technique works well to quantify the average potential across a material
surface, but does not effectively quantify the charge variations associated
with grain boundaries. Vandiver et al. have more recently employed high-
resolution force spectroscopy to measure directly the force required to
separate the probe tip from the surface in HA.12 By collecting these
measurements across several grains at the HA surface, the surface charge
was approximated and found to be higher near grain boundaries. This
increase in surface charge with finer grain sizes should more effectively
attract the negatively charged 0836, and explains the increase in OB
attachment found in nano-grained HA.

To determine whether this is true, dense HA substrates will be
fabricated using the techniques described in Chapters 5 - 8. By using a nano-

scale starting HA powder and varying the sintering time, HA discs will be

168

fabricated in two distinct groups: 100 nm grain diameter and 1 pm grain
diameter (as determined by SEM). These grain sizes are below the accepted
values of GCR for HA and have been chosen to prevent TEA-induced
microcracking. Both specimen groups will be polished to > 1 pm and the
average surface roughness determined using atomic force microscopy (AFM).
By assuring that that the difference in mean surface roughness between
these two groups is not significant, the effect of increased surface roughness
on OB attachment is eliminated.

MC3T3-E1 083 will then be seeded onto both specimen groups and
cell attachment quantified at 0.5, 1, 2 and 4 hours, using optical microscopy.
OBs will be fixed and stained using Hoescht (to image cell nuclei) and
Rhodamine-Phalloidin (to image cell morphology) fluorescent stains, then
imaged using a Leica DM IL optical fluorescence microscope (Leica
Microsystems, Wetzlar, Germany) equipped with a SPOT RT camera system
(SPOT Diagnostic Instruments, Sterling Heights, MI). 083 from 20 fields of
view will be counted, to create an accurate representation of OB attachment
across the HA surface. Students t-test will be used to determine the

significant difference between groups.

Future work in microcracking and DB attachment on HA
The findings included in the previous chapter effectively rule out
microcracking and grain boundary grooving as phenomena that contribute to

increased osteoblast (OB) attachment associated with reduced grain size in

169

dense hydroxyapatite (HA).10'12 However, microcracking has a positive effect
on OB attachment when it is present in larger-grained HA. This new finding13
merits a further investigation into this relationship, and whether or not the
observed increase is tied to the role of microdamage in bone regeneration.
The idea that microcracking in bone results from fatigue and is linked
to the remodeling process was originally presented by Frost in 196014 and is
discussed in detail in Chapters 6 and 7. Targeted remodeling15 occurs at sites
of localized microdamage as observed in studies by Johnson et al., Burr et al.

and Mori et al.1619

where osteoclasts act to resorb damaged bone, including
osteocytes.20 This resorption is regulated by the condition of the osteocyte
synctinium in these local regions and osteocyte damage and subsequent
apoptosis leads to increased osteoclast activity.

The osteocyte is one of three bone cell-types. The other two are the
osteoblast (bone-forming cell) and osteoclast (bone-eroding cell).21
Osteocytes begin life as osteoprogenitors, which differentiate into osteoblasts.
These osteoblasts help in the formation of osteoid. Some osteoblasts become
trapped within the newly formed osteoid and become osteocytes, residing
within the lacunae. As the osteoid matrix forms, the mineralization front
surrounds the osteocytes. Once surrounded, the cells extend long, thin
processes toward the vasculature. It is through these processes, as well as
gap junctions, that the osteocytes form a functional syncytium.22

The osteocyte syncytium is the subject of interest in the proposed

study. It is theorized that a relationship exists between osteocyte health and

170

bone volume, with a higher cell number causing increased extracellular matrix
volume.23 Osteocyte number density has also been tied to the amount of
microdamage accumulated in bone.24 It has also been theorized that
osteocytes undergo apoptosis with the incidence of linear microcracking, prior
to bone remodeling”28 leading to the supposition that apoptosis is a key part
of the resorption process, by serving to direct osteoclasts to damaged sites.29

To determine whether this process is initiated by induced
microcracking in HA, discs will be prepared by uniaxial pressing as described
in Chapters 5 - 8. A starting powder with a diameter small enough to yield a
sintered grain size > GCR will be used to avoid TEA-induced

303‘. Rather, controlled microcracks will be induced using a

microcracking
Vickers microindentor as described in Chapters 6 - 7.

To determine the spatial relationship between osteocyte apoptosis and
microcracking, samples will be prepared by creating discreet regions of
microcracking, using Vickers indentation.

Osteocytes will be cultured in the presence of these dense HA discs.
The number of apoptotic cells will be assessed after a prescribed time and
cell density and apoptosis will be reported as a function of average
microcrack length (mm) and density (cracks/mmz). Cell density will be

determined by staining and counting with an optical microscope. Statistical

significance will be determined using the Student's t-test.

171

Additionally, the proximity of apoptotic cells to regions of microcracking
will be assessed using scanning electron microscopy (SEM), confocal laser

scanning microscopy (CLSM) or fluorescent light microscopy.

172

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Smith IO, Soto MK, Baumann MJ, McCabe L. In vitro stability
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Smith IO, Baumann MJ, Obadia L, Bouler JM. Surface potential and
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Ducheyne P, Kim CS, Pollack SR. Effect of phase differences on the

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15.Burr DB. Targeted and nontargeted remodeling. Bone 2002;30:2-4.

16.Johnson KA, Skinner GA, Muir P. Site-specific adaptive remodeling of
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27.Verborgt O, Gibson GJ, Schaffler MB. Loss of: osteocyte integrity in
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175

APPENDIXA
THREE-DIMENSIONAL MICROSTRUCTURAL CHARACTERIZATION OF

POROUS HYDROXYAPATITE USING CONFOCAL LASER SCANNING
MICROSCOPY (CLSM)

FEI REN, IAN 0. SMITH, MELISSA J. BAUMANN AND ELDON D. CASE
Department of Chemical Engineering and Materials Science Department
Michigan State University, East Lansing, MI 48824

PUBLISHED IN: lntemational J Appl Ceram Technol 2005;2z200-211.

ABSTRACT

The Characterization of porosity IS crucial in the development and
commercialization of ceramic bone replacement technology, Since the pore
size and interconnectivity play a central role in both biological function (bone
ingrowth and nutrient flow) as well as mechanical properties of bone
scaffolds. The ability of Confocal Laser Scanning Microscopy (CLSM) to
image three-dimensional structures with large vertical depths (~ 2 mm) and
fine vertical resolution (~ 1 pm) is utilized in this paper to characterize the 3-D
microstructures of hydroxyapatite (HA) bone scaffold specimens with porosity
ranging from roughly 60 to 70 volume percent. Various CLSM techniques are
applied to image and interpret the HA pore structure, including Z-series

stacking, topographic profiling, and Phi-Z scanning and contour mapping.

176

INTRODUCTION

The Confocal Laser Scanning Microscope (CLSM) has become, over the
last 10 to 15 years, a standard instrument for the study of soft biological
tissues. Many studies of normal and tumor cells1, biofilms 2, proteins 3, brain
tissue 4, skin tissue 5, corneal tissue 6, etc. are published each year. Despite
the widespread use of CLSM for soft tissue research, there is still very limited
use of CLSM for "hard biomaterials" (such as bone and teeth) or more
generally for the field of materials science. This study utilizes the CLSM to
Characterize the porosity in hydroxyapatite (HA) that is the major component
in the burgeoning demand for ceramic bone tissue engineering replacement
materials.

Disease and traumatic injury result in a critical need for bone replacement
materials, as evidenced by the large number of bone grafts (more than
500,000 annually in the U. S. 7. While bone tissue harvested from alternate
sites in the patient (autograft) may often be an ideal replacement for the case
of spongy (cancellous) bone, load bearing autografts are problematic
because removing load bearing (cortical) bone from a donor site in a patient
may require a repair/replacement at that site. Although bone harvested from
other donors (allograft from cadavers) can be ground and used as part or all
of such replacement material, the possibility of disease transmission makes
the use of allograft bone tissue problematic. Thus, as an attractive alternative,
there has been considerable work done to develop artificial bone replacement

materials consisting of bioceramics such as hydroxyapatite (HA). Currently,

177

approximately 60% of the bone graft materials use calcium phosphate-based
materials which are commercially available as a paste, putty, solid matrix and

as granules 7

and are supplied by all of the world’s major orthopaedic
corporations. Because of the biological and mechanical necessity of forming a
strong bone/bioceramic interface, the development and long term success of
these synthetic bone tissue scaffolds is highly dependent on characterizing
the pore size, morphology and interconnectivity, since both the biological and
mechanical functions of such scaffolds depends on the details of the pore
microstructure of these materials.

In this study, we utilize the CLSM confocal reflection mode to
characterize the surface topology of both macropores and micron-scale
features in porous HA bone scaffold materials. In addition, we use the
confocal fluorescence mode to image surface features. As part of this study,
we have begun the process of modifying the experimental procedures
developed by Fredrich et al. (to study porous geological materials) 8 so that
the fluorescence mode can be applied to the considerably larger pores found
in our bone scaffold materials. In total, we have examined 11 HA bone
scaffold specimens (6 with epoxy impregnation and staining and 5 in the as-

sintered state). This study presents CLSM images from 11 specimens having

approximately 60 to 70 volume percent porosities.

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BACKGROUND

Porous HA ceramics: biological and mechanical implications of porosity
To underscore the importance of characterizing the porosity of HA

scaffolds, we Shall first discuss the central role that porosity plays on both the

biological and mechanical properties of these materials. Moreover, for the

biological function of bone, a bimodal pore architecture is needed. The bone

tissue engineering community classifies pores about 200 - 400 um across as

macropores and interconnecting pores roughly 60-70 um in diameter as

“micropores”.

Biological implications of porosity

For synthetic bioceramic bone scaffolds, bimodal porosity is essential for
bone-ingrowth as well as the attachment of the ceramic scaffold to the bone 9
12. These pores assist the process of bone ingrowth into the synthetic bone
scaffold materials (fabricated from porous HA) where the newly forming bone
penetrates into the pores of the scaffold and if a resorbable phase of HA is
used for the scaffold, eventually replace the ceramic with neobone 13. Hing et
al. have found that microporosity is especially important for the early stages of
bone ingrowth (osseointegration), in that increasing the volume fraction of
microporosity can improve the bioactivity of the porous HA 11. In addition to

the "early biological response", the "mechanical properties of a porous

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hydroxyapatite bone graft substitute are highly sensitive to its pore structure"
12

A bimodal pore Size distribution also is important in the controlled release
of antibiotics from porous HA bone scaffold materials 14, especially in terms of

interconnective pores between the micro and macro porosity 10.

Impact of porosity on mechanical properties of bone scaffolds
For ceramic materials, including bioceramics, porosity leads to exponential

decreases in the elastic modulus 15’19, hardness 2° and fracture strength 18'19'

21' 22 along with dramatic decreases in fracture toughness 23 and the critical
strain energy release rate 23. For bioceramics, such as HA in particular,
porosity directly affects the tensile strength 24, the compressive strength 24' 25
and the hardness 20. Thus, the mechanical integrity and viability of the bone

scaffold materials is also closely linked with the details of the pore

architecture.

Using the CLSM to characterize porous HA specimens

One method of examining the pore distribution (size and spatial
arrangement) is by means of serial sectioning. Typically, this is done by
~ mechanical sectioning, where the specimen is sectioned (cut) by a saw to
reveal the sub-surface pore structure. Disadvantages of serial sectioning by
mechanical means of course include the inevitable damage caused by the

sectioning process, including the portions of the specimen that are destroyed

180

in the saw kerf. (For a typical low speed diamond blade that is on the order of
0.35 mm thick, 10 diamond saw cuts will destroy the equivalent of a 3.5 mm
thick slab of the specimen). Mechanical sectioning also is tedious and time
consuming and in addition to the damage caused by the saw kerf, additional
specimen damage is incurred in brittle materials due to chipping and spalling
near the cut.

Of course, it would be preferable if the specimen could be sectioned either
by optical or electronic means, thus eliminating the damage generated by
repeatedly cutting the specimen. For the CLSM, since the image is
constructed from a series of images taken at slightly different focal planes, the
image (rather than the specimen) can be sectioned or rotated electronically.

For topology studies using the CLSM, no surface preparation is required.
However, for standard scanning electron microscope samples, coating the
surface is required as well as placing the specimens in a vacuum. Even for
the environmental scanning electron microscope (ESEM), the specimen
needs to be viewed from within a chamber that requires partial vacuum. In
contrast, specimens can be examined in the CLSM at one atmosphere
ambient pressure, at room temperature and without a conductive coating. Of
course, for the CLSM, if one desires to immerse the specimen in water or
place the specimen in a particular atmosphere, this can be done in principle
using an environmental chamber. In addition, the shape and Size of
specimens observed using CLSM are much more flexible than those using

SEM.

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In addition to the CLSM, there are alternative methods to image porous
bioceramic and bone materials. For example, micro-computed tomography
(uCT) is utilized for non-destructive two-dimensional”: 27 and three-

'28

dimensiona image of bone and HA, with resolutions ranging from about 1 to

100 pm. The best resolution (about 1 pm) for uCT is possible only using

extremely high-flux sources, such as x-ray synchrotron radiation. For
example, the European Synchrotron Radiation Facility was recently used to

image bone ingrowth in calcium phosphate bioceramics in three
dimensions.” The effective resolution was approximately 1.4 um for long rod

specimens 0.6 mm X 0.6 mm X 10 mm, examined perpendicular to the long
axis of the rods.” (It is relatively typical that for three-dimensional images
formed by x-ray synchrotron radiation that the characteristic dimension of the
specimen is on the order of 1 mm or less). Also, it should be noted that
synchrotron x-ray radiation facilities are large, centralized facilities for which
one competes with other researchers to schedule beamtime on the facility.
Thus, with synchrotron radiation, the specimen sizes are relatively small, the
equipment is available only at centralized facilities and the high x-ray flux
rules out applying the technique to studies of bioceramics or bone on which
there are living cell cultures. In contrast, the CLSM is a relatively
inexpensive, the specimen size is quite flexible, the maximum axial and radial
resolutions are on the order of a few microns, depending on the wavelength
of the laser source and the numerical aperture of the objective lens (Section

3.3). In addition, the CLSM is widely used for in-vivo studies in biological and

182

 

biomedical applications30' 31, thus the CLSM offers the potential to study, for
example, osteoblast cell cultures on bioceramics such as HA.

Magnetic resonance (MR) techniques, including high resolution MR
(hrMR) and micro-MR (uMR) are also used to image HA and bone
specimens. However, in 2005, Hudelmaier et al. noted that “clinical high-

resolution (hr) MRI scanners have recently reached a sufficient spatial
resolution (around 100 - 200 pm) to quantify microstructural parameters of
trabecular bone at peripheral Skeletal sites in vivo.” Also, current hrMR and
uMR does not easily allow for the compositing of two-dimensional image

Slices in order to form for three-dimensional images.32 Thus, while the hrMR

and the micro-MR have the ability to study cells in vivo, these techniques lack

the resolution of uCT or the CLSM.

EXPERIMENTAL PROCEDURE
Materials and specimen preparation
The porous hydroxyapatite (HA) specimens used in this study were

fabricated using medical grade, sintered HA powder (HiMed, Old Bethpage,
New York) with an initial particle size of approximately 10pm. For the most

porous samples, approximately 28.5 g of HA powder was combined with 13.7
ml 10'3 M KN03 solution, forming a slurry. To this slurry, 1.1 ml of 30% H202
was added. The modified slurry was then placed into silicone cylindrical
molds. The molds were covered with aluminum foil and placed in a gravity

convection oven at 125°C for 1hr to dry the cylinders. After drying, the molds

183

 

were removed and the (green) samples ejected. The green (unfired)
cylindrical specimens were approximately 1.3 cm in diameter and 1.6 cm in
length. Cylinders were then sintered in air for 4 hours at 1360°C, with a
heating/cooling rate of approximately 10°C/min. This procedure was repeated
for differing sample porosities, with modifications, as discussed below.

The HA specimens included in this study having porosities ranging from
approximately, 60 to 70 volume percent. Porosity was varied by altering the
concentration of H202 used in foaming, as well as the HA/KNO3 ratio in the
slurry. Porosity was determined either by Archimedes method or by
comparing the theoretical mass to the actual mass. For specimens having the
highest porosity, the specimens were foamed using a 30% H202 solution and
an HA/KNO3 ratio of approximately 2:1. Less porous specimens were foamed
using a higher HA/KNO3 ratio and a lower concentration of H202 33.

The specimens were foamed and then sintered. The sintered specimens
were cylindrical, approximately 3 cm in length and about 1 cm in diameter.
The tops of the molds were open during foaming, so that a mushroom-like
dome formed on top of the specimen as it foamed in the mold. Using a
diamond abrasive wafering saw at ~200 rpm, the domed-ends of the sintered
specimens were sliced off and the remaining cylinder roughly sectioned into 2
cm lengths. The dome section was not utilized because of inconsistent
porosity. These highly porous samples were either mounted whole or
sectioned further using a padded specimen holder to avoid crumbling. For

reflective CLSM mode, specimens were sectioned to 1 cm lengths prior to

184

 

viewing while the specimens viewed in fluorescence were cut to 2 mm
lengths. Thus, in this study, the maximum dimensions of the specimens were
determined by the size of the molds used to form the specimens. The
subsequent sectioning was carried out to conserve the amount of epoxy used
in the fluorescence mode specimens and hydroxyapatite used for both the
fluorescence and reflection mode specimens. Viewing much larger samples
is possible with the CLSM (Section 3.2).

Following a procedure developed by Fredrich 8 for imaging the pore
structure in 3-D geological specimens, 3 dye-impregnated epoxy was
prepared by dissolving Rhodamine B into a Spurr resin kit. The Spurr kit is a
four-part epoxy, consisting of Noneyl Succinic Anhydride (NSA),
Vinyleyclohexane (VCD), Dow Epoxy Resin 736 (DER) and
Dimethylaminoethanol (DMAE) normally combined at a ratio of 120:2021221 to
form the epoxy. The viscosity of the epoxy can be customized to the porosity
of the sample, increasing the amount of DMAE increases the viscosity. By
increasing the proportion of DMAE by a factor of 3 relative to that used by
Fredrich 8, we achieved an epoxy with the desired viscosity, allowing for more
effective pore intrusion. These components are then combined at 30:5:3:1
and Rhodamine added to the resulting mixture at 1:2000 by weight. The dye
was incorporated by stirring with a metal spatula for 5 minutes, followed by
transferring the mixture to a centrifuge tube and mixing via vortex for 1
minute. Any air incorporated during mixing was removed by centrifugation at

2000 rpm for 5 minutes resulting in an epoxy ready for use.

185

Ring-shaped sections of polyvinyl chloride pipe (PVC: OD = 2.15 cm, ID =
1.5 cm) were cut to 2.5 cm in length using a coping saw, then attached to
small petri dishes (Becton Dickinson, Franklin Lakes, NJ) using super glue.
HA samples were next placed into the newly formed reservoirs and the epoxy
preparation poured onto the HA. Care was taken to completely cover the HA
specimens with epoxy. The epoxy coated HA samples in their reservoirs were
then placed into a dark vacuum Chamber which was evacuated to 88 kPa
(0.87 atm) for 12 hours. Following the time in the vacuum chamber, the
specimens were placed into a gravity convection oven at 65 °C for 24 hours to
completely cure the epoxy. The bottom of the Petri dish was out along the
outer circumference of the PVC ring without disturbing the super glue bond
between the sample and the Petri dish leaving the sample attached to the flat
Petri dish surface. Samples were then sectioned to 2 mm at the petri dish
end, using a diamond wafering saw operating at 200 rpm. Exposed sample

surfaces were next sequentially polished using a diamond abrasive paste at 6

pm, 1 pm, and 0.1 pm, rinsed with R0 water and dried in air. Samples were
stored in the dark to avoid photobleaching.

Two types of Specimens were observed, namely (1) specimens that were
mounted without epoxy impregnation or staining and (2) specimens that were

both epoxy impregnated and stained.

186

CLSM Observation

Two different CLSM instruments were used for this study: (1) LSM 210
(Carl-Zeiss Inc. Jena, Germany) and (2) LSM 5 Pascal (Carl-Zeiss Inc. Jena,
Germany).

The LSM 210 CLSM is equipped with 488 nm, 514 nm Ar laser and 633
nm HeNe laser illumination. In the study of topology, a confocal reflection
mode was used with the 488 nm Ar laser illumination that has the shortest
wavelength available for this instrument and in turn gives the highest
resolution. To increase the signal to noise ratio, a low pass filter program was
applied to the topographical images after the images were collected.

The LSM 5 Pascal has the following laser sources: Ar laser (458 nm, 488
nm, and 514 nm) and HeNe laser (543 nm and 633 nm). For the study of pore
space impregnated with epoxy and stain, the 488 nm Ar laser was used. The
emission fluorescence was collected via a long pass filter LP560. A line
average of 4 was applied when collecting images to reduce noise.

All images included in this paper are composed of 512 pixels x 512 pixels.
The aspect ratio (X: Y) of the LSM 210 is 3:2, thus images collected on this
CLSM are thus elongated in X direction. The lateral resolution (LR) and axial
resolutions (AR) of the two CLSMs are dependent on the wavelength of the
laser and the numerical apertures (N. A.) of the objective lens. For LSM 5
Pascal CLSM, an objective lens of 20X magnification with N.A. of 0.5 was
used, giving lateral and axial resolutions are 0.69 pm and 1.95 pm,

respectively. For the LSM 210 CLSM, the following objectives lenses were

187

used: 5X (NA. = 0.15, LR = 1.98 pm, AR = 21.7 um), 10X (NA. = 0.30, LR =
0.99 pm, AR = 5.42 pm), 20X (NA = 0.50, LR = 0.60 pm, AR = 1.95 pm),
and 50X (NA. = 0.50, LR = 0.60 pm, AR = 1.95 pm).

The specimen size that can be accommodated in CLSM is very flexible.
The specimen stage of LSM 210 is 90 mm x 60 mm and the maximum
distance between the stage and the objective lens is 30 mm, while the stage
of LSM 5 Pascal is 80 mm x 60 mm large and the maximum distance
between the stage and the objective is 20 mm. Since there are no
impediments at the sides of the CLSM specimen stage, the CLSM could
easily accommodate specimens with larger lateral dimensions that the stage
dimensions, for example specimens on the order to 20 mm to 30 mm thick
and 100 mm in each lateral dimension could easily be accommodated, which
is a considerably larger specimen that can be placed in most SEM sample
chambers. Furthermore, the CSLM stages can be replaced with modified

stages or removed entirely in order to accommodate even larger specimens.

Scanning electron microscope comparison

A series of Scanning Electron Microscope (SEM) images were taken of
several selected areas on the HA specimens that coincided with CLSM
images using a JEOL 6400V SEM (JEOL Corporation, Japan) operated at an
accelerating potential of 2 keV. The specimens viewed in the SEM were
sputter coated with an approximately 21 nm layer of Au using a SC500

sputter coater (Emscope Laboratories, Ltd. Ashford, England).

188

RESULTS AND DISCUSSION

Using a variety of fluorescing labeling agents, CLSM has been used to
image the morphology and focal adhesions of human osteoblast cells on HA
34, TIOz/HA 35 and PMMA cement 36, locate focal cell adhesions of osteoblasts
on HAPEX (HA in HDPE) 36' 37, HA microspheres 34, assess the effect of
surface topography on osteoblast attachment for composites of HAPEX and
AWPEX (a glass-ceramic apatite-wollastonite in HDPE) 38 and estimate the

h 39 on dense hydroxyapatite implants. In a related

degree of bone ingrowt
study, 40 a CSLM was used to investigate the interconnected nature of the
defects arising in plasma sprayed hydroxyapatite coatings after immersion in
simulated body fluid as a function of time. An SEM was used to detect the
presence of cracks and pores but could not distinguish the degree of
interconnectivity. The CSLM however was able to detect an interconnected
defect network in the plasma sprayed HA coating.

However, while each of these studies demonstrates the usefulness of
CSLM to image morphological aspects of HA-cell interactions, to date, there
is little research examining the pore size and pore size distribution, including
the degree of pore interconnectedness, which has been linked to
improvements in neobone formation and ingrowth into 3D HA bioceramic
scaffolds.

In this study, .we employ two CLSM operational modes, namely the (i)

confocal reflection mode and (ii) the confocal fluorescence mode. The

reflection mode and the fluorescence mode allow different types of CLSM

189

 

images to be collected, that is, while the reflection mode is best suited for
imaging the surface profile of rough or porous surfaces, the fluorescence
mode allows one to highlight specific subsurface features by adding
fluorescent tracers.

For example, the CLSM confocal reflection was used by Draeger and
Case 4143 to image surface channels about 500 to 900 pm in diameter in

partially stabilized zirconia (PSZ) specimens. (Surface Channels can be

44 46 and

converted into bulk-penetrating channels by ceramic/ceramic joining
such channels can in turn be used to transfer a variety of fluids including
cooling fluids for electronic substrates 47.) For the surface channels in the

4‘43 it was essential to characterize the details of both the

P82 specimens
meso-scale channels dimensions and shape as well as the micron-scale
topography of the channel surface Since both (i) the meso-scale size and
shape of the channel and (ii) the micron-scale projections on channel walls
influence fluid flow and heat transfer in the channels. The ability of the CLSM
to image meso-scale objects (the channels in the P82 specimens or the
macropores in the bone scaffolds) as well as micron-scale objects (pores and
surface roughness on the walls of the P82 channels and the pores in the
bone scaffolds) is thus vital to characterizing ceramic components whose
function depends simultaneously on structures at the meso and micro scale.
The interest In the flow and storage of liquids and of gases in rocks has
led the geological community to employ the CLSM operated in fluorescence

mode to characterize three-dimensional porosity in geological materials 8' 48'

190

51. AS in this study, a porous specimen is impregnated with a low viscosity
epoxy that contains a fluorescent stain. Since rocks and ceramics do not
fluoresce, the stain can then help to delineate subsurface porosity when
illuminated by the CLSM.

There are both important differences and similarities between the details
of the pore architecture for rocks and bone scaffolds. The porosity in
sedimentary rocks can range from roughly 20 volume percent for some
sandstones to about 60 volume percent for diatomaceous rock 8, while the
volume percent porosity of the bone scaffold specimens examined here is
between about 60 and 70. In addition, the pore architecture for some
geological materials is generally analogous to that observed for ceramics
densified from powder compacts. For example, geologists refer to “nodal"
pores at triple-grain junctions and label as "throats” the interconnecting pores
(which are sometimes “sheet-like”) that run along the grain edges 48. This type
of pore geometry is consistent with the general pore architecture described by
German 52 for sintered powder compacts, although instead of the “sheet”
shaped pores that connect nodal pores in geological materials, tubular pores

tend to connect pores at triple grain junctions in sintered powder compacts 52.

In sandstones, the mean pore size is typically 25 to 40 um, with a

relatively small fraction of pores being as large as 100 to 250 um 48, while in
the bone scaffolds there is a macropore phase (pores having diameters of

200 um and larger) and a “micropore” phase of interconnecting pores of

diameter 60-70 pm. Thus, the successful use of CLSM for characterizing

191

8' 48 indicates that these methods

subsurface porosity in geological materials
should be applicable to ceramic bone scaffolding materials. In this paper, we
have begun the process of modifying and optimizing the geologists’

procedures to accommodate the differences in porosity and surface structure

characteristic of ceramic bone scaffold materials.

CLSM characterization of HA ceramic bone scaffold specimens

Each of the specimens described in this study (Figures 1 - 5) have
been sectioned by a low speed diamond saw prior to observation by either
the CLSM (all Figures except 4a) or a conventional SEM (Figure 4b). For the
Specimen shown in Figure 3, the specimen was first impregnated with epoxy
and then sectioned while all of the other specimens depicted here were
sectioned and observed without epoxy impregnation. This sectioning
accounts for the "flat" profile that is observed at the highest points on each

specimen.

192

 

 

Fig. 1: A series of four reflection-mode confocal laser scanning microscopy images taken at a
single location on hydroxyapatite (HA) specimen HA-70-AS-1: (a) an overlay image showing
both macropores (3150—300 mm) and micropores (860—70 mm), (b) a Phi-Z scan along “cut
line” A-B, with pore dimensions labeled directly on the profile, (c) a “wire mesh” topographic
three dimensional profile highlighting the larger surface features of the specimen while
supressing the fine surface detail, and (d) a contour map that uses a gray scale to indicate
depth relative to a basal plane.

193

 

Fig. 2: A series of reflection-mode confocal laser scanning microscopy images of specimen
HA-70—A8—2, taken at a single location on the specimen: (a) an overlay image, (b) a Phi-Z
scan

194

;1ng- -r,- *W-W’W'gfwfign #3?“an ( ”3|: WW“

 

Fig. 3: For a single location on the epoxy-impregnated and stained specimen HA-70-ES-1, a
series of fluorescent-mode confocal laser scanning microscopy images: (a) an overlay image,
(b) a Phi-Z profiles along the cuts “1 " and cut “2," where the bright regions correspond to the
stained epoxy that fluoresces and the dark regions hydroxyapatite (HA), which does not
fluoresce, and (c) a single 2 section image taken at a depth of about 40 mm beneath the
specimen surface. A pore throat (opening between two adjacent pores) is labeled on the
Image.

 

Fig. 4: A comparison between (a) an overlay confocal laser scanning microscopy (CLSM)
image and (b) a conventional scanning electron microscoope (SEM) image taken at the same
location with identical magnification. Although the SEM image has higher lateral resolution
than the CLSM image, the SEM is unable to simultaneously image the specimen surface and
within the pores because of the limited depth of field of the SEM.

196

 

 

 

 

Fig. 5: A series of reflection-mode confocal laser scanning microscopy (CLSM) images of
specimen HA-70-AS-2 taken with a higher lateral and vertical resolution than the CLSM
images in Figs. 1—5, including a: (a) Phi-Z scan and (b) topographical profile. All images were
taken at the same location on the specimen.

HA-70-AS-1, an as-sectioned (not epoxy impregnated and unstained)
hydroxyapatite specimen with about 70 volume percent porosity, was
observed using a Zeiss 210 CLSM operated in the confocal reflection mode.
Figure 1a shows an overlay image from a "100 z-sectioning series", meaning
that the image has been electronically constructed from 100 individual
confocal images. The total image depth is 1200 pm and the "thickness" of
each section in the image stack is 12 pm, that is, as the images were

collected by the CLSM computer, the focal plane was shifted up by 12 pm

197

after each image was collected. Specimen HA-70-AS-1 includes macropores
with diameters of roughly 150 to 300 um (Figure Ia), which is within the size
range of micropores that is typical of bone scaffold materials. In addition,
Figure 1a shows a portion of a micropore about 60 pm in diameter. (AS

discussed in Section 2.1, the bone tissue engineering community uses the
term “micropore” to describe pores with diameters of about 60 to 70 pm that
serve as connecting pathways for nutrient flow and bone ingrowth).The
details of the topography of an overlay image (Figure 1a) can be analyzed via
a ”phi-Z' scan (Figure 1b), a topographical map (Figure Ic) and a contour
map (Figure 1d). A phi-Z scan (Figure 1b) represents a top surface profile
along an operator-selected line within the field of view of the CLSM image.
For example, the lower half of the image in Figure 1b shows Line A-B placed
across a portion of the image shown in Figure 1a. The upper half of Figure 1b
presents the geometric profile of the specimen along the “cutting” line A-B.
Once the profile has been obtained, the CLSM software allows one to
measure vertical heights between any two points along the specimen profile.
For the pore shown in the phi-Z profile (Figure 1b), the horizontal distance
between points P1 and P2 Is 104 um and the vertical distance between points
P1 and P3 is 148 pm. Thus, the phi-Z scan allows one to select particular
features of interest and then to determine their dimensions.

Figure 10 gives a three-dimensional topographic profile for specimen HA-
70-AS-1 that corresponds to the overlay image presented in Figure 1a. After

an overlay image has been collected, a topographic profile image as well as

198

 

 

 

contour images (Figure 1d) and an arbitrary number of phi-Z scans (such as
Figure 1b) can be constructed electronically from the overlay image. Phi-Z,
contour and topographical images also can be constructed from overlay
images and can be collected independently.

The topographic image has two advantages over the overlay images.
First, the topographic map's “wire mesh image” shows clearly the overall
surface profile of a specimen by suppressing the information on the
specimens fine surface features. A second advantage of the topographic map
is that one can choose to collect the topographical map separate from an
overlay image, since the topographical map itself takes far less time to
accumulate than a complete overlay image. For example, whereas for an
overlay image that takes 10 minutes to accumulate, a comparable
topographical image can be obtained in perhaps five minutes or less. Thus,
one can more quickly and efficiently “survey” the topographical features using
a series of topographical maps as opposed to a comparable series of overlay
images. This implies that the CLSM could efficiently gather topographical
information from specimens of realistic dimensions for medical practice,
where ceramic bone scaffold dimensions can range from roughly a few
millimeters to several centimeters depending on the relative size of the bone
in which the defect occurs. Of course, the overlay image (such as Figure 1a)
has the advantage of displaying far greater detail in the image than does the

topographical profile.

199

A contour map for specimen HA-70-A8-1 (corresponding to the overlay
image in Figure 1a) uses a gray scale to indicate depth and contour lines to
indicate "levels" of constant depth. (Although it is shown as a gray scale here,
the contour maps prepared by the CLSM'S computer are rendered in color.) In
Figure 1d, there is a height difference of 200 pm between adjacent contour
lines. The labels "A", "B" and "C" in Figure 1d denote depths of 1000 pm, 600
um, and 200 pm, respectively above the lowest point in the image
(designated as a height of zero). The contour map's information thus
complements that given by the overlay, phi-z and topographical images.

Figure 2, which was obtained using the Zeiss 210 CLSM in the confocal
reflection mode, shows specimen HA-70-AS-2 (as-sectioned specimen with a
70 volume percent porosity). Figure 2 is an overlay image consisting of a 100-
2 section series with total depth of 2 mm and a thickness of 2.0 pm for each
section. While the overlay image (Figure 2a) depicts the details of the
specimen surface, the phi-Z scan (Figure 2b) again helps us to understand

the three-dimensional nature of the surface. Figure 2b shows (for the cut A-B

across the overlay image), a steep-sided ridge 1325 um tall (Point P1 to P2 in

Figure 2b) and 363 um wide (point P2 to P3 in Figure 2b) that separates two

pore regions with relatively flat or gently sloping bottoms.

Specimen HA-70-ES-1, an epoxy impregnated, Rhodamine B stained
hydroxyapatite specimen with approximately 70 volume percent porosity was
imaged using the Zeiss Pascal 510 CLSM in the confocal fluorescence mode

(Figure 3). Figure 3a is a 100-section overlay image with a thickness of 1.0

200

 

 

pm for each section and total image depth of 100 pm. In Figure 3a, the two
lines that intersect at right angles (labeled at "1" and "2") mark the position of
two phi-z cuts on the image, the projections of which are shown in Figure 3b.
The two profiles depict the thickness and shape of a thin bridge of epoxy,
where the profile along cut 1 shows an arching section of epoxy that varies
from about 90 um down to about 15 pm in the vertical direction (Figure 3b). In
contrast, the bridging section along cut 2 is only about 80 um long with a
maximum thickness of about 20pm.

Figure 3c is an image in the fluorescence mode taken at a depth about 40
um below the specimen surface. Figure 3c depicts a "pore throat", an opening
between two pores, which is about 30 pm across at its narrowest point. The
use of the CLSM fluorescence mode to image subsurface features will likely
prove to be extremely important to the development of a process to
characterize the three-dimensional nature of bone scaffolds, in a manner
analogous to the analysis done for the interconnected porosity in geological
materials 8' 48' 49.

The pair of images given in Figures 4a and 4b provides a direct
comparison between a CLSM overlay image (Figure 4a) and a conventional
SEM (Figure 4b) micrograph of a sintered 70 volume percent HA specimen
(HA-70-A8-3). In order to facilitate comparison, the two micrographs (Figures
4a and 4b) depict the same location on the specimen surface and were taken

at the same magnification. While the SEM micrograph portrays well the

details of the uppermost portion of the specimen surface of the specimen,

201

 

 

there is no information from within the pores due to the limited depth of field of
the SEM. In contrast, the upper surface as well as the bottoms and side walls
of the pores are in focus in the CLSM overlay image. To further determine the
details of the surface topology, one could use the phi-Z scans, topographical
maps or contour maps, as discussed for Figure 1.

Figure 5a and 5b demonstrate the use of the CLSM reflection mode to
examine micron-scale features on the bone scaffold surfaces. The phi-Z scan
(top section of Figure 5a) corresponding to surface Shown (bottom section of
Figure 5a) shows surface roughness on the scale of several microns. For
example, the larger surface features in Figure 5a are on the scale of tens of
microns, where the horizontal distance from P1 to P2 is 27.6 um and the

vertical distance from point P1 to P2 is 31.8 um.

SUMMARY AND CONCLUSIONS

The development of ceramic bone scaffold materials depends on
developing methods to characterize the microstructure of the ceramic
scaffolds. These ceramic materials represent a cost of approximately $2.5
billion annually in the United States alone 7. When combined with the market
in craniofacial and other orthopaedic applications such as the total hip
replacement, bioceramics represent an increasingly popular choice for
biomedical applications. Confocal Laser Scanning Microscopy is a tool that is
well suited to characterizing important macro- and micro- scale features in

bone scaffolds and has the potential to receive wide spread use

202

 

 

characterizing other bioceramic materials such as alumina, zirconia and
Bioglass‘m.

While the CLSM reflection mode is well suited to characterize the 200 to
1000 um diameter surface-breaking pores in HA bone scaffolds, the
fluorescence mode is very powerful since it can image subsurface features.
This study has begun the process of adjusting the CLSM procedures
developed by geologists 8' 48' 49 for materials with a smaller average pore size
(about 25 to 40 pm on average). As work on these procedures progress, the
observation technique for subsurface features will be further refined, and then
complex statistical analysis can be applied to the CLSM images to fully

characterize the three-dimensional pore architecture of bone scaffolds.

ACKNOWLEDGEMENTS

The authors acknowledge the use of the Confocal Laser Scanning
Microscope and SEM facilities at the Advanced Microscopy Center, Michigan
State University along with the assistance of Dr. Shirley Owens and Dr.

Joanne Whallon.

203

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208

APPENDIX B

CONFOCAL LASER SCANNING MICROSCOPY (CLSM) AS A TOOL FOR
IMAGING CANCELLOUS BONE

I.O. SMITH, F. REN, M.J. BAUMANN AND E.D. CASE
Department of Chemical Engineering and Materials Science, Michigan State
University, East Lansing, MI, USA 48824
PUBLISHED IN: Journal of Biomedical Materials Research Part B.
2006;79:185-192.
ABSTRACT

Understanding the bimodal structure of cancellous bone is important
for tissue engineering in order to more accurately fabricate scaffolds to
promote bone ingrowth and vascularization in newly forming bone. In this
study, confocal laser scanning microscopy (CLSM) was used to create
detailed images of the bimodally porous intertrabecular space of defatted and
deproteinized cancellous canine bone taken from the epiphysis of the
humerus. The bimodal pore structure was imaged using both reflective and
fluorescent modes in CLSM, resulting in four different, but complementary
image types: 1) a Z-stack overlay, 2) a phi-Z scan, 3) a topographical map
and 4) a contour map. Submerging the bone in Rhodamine B dye prior to

fluorescent imaging enhanced the pore surface details giving a more accurate

pore size measurement. The average macropore diameter was found to be
260 um i 97 um while the average micropore diameter was 13 um i 10 um.
When compared to common techniques for imaging cancellous bone,

including micro computed tomography (pCT), magnetic resonance imaging

209

(MRI), scanning electron microscopy (SEM) and environmental scanning
electron microscopy (ESEM), CLSM was found to be an effective tool, given
its ability to non-destructively image the surface and near-surface pore

structure.

INTRODUCTION

This study is an extension of a series by the authors on bone and
engineered bone tissue materials, including the role of hydroxyapatite (HA)
surfaces in the regulation of osteoblast (OB) gene expression and
phenotype,1 electrostatic interactions of biomaterials with OBS and whole
bone,2 the electrostatic attraction between OBS and calcium deficient apatites
and OBS and biphasic calcium phosphate3 and the effect of HA on
differentiation and growth of M03T3-E1 0884. Additionally, the authors’
recent work includes studies of the dependence of the hardness5 and
dielectric constant6 on volume fraction porosity of HA, as well as studies of
microcracking,7'8 biaxial flexure strength9 and ceramic/ceramic joining by both
microwave and conventional sintering.”14 Recently, the authors were among
the first to apply Confocal Laser Scanning Microscopy (CLSM) to the study of
porosity in HA specimens,15 based in part on studies on the morphology of
surface channels in zirconia ceramics.1618

The hypothesis of this study IS that CLSM may be used to image
surface and near-surface features in cancellous bone, including macro- and

micropores. Three-dimensional representative images of cancellous canine

bone were created using stacks of individual confocal images (Z-stacks) and

210

 

 

 

 

both reflective and fluorescent modes in CLSM. These methods created
detailed images of the surface and near-surface structure of the macro- and

micropores, using a combination of Z-stack overlays, phi-Z scans, as well as
topographical and contour mapping. CLSM’S high resolution (2 pm) and high

depth of field (~2 mm) makes it suitable to image both the large scale
macropores and the finer interconnective micropores of cancellous bone.
Imaging is enhanced by fluorescent CLSM Z-stack overlay which also

increased the accuracy of pore Size measurements, Showing macropores

having an average diameter of 260 um i 97 um and micropores having an

average diameter of 13 um i 10 um. Three-dimensional images obtained in

this study give a better understanding of the bimodal pore structure of
cancellous bone, which is an essential part of more effectively designing bone

tissue engineering scaffolds.

BACKGROUND
Bone Structure

The structural elements of bone are divided into two types: cortical and
cancellous tissue.19 Cortical (compact) bone has a higher compressive
strength (69-110MPa)20 than cancellous (spongy) bone and is more abundant
in load-bearing bones, such as the femur.21

Cortical bone consists of nanocrystals of apatite, which has a

A,22'23

composition similar to H and collagen fibrils arranged in thin,

concentrically layered lamellar sheets. These are grouped around central

211

 

 

Haversian canals, which provide a pathway for blood vessels and nerve
fibers”21 Each grouping is referred to as an osteon or Haversian system.21
Within individual lamellae are interconnected spaces called lacunae, which
contain osteocytes that traverse each Haversian system, supplying nutrients
to the entire structure.”21

Cancellous (spongy) bone has a greater volume fraction porosity than
cortical bone19 with an open microstructure of trabeculae intermingled with

bone marrow.21

These trabeculae consist of lamellae, but do not form
osteons. Rather, nutrients are exchanged with the surrounding bone marrow
via lacunae contained within the trabecular lamellae.21 In long bones,
cancellous tissue is found primarily in the epiphysis and along the outer
surface of the medullary cavity.21

The bimodal porosity of cancellous bone tissue is an important factor in
its effective vascularization and for bone ingrowth. Microporosity is essential

during the initial stages of osseointegration, with increased microporosity

causing increased bone ingrowth.24

Imaging porous structures using CLSM

Geologists use CLSM to image subsurface pore structures in rock
where it is difficult to image such features without physically removing a
surface layer, thereby damaging the specimen. These subsurface details are
crucial, for example in water or oil flow in porous rock. CLSM has successfully

characterized the subsurface pore size, shape and interconnectivity,”28 as

212

 

 

well as microcracking,” in a wide range of geological materials by
impregnating porous specimens with a low viscosity epoxy containing a
fluorescent stain”:27

Visualizing the detailed shape of pores, including subsurface pore
structures, is also important when designing bone tissue engineered
scaffolds. Non-destructive sectioning of a solid specimen followed by
reconstruction of a detailed three-dimensional image of the pore structure is
beneficial in understanding and in turn mimicking the three-dimensional
nature of bone.

Recently, a modified version of the technique described above by
Fredrich et al.26 and Petford et al.27 has been used to image porous HA
scaffolds having a pore geometry similar to that of cancellous bone.15 While
this technique provided some detailed structural images, It must be further
refined before it can be used to effectively assess the complete subsurface
pore structure. Ren et al. were more successful in imaging the scaffold
structure by using the reflective mode of CLSM to image surface and near-
surface pores.“5 The larger-scale bimodal pore structure of cancellous bone is

similar to the pore structure of the HA scaffolds viewed by Ren et al.,

therefore similar CLSM techniques could be used.
Electron microscopy for imaging bone

Transmission electron microscopy (TEM) is commonly used to image

the micron and submicron features of bone, including cell structure,

213

 

”'30 as well

organelles, extracellular matrix and submicron structural elements,
as the bone-implant interface, studying the remodeling process between HA
and adjacent bone.”34 While TEM gives microstructural detail on the order of
a micron or smaller, it is not an effective tool to image the larger scale pore
structure in bone. Scanning electron microscopy (SEM) and CLSM are better
suited for imaging surface and subsurface pore structures because of their
lower magnification and greater depth of field.

Bone structures imaged using SEM include cancellous bone in human
proximal femora,35 cancellous bone in the human lumber spine,36 the fracture

surface of bovine subchondral bone37 and rat vertebrae.38 While the SEM

gives submicron resolution (3-6 nm best case), it is limited by its depth of field
(«1 pm at high magnification) along with its inability to image subsurface

features.”

MATERIALS AND METHODS
Sample preparation

A canine humerus, obtained from an animal sacrificed for reasons
unrelated to the current study, was stored at -20°C for approximately two
weeks prior to removal of muscle and skin tissue. Exposed bone was refrozen
and sectioned transversely across the epiphysis using a low speed diamond-
wafering blade at 200 rpm, revealing the trabeculae within the bone. A
second cut was then made, parallel to the first, leaving a slice of bone

approximately 10-mm thick. The rough sectioning used to obtain these

214

 

specimens did not adversely damage the specimen pore structure, as the
saw kerf is negligible when compared to the specimen size. Using a scalpel,
the periosteum and loose marrow were then carefully removed and the slice
sectioned into 8 to 10, 10-mm wide pieces, each including regions of cortical
and cancellous bone. Each piece was then defatted and deproteinized
following a procedure devised by Johnson et al.40 Dry samples of defatted
and deproteinized bone were stored in a petri dish until CLSM imaging.

Two specimens were reserved for fluorescent staining. One specimen
was globally stained by immersion in Rhodamine B solution for 12 hours. The
staining of a second specimen was localized by injection with Rhodamine B at

three pore openings along its surface, where only a small amount of dye (10

pl) was used to avoid spreading beyond the single pore.

CLSM Imaging

Two CLSM units were used in this study. One was better suited for
topographical studies in the reflection mode and the second better suited for
fluorescent mode. The CLSM'unit used in reflection mode* used an Ar 488
nm laser for maximum resolution. An optical filter was‘ not utilized; instead
computer program filters (low pass filter and median pass filter) reduced the

noise to signal ratio.

 

-' Zeiss 210 CLSM (Carl-Zeiss Inc. Jena, Germany)

215

The CLSM unit used in fluorescent mode“ was operated using a 488
nm Ar laser. Images were taken with a long pass filter (LP560) and a line

average of 4 to reduce noise. In both cases, 1 to 2-mm thick sections of bone
were optically sectioned into 50 or 100 slices (with thickness of ~10 um) using

CLSM.

SEM Imaging

Images of the bone from SEM were used for comparison purposes.
Specimens were mounted onto aluminum stubs, gold-coated (21 nm thick)
using a sputter coater and imaged at an accelerating voltage of 5 keV. SEM
image was also compared to CLSM image by measuring pore size using

phase-contrast imaging software.“

RESULTS
Reflection Mode

Reflective CLSM images of the defatted and deproteinized bone
surface show near-surface features (Fig. 1), including both macro- and
micropores. Images collected include Z-stack overlay (Fig. 1a), a single phi-Z
scan across the specimen surface (Fig. 1b), a topographical map comprised
of several phi-Z scans (Fig. 1c) and a contour map of the same surface,

showing a maximum depth of ~1.2 mm (Fig. 1d).

 

” Zeiss LSM 5 Pascal CLSM (Carl-Zeiss Inc. Jena, Germany)
AnalySlS v.3.2 (Soft Imaging System GmbH, Lakewood, CO)

216

 

 

 

 

Figure 1c Figure 1d

Figure 1: Four images taken at a single location on cancellous bone specimen, including (Fig.
1a) a Z-stack overlay image showing macropores and interconnecting micropores (indicated
with arrows), (Fig. 1b) a Phi-Z scan cross—sectional cut taken along the line shown, with A-B-
C shown in the upper portion indicating the features of a pore (A-B-C on the lower portion),
and C-B equal to 1.2 mm, (Fig. 1c) topographical map consisting of a wire mesh of
consecutive Phi-Z cuts and (Fig. 1d) a contour map showing pore morphology and depth.

The topographical map and the contour map are different than, but
complimentary to the Z-stack overlay image (Fig. 1). Contour mapping
provides a reference for the shape, local curvature and depth of each
macropore present. Topographical mapping builds three-dimensional images
of surface features, giving the viewer a visual representation of the
information contained in the contour map. When the topographical map and
contour map are viewed together with the Z-stack overlay image, this

combination is an intuitively useful representation of the specimen surface.

217

The surface area shown has several distinct macropores where each
macropore has micropores along its walls, indicated with arrows (Fig. 1a). In
Fig. 1b, the cross section of a macropore is indicated by A-B-C. The top-down

view is shown in the lower half of the figure and the side view in the upper
half. On the side view, the distance A-C is equal to 700 um and the distance

C-B is equal to 1.2 mm.

Topographical Fluorescence
One locally stained macropore, impregnated with Rhodamine B dye

solution was imaged using fluorescent CLSM (Fig. 2a) along with finer details,
including micropores, along its surface (Fig. 2b). This method was used to
image only the pores impregnated with dye. Other non-impregnated pores

could not be fluoresced.

218

 

 

:11 lyffifii

 

Figure 2a

 

Figure 2b

Figure 2: Local staining of cancellous bone imaging (a) the interior of a single macro- and (b)
micropores along the surface of a macropore.

219

 

 

A globally stained bone specimen was also imaged with fluorescent
CLSM, revealing the bimodal pore structure (Fig. 3a). Further images (figures

3b and 3c) collected using this method were used to measure the average

size of the macropores (260 um i 97 um) and the micropores (13 um i 10

pm). The globally stained specimen was more effectively imaged in

comparison to the locally stained specimen because the dye made it possible
to image any region of the structure. This increased available field of view

meant that the size of the pores could be more accurately measured.

220

 

 

Figure 3a Figure 3b

 

Figure 3c

Figure 3: CLSM reflectance images showing a single micropore on the surface of a
macropore for globally stained cancellous bone. Figures 3b and 3c were used in measuring
pore size.

SEM Comparison
The surface of defatted and deproteinized bone was imaged using both

reflective CLSM (Fig. 4a) and SEM (Fig. 4b). Analysis of the SEM image

yielded an average macropore size of 534.9 um i 265.9 um and micropore

size of 137.7 um i 50 um. Analysis of the CLSM image yielded an average

221

macropore Size of 506.0 pm 1- 252.4 um and micropore size of 106.6 um :t

47.7 urn. Differences between reported values are possibly a result of the

inability of SEM to image deep into the pores. The arrow shown on Figure 4a
indicates a deep pore not clearly visible in Figure 4b. Actual location of these
deep features can be ascertained using phi-Z scanning, topographical
mapping and contour mapping (Figure 1b-1d).

Values reported in both cases are near those reported by deGroot
(~200-400 um and micropores <60 pim).41 Compared to these established

ranges, apparent differences in this case are likely a result of the inability of
both techniques to effectively image smaller micropores, because of the
relatively low magnification used.

While the SEM image effectively shows micron-scale details of the
specimen surface, the sub-surface pore structure iS not visible. The CLSM
image (Fig. 4a) shows the pore interconnections while the corresponding
SEM image (Fig. 4b) does not. Further, the pore depth cannot be assessed
from SEM images. CLSM on the other hand, shows details from the surface
to greater depths by combining individual CLSM image slices into Z-stacks

greater than 1 mm in thickness.

222

 

 

 

 

Figure 4a

 

Figure 4b

Figure 4: A comparison between (a) CLSM image and (b) SEM image of the same section of
defatted and deproteinized cancellous bone. Fig. 4a was taken using reflective CLSM,
showing greater detail of pore depth, while Fig. 4b highlights the limitations of SEM imaging
because of lower depth of field.

223

DISCUSSION

Reflective CLSM techniques similar to those used by Ren et al.15

successfully imaged surface features of cancellous bone. Surface images
produced by stacking optical slices Into three-dimensional Z-Stacks have both
a sufficient resolution and a depth of field that is sufficient to show details of
the bone’s bimodal porosity, including macro- and micropores. In particular,
topographical and contour mapping (Fig. 1c and Fig. 1d) aid in visualization of
the specimen surface including details of pore shape, size, depth and
curvature of pore walls. Fluorescent CLSM of rhodamine dye impregnated
bone further improved the quality of Z-stack overlay images of the near-
surface bone structure. The improved contrast between fluorescing bone
mineral and non-fluorescing empty pore space enhanced the image quality
thus more accurate measurement of macro- and micropore size was possible.

CLSM images taken from globally stained bone are more versatile than
locally stained bone. The images shown in Figure 2 are two of three locally
stained pores on the Specimen surface. When selecting a pore for imaging,
these three are the only choices, ruling out the remainder of the surface
pores. This makes it difficult to observe a true representation of the specimen
surface, as each pore is unique. In contrast, globally stained specimens
provide a larger number of pores for imaging. The images shown in Figure 3
are a sampling of dozens of surface-breaking pores on the specimen. This
wider variety of available pores allows the viewer to see a better

representation of the specimen surface. In essence, local staining requires

224

 

the viewer to choose pores of interest before observing them, while global
staining allows the viewer to observe the entire surface when looking for
regions of interest.

Common bone imaging techniques include micro computed tomography

(pCT), magnetic resonance imaging (MRI) and SEM. uCT gives a three-

dimensional image of porous structures with resolutions up to 5 pm using

42-44

benchtop units and approaching 1 pm using synchrotron radiation

sources.45 However, high-flux synchrotron radiation damages living tissue.
Also, synchrotron sources are rare and beam time is not easily obtainable.

CLSM on the other hand images specimens of living tissue, and has a
resolution of ~2 um.15
MRI resolves and images the larger features of cancellous bone (~100-

200 pm), but is unsuitable for smaller features such as micropores (<60

um).46 Furthermore, current MRI techniques do not include the compositing of

two-dimensional images into three dimensions and MRI units are a far more
expensive option, with unit costs over $1 million, compared to ~$200,000 for a
CLSM unit.

While SEM imaged the surface features of bone, those images did not
include detailed features at depths possible with CLSM. Also, SEM does not
image qualities of the pore structure, on the order of those Shown by contour
mapping (Fig. 1d). The lack of these abilities prevent accurate pore depth

measurement in bone using SEM. Additionally, SEM specimens used in this

225

study were free of water, because of the high vacuum required. This
requirement rules out imaging whole bone tissue or bone cells in future
studies?9 The additional requirement that surfaces of SEM specimens also be
conductive requires that ceramic-containing materials be imaged using a low
potential or by adding a conductive coating, further limiting the use of living
tissue.39 Environmental scanning electron microscopy (ESEM) images water-
containing specimens, but still requires a partial vacuum of 1 — 3 kPa.39 Both
SEM and ESEM vacuum Chambers also constrain the specimen size (<100
mm in most cases), while CLSM uses an open stage, which accommodates
specimens up to 20-30 mm thick and >100 mm in length and width.15

CLSM is a high-resolution imaging tool, imaging details as small as 2
pm with high depth of field capability, and is capable of creating three-

dimensional images of known dimensions. In this study, CLSM imaging
resulted in detailed images of both macro- and micropores comprising the
bimodally porous intertrabecular space of cancellous bone. In the design of
bone replacement scaffolds, mimicking the bimodal pore structure promotes
bone ingrowth and proper vascularization in newly forming bone. Imaging
pore structures via CLSM and using the resulting Z-stack overlays,
topographical maps and contour maps, which are a source of valuable
structural data, helps the viewer to better understand this structure. Although
we have successfully imaged pore structures in HA15 and in cancellous bone

(this study) via CLSM, we have not yet imaged the pore structure of cortical

226

 

bone using CLSM. However, the ordered structure of cortical bone is

comparatively well understood”21

CONCLUSION

CLSM has a lower lateral resolution than SEM and a lateral resolution
comparable to uCT. However, the lateral resolution of CLSM is sufficient for

detailed imaging of both macro- and micropores in cancellous bone.
Fluorescent CLSM further improves contrast of these images, increasing the
accuracy of pore size measurement. Additionally, CLSM does not require
special specimen preparation or conditions damaging to living specimens for
operation, which makes it possible to image living cells and tissue, a prospect

for future studies.

The authors acknowledge the invaluable assistance and resources
provided by' the Center for Advanced Microscopy at Michigan State
University, especially that of Dr. Shirley Owens and Professor Emeritus
Joanne Whallon. Additionally, the authors thank Stuart Kaltz, an

undergraduate student, for his assistance in preparing the bone specimens.

227

 

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232

 

 

APPENDIX C

SAMPLE CALCULATION FOR PERCENT OF OSTEOBLASTS ATTACHED
TO DENSE HYDROXYAPATITE SPECIMEN (CONTROL VS.
MICROCRACKED, 4 HR ATTACHMENT).

MC3T3-E1 osteoblasts (OBS) were counted in 10 select fields of view, as
described in Chapter 2. These counts were then used to calculate the number
of cells attached on each specimen. A Microsoft Excel spreadsheet was used
to determine mean attachment values and standard deviation for the control
and microcracked groups.

Table 1: Cell counts taken from 10 fields of view per specimen (9 control specimens, 9
microcracked

 

233

Table 2: Average percent attached osteoblasts for both the control and microcracked group.

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Mean Per cm2 Per sample % of attached
5.1 2454.80 10162.87 21.61
4.5 2166.00 8967.24 19.07
3.4 1636.53 6775.25 14.41
3 1444.00 5978.16 12.71
2.1 1010.80 4184.71 8.90
4.3 2069.73 8568.70 18.22
5.1 2454.80 10162.87 21.61 Raw Count Per cm2 % attached
7.1 3417.47 14148.31 30.08 MEAN 4.27 2053.69 18.08
3.8 1829.07 7572.34 16.10 STDEV 1.45 696.27 6.13
4.5 2166.00 8967.24 19.07
6.4 3080.53 12753.41 27.12
4.7 2262.27 9365.78 19.91
9.8 4717.07 19528.65 41.52
5.6 2695.47 11159.23 23.73
8.3 3995.07 16539.57 35.17
6.9 3321.20 13749.77 29.23 Raw Count Per cm: % attached
6.9 3321.20 13749.77 29.23 MEAN 6.86 3299.81 29.05
8.6 4139.47 17137.39 36.44 STDEV 1.80 864.09 7.61

 

 

 

 

 

 

 

 

 

 

To determine statistical significance in the difference between these
mean attachment values, we performed an unpaired Students t-test, taking n
< 0.05 to indicate statistical significance. The results of this test are shown
below. Group A is the control group and Group B is the microcracked group.

Student's t-Test: Results

The results of an unpaired t-test performed at 22:40 on 27-JAN-2007

= ~3.37

sdev= 1.63

degrees of freedom

hypothesis, is 0.0039

Group A: Number of items = 9
2.10 3.00 3.40 3.80 4.30 4.50 5.10 5.10 7.10

Mean = 4.27

95% confidence interval for Mean: 3.115 thru 5.419

Standard Deviation = 1.45
Hi = 7.10 Low = 2.10

Median =
Average Absolute Deviation from Median = 1.06

4.30

Group B: Number of items = 9
4.50 4.70 5.60 6.40 6.90 6.90 8.30 8.60 9.80

16 The probability of this result, assuming the null

 

 

 

Mean = 6.86
95% confidence interval for Mean: 5.704 thru 8.008

234

 

Standard Deviation = 1.80

Hi = 9.80 Low = 4.50

Median = 6.90

Average Absolute Deviation from Median = 1.38

 

 

235

 

APPENDIX D

X—RAY DIFFRACTION OF DENSE HYDROXYAPATITE USED IN THIS
STUDY

500
450 I
400 j
350
300 -

— HA Disc

200 .
150 .
100

50 '

 

Figure 1: X-ray diffraction data of dense sintered hydroxyapatite (Taihei Chemical).

236

 

APPENDIX E

ELECTROSTATIC INTERACTIONS AS A PREDICTOR FOR OSTEOBLAST

ATTACHMENT TO BIOMATERIALS

I.O. SMITH*, M.J. BAUMANN*, L.R. MCCABE‘t
*Department of Chemical Engineering and Materials Science, *Department of

Physiology, Michigan State University, East Lansing, MI 48824-1226

PUBLISHED IN: J Biomed Mater Res 2004;70Az436-441.

SYNOPSIS
The present study utilizes zeta (8;) potential analysis as an indicator of

bonding of osteoblasts (OBS) and whole bone to various biomaterials.

Common metal alloys (316L stainless steel, CoCrMo and Ti6Al4V) and
bioceramics (HA and B-TCP) used in orthopaedic applications were

suspended in particulate form in physiologic saline (P8), both as-received and
supplemented with bovine serum albumin (BSA). Metal alloys were also

treated with NaOH washing to study the effect of such a surface treatment on

the C-potential. The NaOH wash was found to increase the C-potential for

CoCrMo and Ti6Al4V and a decrease in the magnitude of the C-potential for

316L stainless steel. When the metal alloy powders were suspended in BSA-

supplemented PS, the Q-potential as a function of pH increased, thereby

237

increasing the electronegativity gap and increasing the propensity for bonding
between each of the metal alloys and bone. This increase is likely due to
matrix proteins in the BSA, which adsorb onto the metal alloy surfaces,

promoting bone growth. With the addition of BSA to each bioceramic system,
a uniform decrease in Q-potential was observed. However, the
electronegativity gap remains large in each case, maintaining the anticipation
of bonding. C-potential analysis is an effective predictor of biomaterial

attraction to osteoblasts and bone, providing a useful in vitro method for

predicting such interactions.

INTRODUCTION

The biocompatibility of biomaterials in contact with bone has been
widely examined.1 Such materials, including metals and ceramics, have a
wide array of uses in orthopedic implants. However, the exact mechanisms of
the implant-bone bond are not clearly understood. While surface properties,
such as roughness and porosity, and chemical treatments, such as NaOH
washes affect the extent of the implant to bone attachment, a comprehensive
explanation for the nature of this bond is currently lacking. Biocompatible
alloys such as CoCrMo, 316L stainless steel and Ti6Al4V have been
commonly used in the fabrication of orthopedic implants. Calcium phosphate
ceramic materials, such as hydroxyapatite (HA) and tricalcium phosphate
(TCP), have been a focus in the development of surface reactive bone

scaffolds and implant coatings. These bioceramics have been shown to form

238

an interfacial bond with bone in previous studies and it is therefore
reasonable to assume they will Show a propensity for bonding with osteoblast
(OB) cells, which are the dominant cell type in normal physiologic bone
growth.2 Based on these findings, it is also logical by extension, to
hypothesize that these ceramics will show a propensity for bonding to whole
bone.

Several studies have linked the Q-potential of bone to the piezoelectric
effect:3 Opperrnann et al. examined the stability of the interactions between
bone and various compositions of Bioglass© and hydroxyapatite (HA) using :-
potential data as a function of pH.4 Their results showed that highly stable
interactions were indicative of successful in vivo bonding between optimum
Bioglass© compositions to bone and between HA and bone. An effective
method to quantitatively examine bone-implant bonding is to measure the
mobility of colloidal particles in a suspension, using ultrasonic attenuation
spectroscopy (UAS), from which the surface, or zeta (Q) potential is
calculated. This study will employ Q-potential analysis using UAS, as a
function of physiologically relevant pH, to predict bone-implant interactions in
a variety of bioceramic and biocompatible metal powder materials. The metal
powders also were washed with NaOH in order to measure any change in the
implant to bone surface potentials and relate this to noted improvements in in
vivo bone to implant bonding.5 This approach was verified for the case of OB-
HA interactions by examining in vitro OB cell, growth and proliferation as a

function of time on HA discs.

239

 

 

MATERIALS AND METHODS
Suspension Preparation

M_eta§ Three metal powders, representative of metals used in total
knee and total hip prostheses were examined (316L SS (Ametek, Eighty
Four, PA), 750F CoCrMo and Ti6Al4V (Powder Alloy Corp., Cincinnati, OH)).
To normalize particle size in these powders, each was ball milled dry for 14

days at <2°C with AI203 grinding media prior to analysis, leaving particle sizes
of 10 um, 30 um and 80 pm, respectively. To simulate in vivo exposure to

proteins upon implantation and ionic activity, the ball-milled powders were
next suspended at 1 volume % in physiologic saline (PS) at 0.154M NaCl with
1% (by volume) bovine serum albumin. Samples of each metal powder were
then washed in 0.1M NaOH by suspending each powder and stirring for 15
minutes, then allowing particles to settle for 1.5 hours before removing liquid
and allowing powder to dry under mild heat. Excess NaOH solution was
decanted and the powders air-dried, before subsequent resuspension in PS
for UAS measurement under a N2 blanket. In this study, the O’Brien model,
which takes into effect double-layer distortion, was used to calculate the (-
potential from the UAS mobility data which was measured using the
AcoustoSizer ll ESA apparatus (Colloidal Dynamics, Warwick, RI)."'7
Amplitudes were measured across a range of frequencies, yielding an
attenuation spectrum. For each set of lg-potential data, 30 measurements of
particle size were taken and the Gaussian average of these utilized in t;-

potential calculation at each pH value. To eliminate adsorption of ions from

240

 

 

the atmosphere, all Q-potential measurements for each of the materials and
cells examined in this study were carried out under a N2 blanket.

m Bone stock was taken from the proximal femur of a deer harvested
within 4 hours of death and sectioned for immediate grinding. The soft tissue
and periosteum were removed and the bone wet ground from the sub-
periosteal surface to the mid-cortex, while submerged in P8, using a diamond
conical grinding bit at 3,000 rpm, under moderate hand pressure. The
resulting bone particle suspension was then subjected to immediate Q-
potential measurement. For each material, the mobility was measured and
the resulting Q-potential calculated and plotted as a function of pH. The range
of pH was controlled, under constant stirring, with automated titration using a
complementary acid (1 M HCI) and base (1M NaOH).

Osteoblasts: MC3T3-E1 mouse osteoblast cells were cultured in a minimum
essential medium (a—MEM, Invitrogen Corp., Carlsbad, CA) supplemented
with 10% fetal bovine serum (FBS, Atlanta Biologicals, Norcross, GA) at 37°C
in a humid atmosphere containing 5% CO2 and 95% air. The media was
aspirated off and the cells washed once with 1X phosphate buffered saline.
Cells were then trypsinized and subsequently pelleted by centrifuge at 10,000
rpm for 10 min. The resulting OB pellets were resuspended by aspiration in
P8 in 50 ml tubes. P8 was prepared at three different pH levels (7.3, 7.4, 7.5)
by titrating with complementary acid and base (1M HCI and 1M NaOH). The
resulting suspensions were then transferred to 60 ml syringes and finally,

injected in turn into the UAS cell chamber where instantaneous mobility

241

 

measurements were taken and C-potential values calculated and recorded as

a function of pH. This technique is known as static C-potential measurement

(as the suspension does not constantly flow through the Chamber) and was
chosen because of the small volume of OB suspension available for
measurement.

_Bioceramics: Three calcium phosphate ceramic powders, typical of those
used in bone tissue engineering scaffolds were examined, unsintered HA

(Berkeley Advanced Biomaterials Inc., San Leandro, CA), sintered HA and

sintered B-TCP (HiMed, Old Bethpage, NY) having particle sizes of 0.85 pm,

0.95 pm and 1.15 um, respectively. HiMed HA powder was sieved by the
manufacturer using a proprietary process to achieve a narrower particle size
distribution. These powders were suspended in PS and subjected to ;-
potential measurement with the AcoustoSizer ll under a N2 blanket. For each
material, the Q-potential was collected and plotted as a function of

physiologically relevant pH. The pH range was controlled with automated
titration with complementary acid (1M HCI) and base (1M NaOH) under
constant stirring. For comparison, HA suspensions were also prepared in P8

with 1% (by volume) bovine serum albumin.

242

 

 

RESULTS
Metal Alloys

The q-potentials as a function of pH for the as-received metal alloy
powders are shown in Figure 1. Across three separate data sets, C-potential
date was found to be very nearly identical for each metal. For 316L 88, the Q-
potential ranges from —1.4 at pH 6.3 to —3.2 mV at pH 7.5, while for CoCrMo
and Ti6Al4V, the Q-potential ranges from —1.3 at pH 6.3 to —1.7 mV at pH 7.8
and from —2.8 mV at pH 6.5 to —3.1 mV at pH 7.8, respectively. In each case,
the metal powders possess C-potentials that are significantly more positive
with respect to the bone which has a Q-potential of —75 mV through the entire

pH range examined and is shown in Figure 2. Kowalchuk et al., using bone

particles having a diameter of >5 um, reported c-potential values from 0 to -
15 mV, while the bone particles used in this study have a diameter of 1pm
and a Q-potential of —75 mV across comparable pH ranges.8 Because the

magnitude of the q-potential is indirectly proportional to the diameter of the
particulate material, the C-potential values obtained in this study are therefore
consistent with this earlier reportf"13 as a five-fold increase in particle size
decreases the surface area of the powder by a factor of approximately five
while holding volume constant. With NaOH washing, the C-potentials as a
function of pH for each metal powder remain negative over the pH range and
are shown in Figure 3 where the Q-potential for 316L 88 shifts to more

negative values from -6.0 mV to —10.2 mV, more positive values for CoCrMo

243

 

 

from —0.4 mV to —1.0 mV and for Ti6Al4V shifting to more negative values of

an approximately constant —2.5 mV.

 

 

OI r T 77 fi

_05 2 --+"316L

‘1 A --I---CoCrMo
9 I
E, O. ...... --*--Ti-6Al-4V
a -1.5+ """ I---
:3 ..~ ------- I --------- l
5 ~.
:5 -2: ‘
‘f
U .

-2.5- L ‘ ~-

‘ ......... ‘---::--..
':---I.~
-32 "~c "A
.35-J
pH

Figure 1: g-potential as a function of pH for aS-received 316L, CoCrMo and Ti6Al4V vs. whole
bone suspended at 1% in PS.

§— potential (mV)
:3
o

 

I - - Q - -Whole Bone

 

Figure 2: q-potential as a function of pH for ground whole deer bone suspended at 1% in PS.

244

The C-potentials as a function of pH in BSA for both the as-received
and the NaOH metals are shown in Figures 4 and 5, respectively. For the as-
received metals in BSA, over the indicated pH, the C-potential of 316L SS

ranges from —2.5 to -4.8 mV, for CoCrMo from 3.7 to 2.5 mV and for Ti-6AL-
4V from —1.5 to -—1.7 mV. When comparing to the as-received metal alloy

powders not suspended in BSA, the metal alloy powders dispersed in BSA all

exhibited positive shifts in the C-potential. Dispersing the NaOH washed metal

alloy powders in BSA over the same pH range gave positive shifts in the @-

potentials for 316L 88 from —2.2 mV to —3.6 mV, for CoCrMo from —0.3 to —

1.0 mV and for Ti6Al4V a relatively constant —1.3 mV. Again, in all cases, the

Q-potential values of the metal alloy powders dispersed in BSA are

significantly more positive than those of bone at a C-potential of -75 mV.

 

 

-2 .
I ---------- A ------- A ------------ A
5‘ .4 -
E - . 9- - ~316L + NaOH rinse
g -6 4) ...... - - I- - -CoCrMo+ NaOH rinse
Q ......
é ----- .. . . - - fi- - -Ti-6AI-4V + NaOH rinse
klI‘ '8 4 ...... e
-10 -
7 9
-12 -
pH

Figure 3: Q-potential as a function of pH for NaOH-washed 316L, CoCrMo and Ti6Al4V vs.
whole bone suspended at 1% in PS.

245

 

 

4‘ I--
3- ..... ' ----
. ”“1- ................ I
2- ""
S‘ .
£1
E 0.0— T Y 1
§_1g5 65 7 75 8
g k ------- A --------------- A --------- A
3'2“
e...
-34 “O ------------ . .....
._.~
_4, -+-316L
I “‘o
-5« -I--CoCrMo
‘5 ‘ - i- -Ti-6AI-4V
pH

Figure 4: Q-potential as a function of pH for as-received 316L, CoCrMo, Ti6Al4V and whole
bone suspended at 1% in P8 with 1% (by volume) bovine serum albumin.

 

o . . , .,
..... 6.5 7 7.5 8
-0.5 4 """"" l ‘
‘- .......
-1 -l ....... l
S" A ------------------- A """"""" A ------------ A
g '1.5 “
E - O -316L + NaOH rinse
‘5: -2 «.,
'5 - I -CoCrMo + NaOH rinse
‘1 -2 5 -
d» ' - i- -TI-6AI-4V + NaOH rinse
‘ ‘9
-3 —. “
\“ ...... .- 'o.- I
-3.5 - ' ' . 0

 

I
h
L

pH

Figure 5: Q-potential as a function of pH for NaOH-washed 316L, CoCrMo, Ti6Al4V and
whole bone suspended at 1% in P8 with 1% (by volume) bovine serum albumin.

Bioceramics

For OBS, whole bone, HA and B-TCP in suspension without BSA, the

C-potential as a function of pH is shown in Figure 6. The C-potential of the

246

sintered HA ranges from 4.1 mV at pH 7 to 4.2 mV at pH 7.8. Unsintered HA

has a Q-potential from 17.1 mV at pH 6.7 to 14.5 mV at pH 7.7 and B-TCP
has a relatively constant C-potential from 2.5 mV at pH 6.6 to 2.1 mV at pH
8.1. With BSA additions, shown in Figure 7, the C-potential of sintered HA
Shows little change from the C-potential values of the suspension without
BSA, having a Q-potential from 3.8 mV at pH 6.1 to 3.5 mV at pH 7.6. The
unsintered HA has a Q-potential ranging from 11.8 mV at pH 6 to 10.4 mV at
pH 7.7 while the B-TCP has C-potentials ranging from 1.1 mV at pH 6.1 to —

0.5 mV at pH 8. From pH 7.3 to 7.5, 083 had negative Q-potentials of —29.4

 

 

mV to —52.4 mV.
40 1
20 -
I - - - -- - - - .-
0 f ‘ - I ’r-l I. -L$ =.-.- -T-A 1
9 .
g b 6.5 7 7.5 8 8.5 pH
E
“E 201
9
g - o - HA (Sintered)
J» _40 _ - -I - HA (Unsintered)
- i - Beta TCP
—¢—Osteoblasts '
-60 a +Bone
+ a I
-80

 

Figure 6: Q-potential as a function of pH for HA (sintered), HA (unsintered), g-TCP,
osteoblasts and whole bone suspended at 1% in PS.

247

 

202

 

 

 

9
E -20
:13 - Q - HA (Sintered)
é - -I - HA(Unsintered)
3 ‘40 ‘ - I - BetaTCP
—O—Osteoblasts
-50 - +Bone
-80 I a I
pH

Figure 7: q-potential as a function of pH for HA (sintered), HA (unsintered) and B-TCP in 1%
(by volume) bovine serum albumin in comparison to osteoblasts and whole bone suspended

at 1% in PS.

DISCUSSION

Metal Alloys

Each of the metal alloy powders, over the given pH range, has an average @-

potential value near zero, where NaOH treatment caused an increase in the

C-potential for CoCrMo and Ti6Al4V and a decrease in the magnitude of the

Q-potential for 316L stainless steel. While it is widely held that none of these

common orthopaedic materials exhibits direct bonding to bone, NaOH
treatment has been shown to increase apatite layer formation in vitro on the
surface of Ti alloys.”15 Here, Ti forms an alkali titanate hydrogel layer to
which bone is observed to directly bond, both in vitro and in vivo.14 While the
zero point of charge for Ti occurs at pH 6.2, over the physiologic pH ranges

used in the current study (pH 6-8), it is therefore expected that the Ti6Al4V

248

 

 

 

would be negatively charged as shown in Figure 1.16 With exposure to NaOH,
Kim et al. postulate that the titanium oxide on the Ti6Al4V surface is partially
dissolved by the OH” ions, to form HTiOg’ at the same time as the underiying
base Ti metal is hydrating to form Ti(OH)3+.14 Therefore, with NaOH
treatment, there is an increase in the number of positive ions formed on the
Ti6Al4V surface that corresponds to the noted increase in Q-potential found in
this study and shown in Figure 3. In the case of 316L, it is possible that the
steel’s self-healing oxidation layer may not remain impermeable to continued
dissolution in an OH' -Containing environment. This may contribute to the
larger variation in Q-potential values experienced by the 316L SS in
comparison to either the CoCr or the Ti6Al4V powders.17 Significant
electronegativity differences |~75mV| between whole bone and each of the
metal alloy powder materials were observed. This large gap in
electronegativity suggests a propensity for electrostatic bonding between

each metal alloy and whole bone.

BSA was added to each colloidal system in order to create an in vitro
environment that more closely matches in vivo conditions. Our results
demonstrate that BSA increases the electronegativity gap, and therefore
attraction, between each of the metal alloys and bone. In addition, BSA
contains extracellular matrix proteins, such as fibronectin, that adsorb onto
metal surfaces, thereby providing sites for integrin mediated osteoblast

attachment. Specific integrin and extracellular matrix combinations can

249

influence osteoblast adhesion and contribute to surface attraction and

subsequent promotion of bone growth.
Bioceramics

Unsintered HA exhibits uniformly greater Q-potential values in comparison to
Sintered HA, ~17 mV as compared to ~4 mV, over the given pH range. While
x-ray diffraction results (not shown here) for both the sintered and unsintered
HA powders demonstrate that the only phase present is pure hydroxyapatite,
the Ca/P ratio determined by energy dispersive spectroscopy are similar but
not identical (2.30:0.2 for unsintered HA and 1671008 for sintered HA).
Therefore, differences in powder processing have led to changes in the Ca/P
ratio, which have altered the surface chemistry, leading to differences in the
surface or Q-potential. The differences in the Q-potential in tum necessitate
changes in the experimental protocol if these powders are processed in an
aqueous slurry. For example, the higher magnitude of the Q-potential in the
unsintered HA makes it is easier to disperse and form a uniform slurry. A
uniform slurry is necessary to form a uniform green body, and hence a
uniform scaffold.18 Unsintered HA also has a greater electronegativity
difference with respect to both osteoblast and whole bone in comparison to
sintered HA and B-TCP. Therefore, unsintered HA would be expected to show
a greater propensity for bonding with both osteoblasts and whole bone.
However, as an unsintered powder must still be sintered to form a scaffold, it
is worthwhile to note that significant electronegativity differences are still

present for both sintered HA and B-TCP. Another factor that accounts for the

250

 

differences in Q-potential between these bioceramics are the differences in
particle size. Because the magnitude of the q-potential is indirectly
proportional to the diameter of the particulate material, the differences in Q-
potential can also be partially attributed to the fact that unsintered HA has a
mean particle size that is less than that of sintered HA, 0.85 pm versus 0.95
pm, respectively. The low Q-potential values for the B-TCP (Ca3(PO4)2) may

be understood by noting the lack of hydroxyl groups, which are charge

determining ions on HA (Ca1o(PO4)5(OH)2).

Contrary to the metals examined in this study, BSA additions to the
bioceramic suspensions acted to uniformly lower the Q-potential

measurements for each material. However, the electronegativity difference
between each of these bioceramics with respect to either osteoblasts or
whole bone, is still on the order of 80 mV, meaning that attraction and
bonding between these materials and bone is still anticipated. These findings
reinforce the argument that calcium phosphate ceramics are suitable for bone
scaffold applications, both for osteoblast seeding as well as attachment to
existing bone. Future studies will probe the experimental link between surface

potential and OB attachment and mineralization.

ACKNOWLEDGEMENTS
Funding for this study was granted by the National Science Foundation

(DMR0074439) and the MSU Foundation. Assistance in osteoblast cell

251

 

 

culture was kindly provided by M. Hossain and R. McMullen and for x-ray

diffraction measurements to M. Soto.

252

REFERENCES

1. Park JB, Lakes RS. Biomaterials: An Introduction. 2nd ed. Plenum Press,
NY 1992

2. Vander, Sherman, Luciano. Human Physiology: The Mechanisms of Body
Function. 8‘“ ed. McGraw Hill, NY 2001

3. Fukada E, Yasuda I, On the piezoelectric effect of bone. J. Phys. Soc., 12:
1158-1162, (1957).

4. Oppermann DA, Crimp MJ, Bement DM, In vitro stability predictions for the
bone/bioglass and bone/HA interactions, J. Biomed. Mater. Res., 42:412-416
(1998)

5. Gross K. Fluor—hydroxyapatite solid solutions as alternative bioceramics.
Proceedings from lntemational Society of Ceramics in Medicine, Bioceramics-
14, BB (2001)

6. O’Brien RW. EIectro-acoustic effects in a dilute suspension of spherical
particles. J. Fluid Mech., 190, 71-86 (1986)

7. O’Brien RW. The Electro-acoustic equation for a colloidal suspension. J.
Fluid Mech., 212, 81-93 (1990)

8. Kowalchuk RM, Pollack 8R, Corcoran TA. Zeta potential of bone from
particle electrophoresis: Solution composition and kinetic effects. J. Biomed.
Mater. Res., 29, 47-57 (1995)

9. Lange FF. Powder processing science and technology for increased
reliability. J. Am. Ceram. Soc., 72, 3-15 (1989)

10. Velamekanni BV, Lange FF. Effect of interparticle potentials and
sedimentation on particle packing density of bimodal particle distributions
during pressure filtration. J. Am. Ceram. Soc. 74, 166-172 (1991)

11. Chang JC, Velamekanni BV, Lange FF, Pearson DS, Centrifugal
Consolidation of Al203 and AI203/ZI'02 Composite Slurries vs. Interparticle
Potentials: Particle Packing and Mass Segregation. J. Am. Ceram. Soc., 74,
2201-2204 (1991)

12. Crimp MJ, Oppermann DA. Green Density Variations in Relation to

Colloidal Packing in SiC/Si3N4 Composites. Proc. Am. Ceram. 800.: Science,
Technology and Applications of Colloidal Suspensions, 91-101 (1995)

253

 

13. Crimp MJ. A Novel Approach to lntermetallic Matrix Composite (IMC)
Processing. J. Mater. Sci. Eng., A192/193, 633-639 (1995)

14. Kim H-M, Miyaji F, Kokubo T, Nakamura T. Preparation of bioactive Ti
and its alloys via simple chemical surface treatment. J. Biomed. Mater. Res.,
32, 409-417 (1996)

15. Wang CX, Wang M. Electrochemical impedance study of the nucleation
and growth of apatite on titanium pretreated with NaOH solution. Society for
Biomaterials 27th Annual Meeting, 101 (2001)

16. Boehm HP. Acidic and basic properties of hydroxylated metal oxide
surfaces. Disc. Farad. Soc., 52, 264-275 (1971)

17. Smith, WF. Principles of Materials Science and Engineering. McGraw-
Hill, NY 1986

18. Sacks MD. Rheological science in ceramic processing. Science of

Ceramic Chemical Processing, Ed. L.L. Hench and DR. Ulrich, J. Wiley 8.
Sons, 522-538 (1986)

254

 

APPENDIX F
SURFACE POTENTIAL AND OSTEOBLAST ATTRACTION TO CALCIUM
PHOSPHATE COMPOUNDS IS AFFECTED BY SELECTED ALKALINE
HYDROLYSIS PROCESSING
l.O. SMITH*, M.J. BAUMANN*, L. OBADIA‘, J-M BOULERI,
*Department of Chemical Engineering and Materials Science, Michigan State
University, East Lansing, MI 48824-1226, USA, *lNSERM EM99-03, Faculté
de Chirurgie Dentaire, BP 84215, 44042 Nantes, France

PUBLISHED IN: J Mater ScizMater Med 2004;15:841-846.

ABSTRACT

This study examines the link(s) between the suspension behavior of
calcium deficient apatites (CDAs) and biphasic calcium phosphate (BCP), as
measured by the g-potential, with respect to both whole bone and
osteoblasts. CDA is fabricated by hydrolyzing an acidic CaP such as
dicalcium diphosphate dihydrate (DCPD; CaHPO4, 2H2O) and has a structure
and composition close to bone apatite. Sintering CDA results in the formation
of BCP ceramics consisting of mixtures of HA and B-TCP, with the HA/B-TCP
weight ratio proportional to the Ca/P ratio of CDA. The choice of the base for
the DCPD hydrolysis allows various ionic partial substitution of the formed
CDA. Na for Ca partial substitution is of interest because of the resulting
improvement in mechanical properties of the resulting BCP ceramics and
NH4OH was used as a negative control. The Q-potential was measured for
these materials and the stability of the ceramic to bone interaction calculated.

Q-potential values decrease for CDA(NH4OH) versus CDA(NaOH) and

255

 

increase for BCP(NH4OH) versus BCP(NaOH). While results of these
analyses indicate that NH4OH and NaOH processed CDA and BCP will likely
yield osteoblast attachment in vivo, differences in the C-potentials may explain

varying degrees of cell attachment.

INTRODUCTION

The surface potential of orthopedic biomaterials in contact with bone has
been the subject of recent interest, however, the exact mechanisms of the
implant to bone bond are not clearly understood. [1-3] While surface
functional groups such as carboxyl and hydroxyl ions determine the surface
Charge and hence the surface potential of many common ceramics, the
degree to which such surface groups alter the resulting ceramic to osteoblast
or bone attachment is less understood. [4-5] Bioceramics such as
hydroxyapatite (HA; Caio(PO4)6(OH)2) and beta-tricalcium phosphate (B-TCP;
Ca3(PO4)2), have been a focus in the development of osteoconductive
scaffolds for bone tissue engineering and improved orthopaedic implant
coatings. [6-8] They respectively present hexagonal and rhombohedral crystal
symmetries. [9]

The link between the surface potential, as assessed by the zeta (Q)
potential, of these bioceramics and their resulting attraction to both bone and
osteoblasts, have been shown in previous studies. [1-3] The focus of this
paper is the relationship between modifications in the processing method of

calcium phosphate powders and their resulting Q-potential, and hence

256

suitability for use as bone tissue engineering scaffolds. Several studies have
linked the surface, or q-potential of bone to the piezoelectric effect.
Oppermann et al. examined the stability of the interactions between bone and
various compositions of BioglassTM and hydroxyapatite (HA) using Q-potential
data as a function of pH. [2] Their results showed that unstable interactions
were indicative of successful in vivo bonding between optimum BioglassTM
compositions and bone and between HA and bone. An effective method to
quantitatively examine bone-implant bonding is to measure the mobility of
colloidal particles in a suspension, using ultrasonic attenuation spectroscopy
(UAS), from which the surface, or q-potential is calculated. This study will
employ Q-potential analysis using UAS, as a function of physiologically
relevant pH, to predict bone-implant interactions in a variety of bioceramic
powder materials that have been synthesized via two different processing
methods.

One possible way of forming calcium deficient apatite (CDA) with both
a structure and composition close to bone apatite is by hydrolyzing an acidic
calcium phosphate phase such as dicalcium diphosphate dihydrate (DCPD;
CaHPO4, 2H2O). [10-11] Sintering CDA between 900°C and 1200°C results in
the formation of biphasic calcium phosphate (BCP) ceramics consisting of
mixtures of HA and B-TCP with the HA/B-TCP weight ratio proportional to the
Ca/P ratio of CDA. [12] The choice of the base for the DCPD hydrolysis
allows various ionic partial substitution of the formed CDA. Na for Ca partial

substitution is of interest because of the resulting improvement in mechanical

257

properties of the final BCP ceramics. [13] Although biological fluids contain
large quantities of ionic sodium, the biological role of ionic sodium within
calcium phosphate (CaP) materials, with respect to cell adhesion and activity,
remains unclear. In order to measure the influence of Na-substitution on the
surface potentials of CaP materials, two kinds of chemical bases NH4OH and
NaOH were chosen, knowing that NH4OH does not induce any particular ionic
substitution.

The degree of ionic activity between bone and synthetic materials may
be assessed through the Q-potential as calculated from the acoustophoretic
particle mobility. O’Brien and Oja have extensively described the
electrokinetic phenomena of acoustophoresis. [14-16] The Q-potential is
calculated using the acoustophoretic mobility with a correction for the particle
inertia G(a)‘1, in an alternating field. This correction reduces the velocity
amplitude of particle motion and was derived using the Helmholtz-

Smoluchowski equation. [14-15]
4' = (707/ 8080001)" (1)

Where so is the perrnitivity in a vacuum, a, is the relative perrnitivity, and G(ot)’1
is given by:

 

5(a)"'=[1 ia(3+2Ap/p) ] (2)

—9{1+(1—i)(a/2)“2}

To describe the effects of homo- and heterocoagulation and to
generate a quantitative theory for the overall kinetic stability of a system of

non-identical particles, Hogg et al. developed an expression for interparticle

258

stability (er). This factor, termed the stability ratio is in effect, the ratio of

particle collision to collisions that result in agglomeration. [19]

°° V dr
WU = I expo/j):2 (3)

where the interparticle intercomponentaSIgbility (W,-,-) varies exponentially with
the total force, VT, (attractive and repulsive) and inversely with the square of
the particle separation distance, r, from a separation equal to the sum of the
two particle radii, a,- and a; to an infinite separation distance into the solution.
The repulsive force, which acts as a barrier to coagulation is proportional to
the surface Charge of the particle. However, this surface charge is not readily
measurable because of continuous ion activity at the particle/fluid interface.
Therefore, the surface charge is estimated at the double layer by the @-
potential, which in this study, is measured using ultrasonic attenuation
spectroscopy (UAS) methods. Stable interactions, i.e. suspensions for which

there is no flocculation exist at W > 20, while unstable interactions are

predicted for systems where W< 20.

MATERIALS AND METHODS
Powder Processing Methods
Calcium-deficient apatite (CDA; Calnx (HPO4)x (PO4)6.x (OH)2-x) powders
were obtained by hydrolysing two, 409 batches of dicalcium phosphate
dihydrate (DCPD) powder (Merck, Darrnstadt, Germany) in aqueous alkaline
solutions (500ml with [NaOH] = 0.35 M and [NH4OH] = 0.3M). These two

chemical reactions were performed at 90°C under stable magnetic stirring for

259

4 hours. The CDA powders were washed with distilled water and dried (80°C,
12h). They were then calcinated at 1050 °C for 4 hours to obtain a biphasic
calcium phosphate (BCP), a solid mixture of HA and B-TCP.

Powder Analysis

The purity of the starting and final materials was analyzed by assessing
specific surface area (SSA) and Na content: SSA of both CDA and BCP
powders were measured by BET adsorption and the sodium content was
checked using atomic absorption spectroscopy.

Suspension Preparation

In this study, the O’Brien model, which takes into effect double-layer
distortion, was used to calculate the Q-potential from the UAS mobility data,
which was measured using the AcoustoSizer Il ESA apparatus (Colloidal
Dynamics, Wanlvick, RI). To eliminate adsorption of ions from the atmosphere,
all Q-potential measurements for each of the materials and cells examined in
this study were carried out under a N2 blanket. Measurements were not made
in the presence of an adsorbed protein such as bovine serum albumin as
prior work has shown it to have only minimal effect on the surface potential.
[20] All suspensions were prepared in physiologic saline (PS; 0.154M NaCl)
at a concentration of ~1 volume%.

_B__<_)_n_e_: Bone stock was taken from the proximal femur of a deer harvested
within 4 hours of death and sectioned for immediate grinding. The soft tissue
and periosteum were removed and the bone wet ground from the sub-

periosteal surface to the mid-cortex, while submerged in PS, using a diamond

260

conical grinding bit at 3,000rpm, under moderate hand pressure. The
resulting bone particle suspension was then subjected to immediate t;-
potential and particle size measurement.

Osteoblasts: MC3T3-E1 mouse osteoblast cells were cultured in a minimum
essential medium (a-MEM, Invitrogen Corp., Carlsbad, CA) supplemented
with 10% fetal bovine serum (FBS, Atlanta Biologicals, Norcross, GA) at 37°C
in a humid atmosphere containing 5% CO2 and 95% air. The media was
aspirated off and the cells washed once with 1X phosphate buffered saline.
Cells were then trypsinized and subsequently pelleted by centrifuging at
10,000rpm for 10min. The resulting OB pellets were resuspended by
aspiration in P8 in 50ml tubes. P8 was prepared at three different pH levels
(7.3, 7.4, 7.5) by titrating with complementary acid and base (1M HCI and 1M
NaOH). The resulting suspensions were then transferred to 60ml syringes
and finally, injected in turn into the UAS cell Chamber where instantaneous
mobility measurements were taken and Q-potential values calculated and
recorded as a function of pH. This technique is known as static Q-potential
measurement (as the suspension does not constantly flow through the
Chamber) and was chosen because of the small volume of OB suspension
available for measurement.

Bioceramics: Two types of CDA (hydrolized using NaOH and NH4OH ) along
with two types of BCP bioceramic were assessed for differences in their @-
potential. The average diameters of the CDA (NaOH), CDA (NH4OH), BCP

(NaOH) and BCP (NH4OH) were measured, under identical suspension

261

conditions, as 0.95, 0.85, 1.0 and 0.93pm, respectively. Because the
magnitude of the q-potential is indirectly proportional to the diameter of the
particulate material, the bone was ground until a particle size of 1pm was
reached. These particle sizes were also measured as a function of pH to
detect agglomeration and each of the powders remained essentially
unagglomerated throughout the pH range examined. Each powder was
suspended in PS and subjected to Q-potential measurement with the
AcoustoSizer ll under a N2 blanket. The Q-potential was then collected and
plotted as a function of physiologically relevant pH. The pH range was
controlled with automated titration with complementary acid (1M HCI) and

base (1 M NaOH) under constant stirring.

RESULTS AND DISCUSSION

Powder analysis

FTIR spectra (not shown) verified that the classic P04 and OH vibration peaks
were present and that the presence of carbonate and hydrogenophosphate
bands was minimally present for both CDA powders. [21]

Rietvelt analysis is applied to XRD data for polycrystalline materials in
order to better refine peaks, yielding a more accurate determination of the
crystal structure. In this study, Rietvelt refinement was used to determine the
lattice parameters of the HA and B-TCP in BCP(NaOH) (HA: a = b =
0.9416(2); C = 0.6887(2) l B-TCP: a = b = 1.0408(2); c = 3.7228(9)) and in

BCP(NH4OH) (HA: a = b = 0.9420(2); c = 0.6881(2) / B-TCP: a = b =

262

1.0403(2); c = 3.7368(9)) (all values in nanometers). The HA/B-TCP weight
ratio was found to be 3.48/6.52i0.301 for BCP(NH4OH) and 3.53/6.57:0.352
for BCP(NaOH). This yielded the atomic Ca/P ratios for both
CDA/BCP(NH4OH): 1.56 and CDA/BCP(NaOH): 1.57. Only the c values for [B-
TCP, present in both BCP powders, were significantly different. This indicates
that Na for Ca substitution occurs preferentially in the B-TCP phase.
According to the crystal structure of B-TCP, the only possible Ca site is the
incomplete one noted Ca(4).

Specific surface area (SSA) was determined using BET adsorption.
The SSA is 120.3:0.15 m2/g for CDA(NH4OH), 6073:005 m2/g for
CDA(NaOH) and 2.85i0.01m2/g for both BCP powders. Atomic absorption
spectroscopy measurements give Na content for both CDA(NaOH) and
BCP(NaOH) as 1.97ir0.09%. For both CDA(NH4OH) and BCP(NH4OH), the
amount of Na was found to be 0.45i0.09%. This latter result Is due to Na
impurities present in the dicalcium phosphate dihydrate (CaHPO4-H20, also
known as DCPD), which is used as a starting material in the alkaline
hydrolysis.
C-potential
q-potential data, as a function of pH, for the calcium phosphate powders is
Shown in Figure 1. The iso-electric point (iep) is the pH at which the net
surface potential is zero. The iep for the CDA powders varies from a pH of 6.8
for the CDA (NH4OH) Shifting to a more basic pH of 7.4 for the CDA (NaOH),

while the BCP processed with NH4OH does not exhibit an iep over the pH

263

range tested and the BCP processed using NaOH has an iep at pH 6.4. The
BCP (NH4OH) powder behaves very much like commercially available HA,
having relatively little positive potential and no iep. [20] The Q-potential values
range from 5.1 to —0.7mV for the CDA (NaOH), 1.5 to —4.2mV for the CDA
(NH4OH), 1.5 to —2.8mV for the BCP (NaOH) and 2.9 to 1.5mV for the BCP
(NH4OH) over the approximate pH range from 6 to 8. Sodium substitution in
the CDA powders caused a slight increase in the overall Q-potential curve
while in the BCP powders, it led to a slight decrease in the overall Q-potential
values. This decrease in the BCP Q-potential with sodium ion substitution is
logical considering that Na ions having a +1 charge are substituting for Ca
ions having a +2 charge, thus decreasing the overall net Charge within the
material. The smaller positive shift in the CDA materials may indicate either
that an insufficient amount of Na was substituted into the structure and/or that
NH4OH caused other changes in the surface chemistry of CDA that resulted
in the noted increase in the overall surface potential.

Kowalchuk et al., using bone particles having a diameter of >5um,
reported Q-potential values from 0 to -15mV, while the bone particles used in
this study have a diameter of 1pm and a relatively constant q-potential of —
75mV across comparable pH ranges. [23] Because the magnitude of the Q-
potential is indirectly proportional to the diameter of the particulate material,
the Q-potential values obtained in this study are therefore consistent with this

earlier finding. [3],[17],[24-26]

264

 

Stability

Stability calculations plotted as a function of pH for CDA(NaOH),
CDA(NH4OH), BCP(NaOH) and BCP (NH4OH) are shown in Figures 2-5. At
all pHs tested, the CDA(NaOH) remained unstable (W18 20) and is predicted
to flocculate in suspension in physiologic saline while the bone shows values
of stability, W22 > 20, and is hence stable over all test pHs. Interestingly, even
though there is a difference in both the sign and hence the magnitude of the
zeta potential between the bone and the CDA(NaOH), the interaction
between the bone and bioceramic (W12) is predicted to be stable (W12 > 20)
and hence each component will remain separately stable and unflocculated
over the pH range. Similar results were also found for the CDA(NH4OH),

BCP(NaOH) and BCP (NH4OH).

265

6 , +CDA32(NaOH)
+CDA30(NH4OH)
BCP32(NaOH)
«: BCP30(NH4OH)

 

 

9 10

 

Zeta Potential (mV)
N

 

-10 .
pH

 

'20 r r Y T 1
.

-30 a

mVj

~3I£——— Osteoblasts

-50 . Bone

 

 

-60 n

Zeta Potential (

-70 .

 

 

-8O ,

pH
Figure 1: Q-potential for CDA(NaOH), CDA(NH4OH), BCP(NaOH), and BCP(NH4OH) (2A). @-
potential for MC3T3-E1 mouse osteoblasts and ground bone (ZB).Both figure 2A and 2B are

shown as a function of physiologically relevant pH at solids loadings of ~1 volume % In 0.154
M NaCI electrolyte.

266

3.50902 -
3.00902 -
2.50E+02 -
2.00E+02 1
1.50902 4
1.00902 .
5.005+01 I

LOG(W)

 

-5.00E+01 0

W‘ H H H H n 7‘ H

8

+LOG10W11
+LOG10W12
LOG10W22
0.00E+00 «o—o—o—o—o—o—o—o—o—o—o—o—o—o—o—Hwn
6.5 7 7.5
pH

Figure 2: Stability calculations for CDA(NaOH), denoted W11, bone (W22) and the interaction
between CDA(NaOH) and bone (W12).

3.50E+02 1
3.00E+02 n
2.50E+02 ~
2.00E+02 a
1.50E+02 ~
1.00E+02 -«
5.00E+01 4

LOG(W)

 

5.00901 5

7‘ 7‘ H H H 7‘ H 7‘ H 7‘ 7‘ 7‘ H N H

H

8

+LOG1OW11
+LOGIOW12
LOG10W22
0.00E+00 «Wmc—o—o—o—e—e—e—o—o—v-o—e—o—o—o—o—efi
6.5 7 7.5
pH

Figure 3: Stability calculations for CDA(NH4OH), denoted W11, bone (W22) and the interaction
between CDA(NH4OH) and bone (W12).

267

 

LOG(W)

3.50E+02 1

 

 

3_OOE+02 «n n n n n n n n n n n r! n n n n n n n n

2.50E+02 ~ \

200802 _ +LOG10W11

150902 ‘ —I—LOG10W12

1.00E+02 a LOG10W22

5.00E+01 n

0.00E+00 ~ —

-5.005+01 6 6.5 7 7.5 8
pH

Figure 4: Stability calculations for BCP(NaOH), denoted W11, bone (W22) and the interaction
between BCP(NaOH) and bone (W12).

LOG(W)

3.50E+02 —
3.00E+02 -
2.50E+02 -
2.00E+02 —
1.50E+02 ~
1.00E+02 n
5.00E+01 r

.A—H—i/V‘ 7‘ 7‘ 7‘ 7‘ 7‘ 7‘ 7‘ 7‘. 7‘ 7‘ 7‘ 7‘ 7‘ 7‘

+ LOG1OW11
+ LOG1OW12
LOGIOW22

0.00E+00 “WW

 

-5.00E+01

6.5 7.5

pH

Figure 5: Stability calculations for BCP(NH4OH), denoted W11, bone (W22) and the interaction
between BCP(NH4OH) and bone (W12).

While it is well known that osteoblasts attach to calcium phosphate

surfaces, and indeed bone has been shown to integrate into these implant

materials, the mechanisms for such bonding are not completely understood.

[27] The Q-potential results in this paper agree with those reported by

268

Ducheyne et al. who found small C-potentials of 0-10mV using <5um
hydroxyapatite powders. [28] The powders in this study were submicron in
size, giving comparable magnitudes and Sign for the Q-potential. Using a
comparably sized HA powder (diameter < 50m) purchased from a commercial
supplier, Oppermann et al. report much greater q-potentials on the order of
~130mV. Therefore, while both of these powders are nominally
hydroxyapatite, they have very different surface potentials, which are likely to
result from differences in powder processing and surface area.

So, while the hydroxyapatite-based powders in this study have similar
particle sizes and Q-potential magnitudes, the effect of different processing
treatments, such as sodium substitution on both the surface potential and
SSA, must still be further evaluated. As each of the powders examined in the
current study are calcium phosphate-based, it is logical to expect that their
surface potentials are similar, which indeed, the data from Figure 1 illustrate.
The primary difference between the two groups of calcium phosphate
powders is the nature of the inorganic base (NH4OH and NaOH) in the
hydrolysis reaction during powder processing. Sodium ions are added to
substitute partially for Ca in both the HA and B-TCP phases, where the Ca/P
ratio decreases with Na substitution but the (Ca+Na/2)/P ratio remains
constant, and these Na ions will then alter the surface chemistry. Such cation
substitutions, if concentrated on the powder surface, should yield a
disproportionate change the net overall surface potential, and hence the

resulting g-potential of the powder. However, the specific surface area (SSA)

269

was found to be twice as high for CDA (NH4OH), at ~120 m2/g, in comparison
to CDA (NaOH), at ~60 m2/g. This doubling of SSA with NH4OH may also be
a factor in understanding why sodium substitution did not produce the
expected decrease in the Q-potential for the CDA powders. In addition, if it is
assumed that only the atoms at the surface of the CaP are contributing to the
surface potential, then a 2% addition of Na to the particle only involves
~1.3x10“’% increase in Na at the surface of 10m particles. Therefore, such a
slight change may not be detected by Q-potential analysis. Further work will
include measuring the Q-potential of CDA and BCP powders with increased
Na substitution to verify the nature of the changes in the surface potential.

Differences in the Q-potential in turn necessitate changes in the
experimental protocol if these powders are processed in aqueous Slurries. For
example, the greater the magnitude of the Q-potential, the easier it is to
disperse and form a uniform slurry. A uniform slurry is necessary to form a
uniform green body, and hence a uniform scaffold. [29] None of the ceramic
materials tested in this study have a sufficiently large enough Q-potential to
exist as a dispersed, stable suspension, however all of the calcium phosphate
powders examined have zeta ”potential curves that have a sufficiently large
enough electronegativity gap with respect to both bone and osteoblasts to
expect that each calcium phosphate would favor the initial attachment of
osteoblasts.

Such differences may be the key to interpreting conflicting osteoblast

attachment results. Given the significant difference between the magnitude

270

and Sign between the Q-potentials of each of the calcium phosphate powders
and both the osteoblasts and bone at pH 7.4 (250mV and 270mV,
respectively), it may be that while the stability curves rightly represent the
thermodynamic steady-state, osteoblast attachment to any of the calcium
phosphate powders examined in this paper is more accurately characterized
by differences in the sign of the net Q-potentials at a given pH because
osteoblast attachment occurs over short times, involving multiple organic and

cellular components.

ACKNOWLEDGEMENTS
Funding for this study was granted by the National Science Foundation
(DMR0074439), the MSU Foundation and CPER Pays de la Loire — Axe

Biomatériaux 83.

271

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